High-Performance Hard Carbon Anode: Tunable Local Structures and

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High-performance Hard Carbon Anode: Tunable Local Structures and Sodium Storage Mechanism Yu Jin, Shixiong Sun, Mingyang Ou, Yi Liu, Chenyang Fan, Xueping Sun, Jian Peng, YuYu Li, Yuegang Qiu, Peng Wei, Zhi Deng, Yue Xu, Jiantao Han, and Yunhui Huang ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.8b00354 • Publication Date (Web): 26 Apr 2018 Downloaded from http://pubs.acs.org on April 26, 2018

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ACS Applied Energy Materials

High-performance Hard Carbon Anode: Tunable Local Structures and Sodium Storage Mechanism Yu Jin, Shixiong Sun, Mingyang Ou, Yi Liu, Chenyang Fan, Xueping Sun, Jian Peng, Yuyu Li, Yuegang Qiu, Peng Wei, Zhi Deng, Yue Xu, Jiantao Han* and Yunhui Huang State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science, Wuhan 430074, P. R. China * E-mail: [email protected] KEYWORDS: hard carbon, structure model, sodium storage mechanism, local structure, sodium ion battery

ABSTRACT: Hard carbon (HC) is one of the most promising anode materials for sodium-ion batteries (SIBs) due to its suitable potential and high reversible capacity. At the same time, the correlation between carbon local structure and sodium-ion storage behavior is not clearly understood. In this paper, the two series of HC materials with perfect spherical morphology and tailored microstructures were designed and successfully produced using resorcinol formaldehyde (RF) resin as precursor. Via hydrothermal self-assembly and controlled pyrolysis, RF is a flexible precursor for high-purity carbon with a wide range of local structure variation. Using these processes, one series of five representative RF-based HC nanospheres with varying degrees

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of graphitization were obtained from a RF precursor at different carbonization temperatures. The other series of HC materials with various microscopic carbon layer lengths and shapes was achieved by carbonizing five-RF precursors with different crosslinking degrees at a single carbonization condition (1300 °C and 2 h). Based on the microstructures, unique electrochemical characteristics, and atomic pair distribution function (PDF) analyses, we proposed a new model of “three-phase” structural for HC materials and found tri-region Na-ion storage behavior: chemi-/physisorption, intercalation between carbon layers, and pore-filling, derived from the HC phases, respectively. These results enable new understanding and insight into sodium storage mechanism in HC materials and improve the potentiall for carbon-based SIB anodes.

Introduction Lithium-ion batteries (LIBs) are widely used as rechargeable batteries for portable electronics and electric vehicles.1, 2 However, lithium is expensive and batteries made from it are unsafe, which inhibit LIBs use in large-scale energy storage systems.3, 4 Sodium and lithium have similar chemical properties, yet sodium is an abundant and cheap resource.5-7 For these reasons, sodium is currently the most desirable element to replace lithium in the rechargeable batteries for energy storage systems. And advanced electrode materials are key to developing sodium-ion batteries. In the past few years, many SIB cathode materials have been investigated, including layered oxides,8 tunnel-type oxides,9 polyanionic sodium salt,10, 11 and Prussian blue analogues.12-14 But equally importantly, anode materials are very limited due to the large ionic radius of Na (1.02 Å compared to 0.76 Å for Li). Because of sodium’s large ionic radius, the carbonate electrolyte in sodium-based anodes can’t use the graphite (interlayer d-spacing of 0.335 nm) widely used in LIBs.15-17 Other anode materials have been extensively investigated, including nongraphitizable


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HC,18-20 alloys,21 oxides,22 and organic compounds.23 Most of them suffer huge volume expansion during sodium-ion insertion, resulting in irreversible capacity fading. Currently, HC materials with randomly oriented graphitic layers are among the most promising candidates for SIB anode materials due to their high reversible capacity near 350 mAh g-1, their suitable average potential of ~0.15 V versus Na/Na+, and their excellent cycling stability.24,

