Bio-Templating Growth of Nepenthes-Like N-Doped Graphene as

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Bio-Templating Growth of Nepenthes-Like N-Doped Graphene as Bifunctional Polysulfide Scavenger for Li-S Batteries Qiucheng Li, Yingze Song, Runzhang Xu, Li Zhang, Jing Gao, Zhou Xia, Zhengnan Tian, Nan Wei, Mark H. Rümmeli, Xiaolong Zou, Jingyu Sun, and Zhongfan Liu ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b05246 • Publication Date (Web): 11 Sep 2018 Downloaded from http://pubs.acs.org on September 12, 2018

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Bio-Templating Growth of Nepenthes-Like N-Doped Graphene as Bifunctional Polysulfide Scavenger for Li−S Batteries Qiucheng Li1,2†, Yingze Song1,2†, Runzhang Xu3†, Li Zhang1,2*, Jing Gao1,2, Zhou Xia1,2, Zhengnan Tian1,2, Nan Wei1,2, Mark H. Rümmeli1,2, Xiaolong Zou3, Jingyu Sun1,2*, and Zhongfan Liu1,2,4*

1

College of Energy, Soochow Institute for Energy and Materials InnovationS (SIEMIS), Soochow

University, Suzhou 215006, P. R. China 2

Key Laboratory of Advanced Carbon Materials and Wearable Energy Technologies of Jiangsu

Province, Soochow University, Suzhou 215006, P. R. China 3

Shenzhen Geim Graphene Center (SGC), Tsinghua-Berkeley Shenzhen Institute (TBSI), Tsinghua

University, Shenzhen, Guangdong 518055, P. R. China 4

Center for Nanochemistry (CNC), Beijing Science and Engineering Center for Nanocarbons,

Beijing National Laboratory for Molecular Sciences, College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, P. R. China

KEYWORDS: nitrogen-doped graphene, chemical vapor deposition, bio-templating growth, polysulfide scavenger, Li−S batteries.

ABSTRACT: The practical application of lithium–sulfur (Li–S) batteries is hindered by their poor cycling stabilities that primarily stem from the “shuttle” of dissolved lithium polysulfides. Here, we develop a nepenthes-like N-doped hierarchical graphene (NHG)-based separator to realize an 1

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efficient polysulfide scavenger for Li–S batteries. The 3D textural porous NHG architectures are realized by our designed bio-templating CVD approach via the employment of naturally abundant diatomite as the growth substrate. Benefiting from the high surface area, devious inner-channel structure and abundant nitrogen doping of CVD-grown NHG frameworks, thus-derived separator favorably synergizes bi-functionality of physical confinement and chemical immobilization toward polysulfides, accompanied by smooth lithium ion diffusions. Accordingly, the batteries with the NHG-based separator delivers an initial capacity of 868 mAh g–1 with an average capacity decay of only 0.067% per cycle at 2 C for 800 cycles. A capacity of 805 mAh g–1 can further be achieved at a high sulfur loading of ~7.2 mg cm-2. The present study demonstrates the potential in constructing high-energy and long-life Li–S batteries upon separator modification.

Lithium−sulfur (Li−S) battery has gained considerable attention as a promising alternative to lithium-ion batteries due to its exceptional theoretical capacity (1672 mAh g−1), high specific energy density (2600 Wh kg−1), natural abundance of sulfur and environmental benignancy.1-4 Despite these conspicuous merits, the practical application of Li−S batteries has still been far from satisfactory, which originates from the poor conductivity of sulfur and hence limited utilization, the free migration of dissolved lithium polysulfides (Li2Sx, 4 ≤ x ≤ 8) between electrodes, as well as the noticeable volume expansion of the sulfur cathode (ca. 79.2%) during lithiation.5,6 In particular, the dissolution of polysulfide intermediates would cause severe shuttle effect, consequently resulting in a rapid capacity fading, depressed rate capability, and marked anode corrosion. The key to high-performance Li−S battery lies primarily in the effective confinement of soluble 2


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polysulfides within the cathode side. As such, great efforts have been devoted to alleviate the shuttle effect, mainly pertaining to the strategic modification of sulfur host materials and advanced design of separators/interlayers to capture Li2Sx via structural blockage and/or chemical entrapment. One straightforward solution deals with the uniform dispersion of sulfur active materials into host matrix containing diverse micro/meso/macro-porous carbon,1,7,8 conductive polymers,9,10 and metal oxides/sulfides/nitrides.11-13 The resulting host composite cathodes have demonstrated suppressed polysulfide shuttling, augmented sulfur utilization, and enhanced rate capability. However, such material incorporations would inevitably limit the sulfur loadings because of the porous nature of the hosts, thereby handicapping the advantage of high volumetric energy density of Li−S battery. Building a thin and functional interlayer onto the separator is another promising strategy to obstruct the migration of polysulﬁdes and guarantee their recycling in the cathode side. In this regard, an ideal interlayer should meet the following conditions: (1) the presence of devious physical architectures and strong chemical anchors for polysulfide entrapment, (2) excellent conductivity for Li2S precipitation and recycling, (3) compact structures with scarcely existed cavities that inhibit free leakage of soluble polysulfides, (4) superior wettability for electrolyte permeation. Along this line, modified separators based on black phosphorus,14 MoS2,15 TiO2-TiN,16 metal-organic framework,17 and covalent organic framework18 ad-layers have been adopted as polysulfide reservoirs due to their strongly coupled interfaces or precisely tuned pore sizes, leading to improved cyclability of Li−S batteries. However, the poor conductivity of these materials would certainly jeopardize the recycle of captured polysulﬁdes, with their tedious/costly fabrication process further hindering uses in practical high-rate Li−S batteries. As the cost-effective and electrically-conductive materials with enriched morphologies, carbon 3


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nanostructures possessing one-dimensional (1D, such as carbon nano-fibers/tubes),19-21 2D (graphene sheets and films)22-27 and 3D (carbon aerogel, super P, and carbonized eggshell membrane)28-30 have readily demonstrated their great potential in the rational construction of separators/interlayers. Specifically, 3D carbon nanostructures with gloriously complex inter-channel structure would endow them with enhanced trapping possibility of Li2Sx species. More importantly, the permselectivity for lithium ions vs. polysulfides of 3D architectures is higher as compared to that of randomly oriented 1D networks (i.e., CNT to allow polysulfide migration) and scarcely defective yet crumpled 2D planar sheets (i.e., graphene to block ion transport), which results in selectively sieving lithium ions whilst impeding polysulfides. To further retard the polysulfide migration, heteroatom doping (e.g., nitrogen, boron) has been developed as an effective strategy for creating chemisorption sites to better anchor Li2Sx.31 In this respect, single-step, one-batch production of heteroatom-doped 3D graphene nanostructures

is

indeed

desirable.

Chemical

vapor

deposition

(CVD)

employing

heteroatom-containing carbon precursor offers the opportunity for directly synthesizing doped graphene with tunable defect density and versatile 3D morphology.32,33 This can be simply realized by conformal growth of 3D graphene powders with naturally abundant biosilica serving as sacrificial templates, which, nevertheless, has not been reported thus far in the field of Li−S batteries. Herein, we design a nepenthes-like NHG membrane as a promising polysulfide scavenger that synergizes bi-functionality of physical confinement and chemisorption of Li2Sx for Li−S batteries. The NHG materials are generated via direct CVD of N-doped graphene on natural diatomite template. The delicate nepenthes-like architectures of NHG are perfectly inherited from the bio-templates experiencing CVD reactions, with such conformal graphene coatings simply retained after the template removal. This bio-templating CVD approach enables batch production and precise control 4


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of the doping concentrations/configurations of high-quality NHG, which is distinguished from the conventional, exfoliation-based graphene materials (e.g., reduced graphene oxide, rGO). Benefiting from the high surface area, versatile pore structure and abundant nitrogen doping of NHG materials, the thus-derived separators manifest an efficient performance for polysulfide trapping. Moreover, the excellent conductivity of CVD-grown NHG framework is beneficial to expediting the catalytic conversion of long-chain Li2Sx into insoluble Li2S2/Li2S, offering an additional pathway to obstruct polysulfide shuttling. Our work puts forward a bio-inspired strategy for the design of fascinating barrier architecture for highly efficient polysulfide trapping towards practical Li−S battery with favorable rate capability and long lifespan.

Results and Discussion It is indeed the wisdom of nature that many organisms rely upon hierarchical bio-assemblies with periodic micro- and nano-scale features formed throughout complex cellular processes. For instance, diatomite is a bio-enabled natural silica material possessing versatile hierarchical morphologies. In this work, a specific type of diatomite, coscinodiscus wailesii, normally constituted by two valves (namely epitheca and hypotheca) combining together as a “petri dish”,34 was selected as the template for the growth of NHG due to its high natural abundance, low cost and multiple pore structures. As with such biosilica templates, a variety of hierarchical structures could be observed in detail, as shown in Figure S1, Supporting Information. The NHG materials were synthesized in a 4-inch low-pressure CVD (LPCVD) system with pyridine employed as both the nitrogen and carbon source (Figure S2). During the CVD reaction, the decomposed CHx and N species from pyridine precursors nucleated on the surface and inside the 5


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pores of diatomite, consequently leading to the conformal coating of 3D N-doped graphene on such templates (Figure 1a). After growth, the originally wheat-colored diatomite powders turned grey in macroscopic quantity, as shown in Figure S3. The synthetic details are further revealed in Methods and Figure S4. After the removal of the diatomite templates by acid etching and subsequent cryo-drying process, 3D porous NHG products can be readily obtained. The as-fabricated NHG architecture resembles a round-shape microplate (10-20 µm in diameter and 50-100 nm in thickness), which is featured by two representative hierarchical pores locating at the edge (bordered pores) and at the center (central pores), respectively (Figure 1b, c). The existence of ample pores within the graphene frameworks can be witnessed by scanning electron microscopy (SEM) and atomic force microscopy (AFM), as depicted in Figure 1d and Figure S5 and 6. Due to the 3D textural porous structure and light-weighted feature of NHG powders, the specific surface area value markedly outperforms that of original diatomite templates (3.8 m2 g-1), reaching 424.4 m2 g-1 based on the Brunauer-Emmett-Teller result (Figure 1e). Interestingly, the weak van der Waals interaction between the non-planar NHG architectures guarantee their uniform dispersion in suitable solvents (e.g., N-methyl-2-pyrrolidone, ethanol), as evidenced by the conspicuous Tyndall effect (Figure S7). As such, homogeneous NHG film can be attained by a simple vacuum-filtration method, demonstrating outstanding flexibility and uniformity at a macroscopic scale (Figure 1f). Moreover, the thus-derived NHG layer manifests favorable electrical conductivity with an average sheet resistance of ~19.3 Ω sq-1 (Figure 1g), according to the sheet resistance mapping results acquired by a four-point probe technique. One striking feature of the NHG architectures lies in the multiple and versatile pore structures. To systematically examine the multifariously hierarchical pore characters, high-resolution 6