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Furthermore, the fundamentals of the sodium storage mechanism in HC materials have also attracted much attention from researchers. In 2000, Stevens and Dahn first presented the “card-house” model for HC materials. The model displays two kinds of structural features,26 i.e., graphite-like microcrystallites and an amorphous region. The team revealed sodium-ion storage in glucose-pyrolyzed HC and assigned the sloping region capacity (SRC) to sodium-ion insertion between the graphene sheets, and the low-plateau region capacity (PRC) to sodium filling/plating into nanopores (herein referred to as the “intercalation-adsorption” mechanism). Meanwhile, Komaba et al.27 corroborated this mechanism by ex situ X-ray diffraction (XRD) and angle X-ray scattering (SAXS), and further confirmed that SRC was related to sodium-ion insertion between the graphene sheets. Later, Cao et al.28 presented exactly the reverse of the adsorption-intercalation mechanism: Based on polyaniline (PANI)-pyrolyzed HC structural and electrochemical characteristics, SRC was assigned to sodium-ion adsorption on surface active sites and PRC was related to sodium-ion insertion into graphite-like microcrystallites. The resulting literature about the two different mechanisms focus on the type of materials, the synthetic conditions, and the special characterization techniques, such as in/ex situ XRD, X-ray photoelectron spectroscopy (XPS), temperature-programmed desorption coupled with mass spectrometry (TDP-MS), highresolution transmission electron microscopy (HR-TEM), and so on.29, 30


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Recently, Ji et al.31-33 expanded the view of sodium storage mechanism using the galvanostatic intermittent titration (GITT) technique. They found that the glucose-pyrolyzed HC features a higher Na+ ion diffusion coefficient (DC) in the sloping region rather than in the plateau region, mostly due to the higher accessibility of the surface-active sites compared to interlayer space. Thus, the team proposed that sodium storage mechanism in the sloping (0.1 – 1 V) part of the galvanostatic discharge curve could be explained as being due to sodium-ion storage at defective sites, rather than due to intercalation within graphene sheets. (Presently, there is general agreement on this point.) Simultaneously, the team observed that the Na+ ion DC in the plateau range has a minimum value at ~0.05 V (later see our value at ~0.03 V), the same voltage where dQ/dV value reaches a maximum. From 0.05 V to the cutoff voltage, the Na+ ion DC gradually increases, while the dQ/dV value decreases. This suggests that the sodiation mechanism is changing in the plateau range before reaching the cutoff potential. Therefore, Ji’s team postulated that the storage mechanism in the low-voltage plateau region is related to Na+ ion intercalation between graphene sheets followed by Na filling pores in the carbon structure (herein referred to as the “intercalation-filling” mechanism). But this is discrepant with the “card-house” model. A similar experimental phenomenon was also observed in our previous work.34 Interestingly, the Na+ ion DC of the plateau region is not continuously decreasing, and it has an inflection point located at ~0.03 V. More importantly, utilizing this point where the Na+ ion DC gives a minimum, PRC can be divided into two parts. Therefore, some crucial questions are raised: Are there rules between the two parts? If SRC is also considered, what kinds of relationships exist? Are the three parts associated with three respective characteristic behaviors: absorption, intercalation, and pore-filling?


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It would be bold to hypothesize that the three sodium storage behaviors relate to three distinct structural features of HC materials. So, we first present a new “three-phase” structure model for HC, i.e., defect surface containing active sites and edges, graphene sheets with variable dspacing forming graphite-like microcrystallites, and confined spaces formed by the randomly stacked graphite-like microcrystallites (called “amorphous region” in the card-house model). We present the three-phase model first to further establish an understanding of sodium storage mechanism corresponding to three sodium-storage behaviors: absorption, intercalation, and porefilling. It’s worth noting at the outset that, coincidentally, two parts of PRC divided by the inflection point of Na+ ion DC can be assigned to intercalation and pore-filling, respectively. Crucially, the three-phase structure model assigns SRC only to Na-ion storage in defect surface, unrelated to HC internal constructions, while ascribing the two parts of PRC to graphene sheets and confined spaces, respectively, as illustrated in Figure 1. Actually, much literature has mentioned the three distinct chemical environments for Na-ion storage, one of which is edge/defect sites on the edges of graphite-like microcrystallites. But none have claimed that SRC is only assigned to the surface of HC particles. Therefore, this point is key to understanding the new three-phase structure model of HC materials for sodium-ion storage. In this study, RF resin was chosen as HC precursor. Via hydrothermal self-assembly and controlled pyrolysis, flexible RF can be used to make carbon material with a wide range of tunable local structures. First, a typical RF precursor (noted RF-H2O) was carbonized at 800 – 1600 °C for 2 h to obtain HC materials with different interlayer d-spacing (d002) and length of graphitic layers, signed as HC-T (T is temperature value, i.e., HC-800, HC-1100, …, HC-1600). Then, one series of RF precursors – marked as RF-Solvents (RF-Ss), i.e., RF-H2O, RF-CH3OH, RF-C2H5OH, RF-C3H7OH, and RF-C4H9OH – with different crosslinking degrees were achieved