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transmission electron microscopy (HRTEM) were carried out. Figure 1h reveals a tilt view of the biomorphic graphene, clearly showing the presence of the macro-pores (diameters of ~200 nm) and tortuous channels. Further close-up TEM views in Figure 1i and j reveal the length of interconnected pore channels ranges ~100-600 nm. It is noted that the channel lengths in TEM images are usually shorter than the real scenarios, since they could be folded on the copper grid upon sample preparation. Notably, meso-pores (~4-20 nm) are also scattered at the entrance of the channels (Figure 1k), owing to the perfect replication of the inner detailed structures of the diatom frustule. Exhaustive TEM observations disclosing the diversity of macro/meso-pore structures were carried out (Figure S8). Benefiting from such features with defective pores serving as ion tunnels, hierarchical structures acting as multi-leveled polysulfide obstructions, and nitrogen dopants functioning as polysulfide anchors, the NHG frameworks are anticipated to effectively capture polysulfides without compromising facile Li+ ion transport. By directly incorporating nitrogen atoms into the graphene lattice during growth, pyridine precursor-based CVD approach enables the production of graphene materials with high doping density and large-scale dopant uniformity. The doping level of the as-fabricated NHG can be comprehensively dictated by the growth temperature, possibly throughout the temperature-dependent competition between the formation of C-C and C-N bonds.35 Lower temperature appears to be favourable for the higher doping level, whilst elevated temperature is beneficial to a better crystalline quality, as confirmed by the X-ray photoelectron spectroscopic (XPS) and Raman spectroscopic measurements of the NHG grown at 1000 oC, 900 oC, and 800 oC, respectively (Figure S9 and 10). As a result, a growth temperature of 900 oC was selected to realize a delicate balance between the doping concentration and crystalline quality, which was further verified by structural and elemental 7


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characterizations. Figure 2a exhibits an optical microscopy (OM) image of an individual NHG architecture transferred on a SiO2/Si surface, showing a flower-like shape with detailed porous structures. The contour of the NHG can be directly reflected by Raman mapping (G band) result shown in Figure 2b, which displays a homogenous contrast with respect to the uniform graphene quality. Corresponding Raman spectrum of the NHG film (red line in Figure 2c) exhibits characteristic peaks of graphene located at 1339 cm-1 (D band), 1581 cm-1 (G band), and 2687 cm-1 (2D band), respectively. In comparison with the intrinsic hierarchical graphene (HG) generated by the identical temperature (900 o

C; 200 sccm Ar, 50 sccm H2, 10 sccm CH4 for 3 h), the 2D peak intensity of NHG is apparently

declining, which can be attributed to the topological defects originated from the N-atom insertion in the sp2 carbon lattices.36 The information on the I2D/IG intensity ratio and the full width at half maximum of 2D band implies that the NHG framework is of few-layer nature, in good agreement with the HRTEM inspection of a five layered NHG (Figure 2f). XPS analysis was carried out on the NHG samples (with a nitrogen content of 5.3 atom%) to examine the nitrogen doping information in graphene. The asymmetric C 1s peak (Figure 2d) can be fitted into three peaks at 284.8 eV, 285.9 eV, and 287.1 eV, corresponding to the graphitic sp2 carbon, two different C−N bonding structures with sp2 and sp3 hybridized carbon, respectively.37 As for the high-resolution N 1s signal in Figure 2e, the peak locating at 398.3 eV, 400.7 eV, and 402.6 eV can be ascribed to pyridine, graphitic, and pyrrolic N, respectively.38 Amongst these three N-doping configurations, the electron-rich donors of pyridine and pyrrolic N account for 35.8%, which are expected to form strong binding with polysulfides through dipole-dipole interactions.26 The spatial distribution of nitrogen dopants was further inspected with the aid of scanning 8


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transmission electron microcopy (STEM) imaging and energy-dispersive X-ray spectroscopic (EDS) mapping. The STEM image of a representative area (Figure 2g) and corresponding EDS maps of carbon (Figure 2h) and nitrogen (Figure 2i) manifest that nitrogen dopants are homogenously distributed in the porous NHG framework, indicative of a successfully uniform spatial distribution of nitrogen doping upon CVD growth. Along this line, few-layer NHG materials with escalated doping content and uniform doping distribution can be readily obtained via a direct bio-templating CVD approach using pyridine precursor. To draw a comparison of the polysulfide trapping ability between NHG and other carbon materials, permeation experiments were systematically carried out with the aid of homemade H-shaped quartz devices to visualize the polysulfide diffusion (Figure 3a-c). The NHG powders were vacuum-filtrated onto polypropylene Celgard separators with a mass density of 0.1 mg cm−2 (with preparation details displayed in Methods). The top-view SEM and cross-sectional SEM/EDS characterizations confirms the one-side coating of NHG frameworks on Celgard separator with a thickness of ~2.2 µm (Figure S11). Accordingly, the modified separators were jointly sealed in the middle of the H-shaped device, where the left part of the tube was filled with polysulfide solution (7 mmol L−1 Li2S6 in a dioxolane/dimethoxyethane mixture) and the right side injected with blank electrolyte. The permeation experiments with the HG and rGO separators were performed under identical conditions. As for the NHG-based separator, it is evident that the migration of the polysulfide solution was refrained within a period of 48 h (Figure 3a), illustrating the effective blockage towards polysulfide diffusion. In comparison, the permeation experiment with HG separator demonstrated that polysulfides started to penetrate through the separator within 24 h (Figure 3b), indicating that the polysulfide trapping effect of HG would be solely stemming from 9


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physical blocking. By contrast, the rGO-based separator even malfunctioned at the initial stage of test (i.e., < 1 h) when the polysulfides began to diffuse out of the separator. In turn, the fresh electrolyte at the right sided tube turned completely yellow within 48 h (Figure 3c). This result directly corroborates that the NHG framework is beneficial to inhibiting the polysulfide diffusion in an efficient fashion because of its physical interception and chemical immobilization of polysulfides. Additionally, visualized adsorption tests by adding equal amount (20 mg) of NHG, HG, and rGO powders into Li2S6 solution were conducted (Figure S12). The Li2S6 solution were de-colored in 3 h by the dosage of NHG, implying competitively strong anchoring of polysulfides as compared with other reported host materials.39-42 In contrast, the solution with the addition of HG only faded to pale yellow after a long period of 48 h, whereas the solution containing pure rGO remained the same yellow as initial. The concentration differences of Li2S6 after adsorption for 48 h were further probed by ultraviolet-visible (UV-Vis) absorption measurements (Figure S13), with the absorption peak of Li2S6 in the visible light range completely vanished in the case of NHG dosage. This phenomenon confirms that the hierarchical porous structures indeed exert key efforts on the multi-level confinement of polysulfides, which is in close coordination with the strong entrapment of polysulfides by nitrogen dopants in the NHG frameworks. The NHG materials were further examined by STEM-EDS after Li2S6 adsorption (Figure 3d-g). The STEM image in Figure 3d reveals that the porous structure can be well preserved after adsorption of polysulfides. Corresponding EDS maps of carbon (Figure 3e), nitrogen (Figure 3f), and sulfur (Figure 3g) demonstrate a uniform distribution of anchored polysulfide species on N-doped graphene films. Moreover, XPS characterization of NHG powders after Li2S6 adsorption clearly manifests S peaks with deconvoluted signals of polysulfide, thiosulfate (S2O32-), and 10


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polythionate [O3S2-(S)x-2-S2O3] (Figure S14).43 This result verifies advanced anchoring ability of NHG frameworks with respect to sulfur species, which accords well with visualized experimental observations. To further substantiate the anchoring effects of NHG frameworks for polysulfides, the binding interaction between polysulfides and three different substrates, namely, pristine graphene, pyridine N- and pyrrolic N-doped graphene, were systematically investigated by DFT simulations. Figure 4a displays the calculated binding energies for S8 and Li2Sn (n = 1, 2, 4, 6, and 8), with optimized configurations shown in Figure S15. Compared to pristine graphene, N-doped graphene significantly increases the binding strength for polysulfides, especially for those containing Li (Li2Sn). Specifically, the binding energy of Li2S, Li2S2, Li2S4, Li2S6, and Li2S8 on pyridine N-doped graphene reaches 2.29, 2.40, 1.85, 1.73, and 1.99 eV, respectively. With regards to the pyrrolic N-doped graphene, the corresponding binding energy is calculated to be 1.68, 1.81, 1.40, 1.17, and 1.81 eV, respectively. Note that the binding energies for polysulfides are comparable to those on a variety of reported polar 2D materials, including MoS2, TiCl2, ZrCl2, NbS2, VS2, ZrS2, and TiS2.44 The underlying mechanism for the enhanced binding of Li2Sn on N-doped graphene can be visualized from the Bader charge transfer analysis. Herein the adsorption of S8 and Li2S4 on pyridine and pyrrolic N-doped graphene are selected as examples with the charge difference after adsorption plotted as Figure 4b (more details in Figure S16). As for the interaction between N-doped graphene and unlithiated S8, there is hardly any electron transfer from S8 to substrates, and the distance between S8 and the substrates remains almost the same (Figure S17). These results indicate that S8 adsorption involves purely van der Waals interaction and the binding strength is independent of nitrogen doping. In contrast, for Li2S4 case, electrons transfer occurs from S atoms to substrates 11