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using their sensibility to synthetic solution polarity. The precursors were carbonized at the same carbonization condition of 1300 °C for 2 h to obtain different HC materials (named HC-Ss, i.e., HC-H2O, HC-CH3OH, HC-C2H5OH, HC-C3H7OH, and HC-C4H9OH, respectively) with different internal constructions, containing contents, and sizes of graphene sheets and confined spaces, but roughly with the same sphere morphology, particle size, and surface status. Results and Discussion Figure 2 shows XRD patterns and Raman spectra of RF-H2O-based HC-Ts carbonized at different temperatures and RF-Ss-based HC-Ss carbonized at the same temperature of 1300 °C, respectively. All XRD patterns exhibit two broad peaks around 23° and 43°, which can be assigned to the crystallographic planes of (002) and (101) in the disordered carbon structure. The weak and broad peaks indicate that these materials are non-graphitized carbon with a low degree of graphitization. Figure 2a shows that with increasing carbonization temperatures, the peak (002) of HC-Ts becomes sharper and significantly shifts to a high angle, which reveals not only a local-structure development forming a short-range ordering, but also decreased interlayer dspacing (002). To accurately understand HC structures, we must consider differences in the structure and chemical composition of organic precursors – especially during the carbonization process – that decompose and form planar graphene sheets with remaining the C atoms. If the organic precursor sufficiently crosslinks, then those planar graphene sheets cannot align vertically. Such carbon materials are difficult to graphitize, even at high carbonization temperatures, and thus are called “hard carbon.” The XRD results of HC samples derived from RF-Ss with different crosslinking states are very surprising (see Figure 2b). Surprising in that the peak (002) of HC-Ss


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has almost the same shape and position, which indicates the interlayer d-spacing of parallel graphene sheets is very close. For HC-Ss under high calcination temperature, the length and amount of the planar graphene sheets grow gradually into the bigger graphite-like microcrystallites, and they are stacked randomly, separated by profuse small-size confined spaces. Although the five HC-S samples have the same interlayer d-spacing, the size and content of their graphite-like microcrystallites and confined space highly depend on the crosslinking degree of the RF precursors. Figure 2b and d show Raman spectra of HC-Ts and HC-Ss, respectively, and all the Raman spectra exhibit two peaks: 1350 cm-1 represents A1g vibration mode of sp2 carbon rings caused by defects, and 1580 cm-1 represents E2g vibration mode of sp2 carbon atoms. Further Raman spectra details and fitting results are shown in Figure S4 and S5. Usually, ID/IG can be applied to indicate the disordering degree of the carbon materials.35 We have deconvoluted the Raman spectra and got the ratio of ID/IG. According to the formula:

.31, 36, 37, we calculated the parameters of all the HC graphitic structures. Table 1 shows that the degree of graphitization obviously increases and the interlayer d-spacing remarkably decreases with increasing carbonization temperature. However, these two parameters reveal a slightly increasing trend with enhanced solvent polarity from C4H9OH to H2O, which well comport with the XRD results. Figure S1a shows that the FT-IR spectra of RF-Ss are similar with typical RF resin characteristics. The characteristic C-O-C absorption band is not observed, which suggests that


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the methylene group is mainly formed instead of an ether bond in polycondensation of resorcinol and formaldehyde. The biggest differences between the characteristic bands are located at 1610 cm−1 (aromatic ring C=C), 1460 cm−1 and 1480 cm−1 (methylene peak).38 We find that the methylene and C=C content of RF-H2O are the highest among RF-Ss. Significantly and based on the principle of the dissolution in the similar material structure, the role of high-molecularweight polymers in enhancing polar solvency can be explained in terms of polymer coil overlap, which results in a large increase in the degree of polymerization and crosslinking, and therefore RF-H2O shows the highest crosslinking degree. During the carbonizing process, the crosslinking structure hinders the movement of the molecular chain, impeding the carbonization process, also preventing the growth of ordered graphitic carbon layers. Figure S1b shows TG curves both of RF-H2O and RF-C4H9OH. After heating to 1000 °C in argon, the RF-H2O leaves more residue than RF-C4H9OH, which supports the FT-IR results. In order to further observe the structural differences between RF-H2O and RF-C4H9OH, ex situ FT-IR was carried on from 300 °C to 600 °C. Figure S2 show that both of FT-IR spectra exhibit no obvious change below 400 °C, which is mainly due to the process of dehydration and dehydration of small molecules. When up to 500 °C, a significant change happens, which suggests that the stability of the methylene structure as shown at 1450 cm-1 and 1470 cm-1 remarkably influences carbon graphitization at higher temperatures. Figure S2a and b show that the structure of RF-H2O is much more stable than that of RF-C4H9OH, which suggests that RFH2O with a higher crosslinking degree has more difficulty rapidly forming carbon-ring conjugate networks. Due to a low degree of crosslinking, the structure of methylene in RF-C4H9OH was destroyed completely at 600 °C, thus causing less residue for RF-C4H9OH.