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(mainly N atoms) throughout Li atoms, leading to the stronger binding on N-doped graphene. The number of electrons accumulated at N atoms are calculated to be 0.59 and 0.47 for pyridine and pyrrolic N-doped substrates, respectively. Moreover, the bond lengths between Li and the nearest atoms from substrates (Table S1) also demonstrate enhanced binding ability for pyridine and pyrrolic N, in good agreement with the binding energy calculations. To sum up, the introduction of N dopants breaks up the neutral π network of pristine graphene, thereby effectively promoting the charge transfer from polysulfides to N-doped substrates with Li atoms serving as bridges. In operando Raman spectroscopic measurement was carried out to investigate the intermediate polysulfide species in the electrolyte during battery cycling. To achieve real-time monitor of discharge products, a designed cell based on a punched cathode in combination with an ultrathin, optical quartz window, was assembled. Figure 4d displays in operando Raman spectra of the electrolyte within the cathode side obtained at 0.2 C in the first cycle of discharging process with the presence of NHG separator, showing gradually attenuated signals pertaining to soluble Li2S8 (~119, 279, and 454 cm−1) and Li2S6 (~178, 339, and 397 cm−1) species in the course of discharging. This phenomenon indeed confirms that NHG framework enables the strong anchor of polysulfides at the separator and effective suppression of the shuttling effect in Li−S batteries (Figure 4c). To systematically evaluate the impact of NHG-based separator exerted on the electrochemical performances of sulfur/graphene (S/G) cathodes, typical coin cells with a sulfur content of 61.2 wt% (Figure S18) for the whole electrode and a sulfur mass loading of 1.4-1.6 mg cm−2 were assembled. For comparison, batteries with HG and rGO-coated separators were accordingly fabricated under the identical conditions. Figure 5a presents the representative cyclic voltammetry (CV) profile of the batteries with the NHG separator for the first cycle at a scan rate of 0.05 mV s−1. It is evident that 12


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two featured cathodic peaks appear at the potential range of 2.43-2.27 V (Peak I) and 2.07-1.83 V (Peak II), which can be attributed to the reduction of S8 to long-chain soluble polysulfides and reduction towards insoluble Li2S2 and Li2S, respectively. The characteristic single anodic peak at the potential range of 2.25-2.58 V is associated with the oxidation of Li2S to S8. To figure out the charge transfer resistances (Rct) of batteries incorporating NHG, HG, and rGO separators, electrochemical impedance spectroscopy (EIS) curves were collected (Figure S19). The battery incorporating an NHG separator achieves a lowest Rct value of 22.8 Ω as compared to that of HG separator (31.3 Ω) and rGO separator (57.9 Ω). Such highly conductive NHG frameworks would be beneficial to expediting the Li2S precipitation and recycling process especially under high-rates. The rate capacities of batteries with the NHG, HG, and rGO-coated separators were tested by elevating the current density from 0.2 C to 2 C for 5 cycles each and returning to 0.2 C (1 C = 1672 mA g−1). As shown in Figure 5b, when cycled at 0.2, 0.5, 1 and 2 C, the batteries with the NHG separator delivers capacities of 1474, 1358, 1175, 987 mAh g−1, respectively. When backing to 0.2 C, a high capacity of 1409 mAh g−1 can be achieved, demonstrating that NHG separator endows sulfur cathode with excellent rate capacities and stability. This is in stark contrast to the rate performances of batteries employing the HG and rGO separators. In response, Figure 5c discloses the galvanostatic charge/discharge curves of NHG- separator-integrated battery at various current densities. Despite of obvious polarization of the sulfur cathode at higher rates, the typical dual voltage plateaus located at ~2.3 V and 2.0 V could still be clearly observed, in good agreement with the stepwise reduction process of polysulfides revealed by the CV profile. Figure 5d compares the cycling performances of batteries with the NHG, HG, and rGO separators at 0.3 C. The initial discharge capacity of 1415 mAh g−1 could be gained with the NHG 13


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separator, suggesting a high sulfur utilization. After 100 cycles, it retains nearly 100% Coulombic efficiency and a reversible capacity of 1092 mAh g−1 with a capacity decay of only 0.23% per cycle, manifesting an outstanding cycling performance at the low current density. HG separator enables a slightly inferior initial capacity (1374 mAh g−1) and cycling stability (0.35% per cycle within 100 cycles) due to its physical confinement of polysulfides. In contrast, the battery incorporating the rGO separator otherwise delivers a much lower initial discharge capacity (1060 mAh g−1), which fades quickly for 100 cycles with a high decay rate (0.69% per cycle). It is realized that cycling at a lower current density would lead to more conspicuous shuttle effect than at high rates.45 As expected, evident disparity of cycling performance at 0.3 C has been witnessed especially between NHG and rGO separator based batteries. This can be ascribed to the synergetic effect of physical trapping and chemical anchoring of polysulfides by NHG that results in a restrained polysulfide shuttle. Additionally, the highly conductive networks of NHG could serve as a platform for high-efficiency conversion of polysulfides, thereby contributing to the excellent cycling performances. The cycling performances of batteries with the NHG, HG, and rGO separators at 1 C for 250 cycles are displayed in Figure 5e. The NHG separator-based cell delivers a high initial capacity of 1234 mAh g−1 with a low capacity decay of 0.15% per cycle, obviously outperforming those of HG and rGO based batteries. Such improved cycling performance of sulfur cathode stems from the strong entrapment and smooth conversion of polysulfides by the coating of NHG layer. Post-mortem inspection of the NHG separator after cycling at 1 C was conducted to probe its structural stability after long-term cycling, where the SEM image and corresponding EDS maps show intact porous NHG structure with a uniform distribution of sulfur-containing species (Figure S20). It is worth noting that the S/G hybrid cathode was used to evaluate the effect of NHG-based separator on the electrochemical 14


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performance of Li−S systems in this work because of its favorable conductivity and high material utilization. For comparison, pure sulfur cathode was also prepared by simply mixing 80 wt% sulfur, 10 wt% super P carbon black and 10 wt% LA132 binder. As expected, the pure sulfur cathode exhibited significantly lower initial discharge capacities of 885 mAh g-1 and 713 mAh g-1 with the NHG separator at 0.3 C and 1 C, respectively (Figure S21). Despite the employment of NHG-based separator, it is evident that the material utilization with respect to pure sulfur cathodes remains still quite low. Additionally, the effect of different mass loadings of NHG on the electrochemical performances was also investigated. Shown in Figure S22, the Li−S battery using NHG-based separator with a mass loading of 0.10 mg cm-2 delivers enhanced rate capability and cycling stability as compared to those based on mass loadings of 0.05 and 0.20 mg cm-2. This can be mainly attributed to the insufficient polysulfide trapping capability of the thinner NHG layer and sluggish Li ions transport in the thicker NHG film. To further highlight the superiority of NHG separator, prolonged cycling performance at a relatively high rate of 2 C was evaluated (Figure 5f). The battery presents an initial capacity of 868 mAh g−1; the great capability to restrain polysulfide shuttling has enabled the maintenance of discharge capacity at 403 mAh g−1 after 800 cycles, accompanying with an average capacity decay of 0.067% per cycle. The NHG separator based cells demonstrate good cycling stability under high current rates, which is crucial for practical Li−S batteries. In addition, the electrochemical performances with our NHG separator herein manifest high initial discharge capacity and capacity retention rate, which compares favorably with many state-of-the-art carbon-based separators reported thus far (Table S2). In view of commercialization requirements, operation stability with higher sulfur mass loadings 15


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was further attempted involving our NHG separator. Along this line, S/G cathodes with sulfur loadings of 3.8 and 7.2 mg cm−2 were fabricated. As shown in Figure 6a, the batteries with sulfur mass loading of 3.8 mg cm−2 delivers an initial capacity of 1131 mAh g−1 at 0.3 C with a capacity decay rate of 0.97% per cycle within 30 cycles. At the sulfur loading of 7.2 mg cm−2, the battery exhibits an initial capacity of 805 mA h−1 with a capacity retention of 84.1% at 0.1 C after 30 cycles. Note that the stability of S/G cathodes based on the NHG separator is not extraordinarily high; but there remain considerable opportunities in the future to optimize the electrochemical performances. From a practical perspective, usage safety under elevated temperatures has also long been a big concern for Li–S batteries. In light of this, we simply evaluated the rate performance of battery with the NHG separator under an elevated temperature of 50 oC. As shown in Figure 6b, the battery delivers capacities of 1206, 1072, 923 mAh g−1 at 0.5, 1, and 2 C, respectively. When the current density was shift back to 0.5 and 0.2 C, it retains high reversible capacities of 1110 and 1170 mAh g−1, respectively, manifesting favorable heat-resistant rate capability under an elevated temperature. Figure 6c displays corresponding galvanostatic charge-discharge curves at 0.5, 1, and 2 C under 50 o

C. In addition, to envisage practical applications in high-energy-density wearable energy storage

fields, a proof-of-concept wearable battery device by integrating segmented Li-S cells is shown (Figure 6d), where a light-emitting diode (LED) can be steadily powered even under various bending states.

Conclusions In summary, we have developed a bio-templating CVD approach for the direct synthesis of 3D NHG architectures via the employment of naturally abundant diatomite as the growth substrate. The 16


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thus-derived NHG perfectly preserve the hierarchical structures of original diatom frustules, which is featured by the fascinating nepenthes-like architectures with multifarious macro/meso-pores and inter-connected channels. Uniform nitrogen doping with tunable dopant concentrations can also be realized by such CVD process, manifesting an efficient performance for polysulfide trapping. The high surface area, devious inner-channel structures and abundant nitrogen doping of NHG material were utilized as polysulfide scavenger with synergetic bi-functionality of physical confinement and chemical immobilization for polysulfides. A simple coating of the as-obtained NHG materials onto commercial Celgard separator renders the construction of a thin, functional NHG separator. The polysulfide trapping mechanisms of such NHG-based separator have been systematically revealed by structural and elemental analysis, permeation and adsorption tests, DFT simulations, as well as electrochemical performance evaluations. Moreover, the NHG separator based Li–S batteries have demonstrated favorable application potentials in terms of high sulfur loading and elevated temperature operations. This strategy reported in the present study could offer insights into designing practical energy-storage devices with bio-inspired 3D hierarchical structures.