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Figure 3 shows HR-TEM and selected area electron diffraction (SAED) images of the intuitive developed graphitic crystallites at different carbonization temperatures. With increasing carbonization temperature, the ordered graphitic crystallites gradually form and grow. The insert of Figure 3a shows that the diffraction ring of SAED at 800 °C is not clear enough, and then the diffraction ring becomes increasingly apparent, demonstrating an improved ordering structure of HC-Ts. Figure 4 shows HR-TEM and SAED images of HC-Ss. With increasing solvent polarity of hydrothermal self-assembly, we observe carbon layer lengths shortening, and their graphitelike microcrystallites average curvature radius decreases, with the number of confined spaces (surrounded by stacked graphite-like microcrystallites) with a size of less 1 nm increasing, as shown in Figure 3c, which might relate to the length of carbon layers with respect to other carbonization temperatures, and that well comports with the results of XRD. The change of ordered structure and the evolution of graphite-like microcrystallites and confined spaces in HCSs may greatly influence the electrochemical properties of HC materials, which will be discussed later. Figure 3f shows that the morphology of RF-H2O is uniform, perfectly spherical with a diameter of ~1.5 um. After carbonization, the diameter of HC-H2O spheres was reduced to ~600 nm (Fig. 3g). Figure 4(f - i) reveal SEM images of RF-Ss. The results show that the sizes of RFSs were reduced, and the spherical morphology was partially destroyed by the growth of alkyl chain length of the hydrothermal solvents, i.e., polarity depression. This phenomenon could be ascribed to water promoting the polymerization rate of the RF precursor: The hydrogen bonding between water and reactants is fairly strong, and thus the reactants can quickly gather into small emulsion droplets and aggregate into perfect homogeneous spheres. The spherical shape has great advantages of high tap density, low specific surface area, and stable structure. The longer


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alkyl chains make the reaction solution more viscous, and then the reactants cluster irregularly. Furthermore, the effect of the reaction time in the Stöber process on the morphologies of RF resins were also studied. As shown in Figure S7, when the co-precipitation time is 20 h, and hydrothermal time is 30 h, the morphology of RF resins is the most uniform. If the reaction time is too long, the RF resin tends to reunite. Figure S9 shows that the pH value also greatly affects resin morphology, and a high pH value is favorable for obtaining large particle size. Figure 5a and c show the galvanostatic discharge/charge curves of HC-Ts and HC-Ss with a current density of 20 mA g−1, respectively. The specific capacities of HC-Ts have a maximum value of 310 mAh g–1 at 1300 °C, with excellent capacity retention of ~300 mAh g−1 after 100 cycles (Figure 5b). As referred to in previous reports, the whole reversible capacity was usually divided into two parts, SRC and PRC, based on the basic shape of charge/discharge curves. However, according to the variation and evolution mode of Na+ ion DC, the whole capacity can be divided into three parts, absorption capacity (AC), interlayer intercalation capacity (IC), and filling-pose capacity (FPC) (later, see Figure 6 and Figure 7). As shown in Figure 5a, with obvious trends for HC-Ts from region to region with increasing carbonization temperature, SRC clearly decreases while PRC first increases and then decreases, with a maximum value at 1300 °C as a result of the continuous structural change of the RF-H2O precursor at the different carbonization conditions. During the early stage of carbonization, below 800 °C, RF-H2O precursor decomposes and emits gases that contain carbon, such as CO and CH4, the RF-based carbon spheres are terminated predominantly with many heteroatoms, such as H or O, and a lot of defected graphitic sheets on spherical surfaces. With elevating carbonization temperature, the active sites tapped on these surfaces decrease significantly, resulting in a reduced SRC, and the carbon layer d-spacing, ~0.38 nm, shrinks gradually while PRC shows a trend of increasing first