Methods CVD synthesis of NHG powders. In a typical procedure, the as-received diatomite powders (Alfa Aesar) were thoroughly cleaned with nitric acid (2 mol·L−1) and deionized water to remove the metal impurities. After further filtration and drying treatment, the diatomite powders were loaded into a 4-inch tube furnace for CVD process. The CVD system was pumped with a base pressure of 1 Pa and then purged with 60 sccm H2 and 150 sccm Ar as carrier gases. The diatomite powders were 17


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heated to 900 oC and exposed to evaporated pyridine precursors, which was placed in a steel cylinder connected to the growth tube. After synthesis, the system was cooled down to room temperature under H2 and Ar. The obtained NG/diatomite powders were immersed and stirred in a diluted hydrogen fluoride (HF) etching solution (HF: H2O: EtOH: 1: 7: 2; v/v/v) at room temperature to remove the silica templates. The pure NHG powder was obtained by virtue of thorough rinsing with deionized water and freeze-drying. Preparation of sulfur-graphene (S/G) hybrids. GO sheets were synthesized from natural graphite by employing a modified Hummers’ method. The as-synthesized GO sheets were kept under 950 oC for 5 h in an Ar/H2 flow for preparing rGO. Sublimed sulfur and rGO with a mass ratio of 4:1 were mixed and heated to 155 oC for 12 h in a sealed vessel to prepare S/G hybrids. Preparation of NHG-based separator. The NHG separators were achieved by vacuum filtration of the NHG dispersion. Typically, 100 mg of NHG was dispersed into 100 ml of 0.02 wt% Triton X-100 aqueous solution by ultrasonication for 20 min. Then 20 mL of the NHG dispersion, 2 mL of 0.1 wt % LA132 aqueous binder were dispersed into 78 mL of the deionized water/ethanol mixture with a volume ratio of 28:50, followed by vacuum filtration onto Celgard 2400 separator. After rinsing and vacuum drying, the NHG separator was finally obtained and the mass of NHG on separator was ~0.1 mg cm−2. For comparison, HG and rGO separators were also gained by the similar procedure. Electrochemical measurements. Electrochemical tests were operated on coin-type (LIR 2032) batteries. S/G hybrids were firstly mixed with super P carbon black and LA132 aqueous binder (mass ratio 8:1:1) with the assistance of a twin-shaft mixer. The as-prepared slurry was coated onto an Al foil. The obtained laminates were punched into circle discs with a diameter of 13 mm and dried under 50 °C in a vacuum atmosphere for 12 h. As a reference, the pure 18


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sulfur cathodes were prepared by casting the slurry containing 80 wt% sulfur, 10 wt% super P carbon black and 10 wt% LA132 binder onto Al foil, followed by the same post-treatment procedure. The batteries were fabricated with S/G cathode, NHG-modified separator, lithium anode, and electrolyte in an argon-filled glove box. The electrolyte was 1.0 M LITFSI in a mixture of dimethoxyethane and 1,3-dioxolane (volume ratio 1:1) with 5 wt% of LiNO3 as the additive. The sulfur content was 76.5 wt% in S/G hybrids and 61.2 wt% in the entire cathode. The electrolyte/sulfur ratio was 10:1 µL mg−1 for typical-sulfur-loaded (1.4-1.6 mg cm−2) cathodes and 7.5:1 µL mg−1 for high-sulfur-loaded (~7.2 mg cm−2) cathodes. Characterizations. Characterizations were performed with the aid of SEM (Hitachi SU8010), AFM (Veeco Nanoscope IIIa, tapping mode), Raman spectroscopy (Horiba, HR Evolution, 532 nm), TEM (FEI Titan Themis G2, acceleration voltage 30-300 kV), XPS (Escalab 250Xi Spectrophotometer using a monochromatic Al Kα X-ray source), UV-vis absorption spectroscopy (Perkin-Elmer Lambda 750S UV/Vis/NIR Spectrophotometer), N2 adsorption/desorption isotherm analyzer (Tristar II 3020), TGA (7300 Thermogravimetric/Differential Thermal Analyzer), and four-probe resistance measuring system (Guangzhou 4-probe Tech Co. Ltd., RTS-4). A Land CT2001A battery testing system was utilized to test the electrochemical performances of batteries including rate capacities, charge-discharge curves and cycling performances. The cyclic voltammograms (CV) and the electrochemical impedance spectroscopy (EIS) were measured on an Autolab potentiostat PGSTAT302N. The scan rate for CV test was 0.5 mV s−1 in a voltage range from 1.7 to 2.8 V. The EIS were measured in the frequency range of 100 kHz to 0.01 Hz. DFT simulations. Vienna ab-initio simulation package46,47 was employed, with the ion-electron interactions and the exchange correlations treated by projector augmented wave potentials48,49 and 19


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the Perdew-Burke-Ernzerhof functional.50 The energy and force convergence criteria were set to be 10−5 eV and 0.01 eV/Å, while plane wave cutoff was chosen as 400 eV. A 7×7×1 supercell for (N-doped) graphene was constructed as our models, and for N-doped graphene, three pyridine N and three pyrrolic N were introduced around vacancies. A 1×1×1 k-point sampling was used for relaxation of polysulfide molecules, while a 3×3×1 sampling was applied for substrates with and without absorbed polysulfides. The energies of the polysulfide absorbates (Eabs), the underlying substrates (Esub), and the polysulfide-absorbed substrates (Etotal) were respectively calculated, using van der Waals interaction by Tkatchenko-Scheffler method.51 The binding energy Eb can then be calculated by ‫ܧ‬௕ = ‫ܧ‬௔௕௦ + ‫ܧ‬௦௨௕ − ‫ܧ‬௧௢௧௔௟ .

ASSOCIATED CONTENT Supporting Information Experimental set-up, sample characterizations, visualized Li2S6 adsorption tests, and detailed electrochemical performances. This material is available free of charge via the Internet at http://pubs.acs.org. The authors declare no competing financial interest.

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] (J. Y. Sun) *E-mail: [email protected] (L. Zhang) *E-mail: [email protected] (Z. F. Liu) 20


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†These authors contributed equally to this work.

ACKNOWLEDGMENTS This work was financially supported by the MOST (2016YFA0200103), the National Natural Science Foundation of China (51702225, 51675275, 21473119, 51520105003, and 51432002), the Beijing Municipal Science and Technology Commission (No. Z161100002116020), and the Jiangsu Youth Science Foundation (BK20170336). Q.C.L., Y.Z.S., L.Z., J.G., Z.X., Z.N.T., N.W., M.H.R., J.Y.S., and Z.F.L. acknowledge the support from Suzhou Key Laboratory for Advanced Carbon Materials and Wearable Energy Technologies, Suzhou, China. Y.Z.S acknowledges the support from the Postgraduate Research & Practice Innovation Program of Jiangsu Province (KYCX17-2023). X.L.Z. and J.Y.S. acknowledge the support from the Thousand Youth Talents Plan of China. R.Z.X. and X.L.Z. acknowledge the Development and Reform Commission of Shenzhen Municipality for the development of the “Low-Dimensional Materials and Devices” Discipline. The computations were carried out at National Supercomputer Center (TianHe-1A) in Tianjin.

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Figure 1. Bio-templating CVD growth of N-doped hierarchical graphene. (a) Schematic illustration of bio-templating CVD growth of N-doped graphene on a diatom frustule. (b) TEM image of an NHG architecture, showing two types of pores located at the edge and at the center of the biomorphic graphene. (c) AFM image of an NHG on SiO2 substrates with the thickness ranging from 50-100 nm along the porous frameworks. (d) False-color SEM image showing the various pores in the NHG material. (e) Nitrogen adsorption–desorption isotherms of NHG and diatomite powders. (f) Photograph of the as-obtained NHG film on Celgard separator with excellent flexibility. (g) Spatial distribution of the sheet resistance (3 cm × 3 cm), with the inset showing the distribution of sheet resistance (collected from 81 points) of the NHG film. (h-k) HRTEM images of the pore structures of NHG, detailing upper macro-pores, interconnected channels, and inner meso-pores, 28


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respectively. Note that the morphology of such biomorphic NHG resembles that of nepenthes displayed in the inset of j.

Figure 2. Structural and elemental characterizations of N-doped hierarchical graphene. (a) OM image of an individual NHG architecture grown at 900 oC with evaporated pyridine under a total pressure of 1.0×102 Pa. (b) Raman mapping image (G peak from 1460 to 1650 cm-1) of the NHG shown in a. (c) Raman spectra of NHG and HG powders grown at 900 oC in comparison with rGO powders. (d, e) XPS spectra of the C 1s and N 1s signals for NHG, respectively. (f) HRTEM image of NHG with a thickness of five layers. (g-i) Corresponding STEM image (g) and elemental maps of carbon (h) and nitrogen (i). 29


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Figure 3. Polysulfide permeation measurements. (a) H-shaped permeation device equipped with NHG separator, enabling the effective suppression of polysulfide diffusion for over 48 h. (b) H-shaped permeation device equipped with HG separator, showing gradual permeation within 24 h. (c) H-shaped permeation device equipped with rGO separator, indicating the failure of polysulfide blockage occurred just within 1 h. (d-g) Corresponding STEM image (d) and elemental maps of carbon (e), nitrogen (f), and sulfur (g).

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Figure 4. Binding affinity between nitrogen-doped hierarchical graphene and Li2Sx. (a) Binding energy between Li2Sn and pyridine N, pyrrolic N, and graphene. (b) The charge transfer after the absorption of S8 and Li2S4 on pristine graphene, pyridine N doped, and pyrrolic N doped substrates. The pink (green) isosurfaces represent charge gain (loss) in real space, with isosurface value set as 0.003 e. The charge differences of Li, S, and N atoms are also marked, with positive (negative) values denoting the number of charge gained (lost). (c) Schematic cell configuration of Li–S batteries using NHG-decorated separator. (d) In operando Raman spectra based on NHG separator collected upon the first discharge at 0.2 C.

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Figure 5. Electrochemical performances of the batteries with NHG separator. (a) CV profile of the battery with the NHG separator in a potential range from 1.7 to 2.8 V. (b) Rate performances of the batteries with NHG, HG and rGO separators. (c) Galvanostatic charge/discharge curves of the battery with NHG separator at different rates. (d, e) Cycling performances of the batteries with NHG, HG and rGO separators at 0.3 C for 100 cycles (d) and 1 C for 250 cycles (e). (e, f) Cycling performance of the batteries with NHG separators at 2 C for 800 cycles.

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Figure 6. High-sulfur loaded, heat-resistant and wearable batteries with NHG separators. (a) Cycling performances of the batteries with the NHG separators at 0.3 C with the sulfur loadings of 3.8 mg cm–2 (upper panel) and at 0.1 C with the sulfur loading of 7.2 mg cm–2 (lower panel). (b) Rate performance of the battery with the NHG separator under 50 oC. (c) Galvanostatic charge/discharge curves at different rates under 50 oC. (d) Demonstration of wearable Li–S batteries based on the NHG separator for powering LED under different bent states.

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TOC graphic

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Bio-Templating Growth of Nepenthes-Like N-Doped Graphene as

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