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and then decreasing, which suggests that the Na-ion storage in the plateau area could not be a single step. In contrast, Figure 5c shows a distinct trend that SRCs for HC-Ss are very close (the sloping curves approximately overlaying), while PRC appears to trend toward continuous increase with increasing solution polarity. Furthermore, the CV curves of HC-H2O and HCC4H9OH at a scan rate of 0.5 mV s–1 in a voltage range of 0 – 2 V are shown in Figure S6a, and the electrochemical properties of other HC materials assembled in different Stöber time, PH value, and hydrothermal temperatures are shown in Figure S (8 – 11). For insight into the mechanism of the sodiation/desodiation process in HC anodes, GITT was used to measure Na+ ion DC. During the GITT tests, sodiation/desodiation was carried out at a constant current density of 15 mA g-1 for an intermittent 30 min followed by an open-circuit stand for value (Es). All the Na+ ion DCs are determined by solving Fick’s second law of diffusion through a series of assumptions and simplifications as showed in supporting information.31,

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Figure 6 and Figure 7 show the GITT curves of HC-T and HC-S anodes,

respectively, with the first and second cycles from 0 to 2.5 V. All the GITT results show a similar trend that the range of the discharge curve is theoretical divided into three available intervals according to the variation law of Na+ ion DCs, based on our new three-phase model, which can be assigned to AC (1.0 – 0.1 V), IC (0.1 – 0.03 V), and FPC (0.03 – 0 V), respectively. Meanwhile, the three-part capacities could relate to those three-phase structures, defect surface, graphene sheets, and confined spaces in HC materials, and the change rules of these three parts of HC-T and HC-S anodes are shown in Figure 6f and Figure 7f, respectively. First, Figure 6f shows the ACs are decreasing with increasing carbonization temperature, owing to losing active sites, such as edges, defects, and functional groups containing O, H, etc. on the surfaces of HC-T spheres at high temperatures. In order to prove that AC is mainly related to the


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surface status of HC spheres, and independent of their internal structures, we design and synthesize a series of HC-Ss with significantly different internal structures, i.e. the size and content of graphite-like microcrystallites, meanwhile forming the corresponding confined spaces, because the higher crosslinking degree of RF precursors are more likely to have greater numbers and smaller sizes of graphite-like microcrystallites in HC-Ss spheres, while with significantly similar surface status due to the same calcination process. Importantly, Figure 7f shows the HC-S anodes give almost the same ACs, which fully proves our hypothesis is correct, that AC (i.e. SRC) is mainly related to the surface status of HC spheres. Second, based on the results of TEM and XRD, the carbon layers and graphite-like microcrystallites of HC-Ts show a process of formation and growth with increasing carbonization temperature. Meanwhile, the ICs increase gradually, as shown in Figure 6f. However, the FPCs have a maximum value at 1300 °C, which is due to an evolution of the gradual formation of confined spaces that then shrink with rising temperature. Figure 8a shows the schematic illustration of the change rule of the three-part capacities in HC-Ts. In addition, Figure 7f shows that the FPCs of HC-Ss increase with increasing crosslinking degree of RF-Ss, which is attributable to the increased amount of confined spaces. However, the ICs interestingly have a maximum value at RF-C2H5OH, which is due to a balance of Na+ ion storage capacity between numbers and size of graphite-like microcrystallites. In the beginning, with increasing crosslinking degree, the ICs are dominated by the increasing numbers of graphite-like microcrystallites; but subsequently, graphite-like microcrystallites become too small so that their sodium capacities dropped sharply. Figure 8b clearly shows the whole process of embedding Na ions into HC anodes. First, Na ions are captured in active sites on the surface of HC particles, which corresponds to an AC (0.1

– 1.0 V). As Na-ion content increases in the active sites, a large repulsive force is generated


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among Na ions, which hinders further adsorption of Na ions, resulting in slowly declines of Na+ ion DC. Meanwhile, it pushed Na ions between the graphene sheets of graphite-like microcrystallites with larger d-spacing (0.36 – 0.38 nm), which is more suitable for Na+ ion insertion and diffusion than that of graphite. While the interlayer diffusion of Na+ ions is still a little harder than surface absorption, it has been proven that Na-graphite compound has difficulty forming stable graphite intercalation compounds, resulting in sharp declines of Na+ ion DC while switching from surface absorption to interlayer intercalation.16,

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Meanwhile, the constant

intercalation of Na+ ions forming graphite intercalation compounds to overcome the higher activation energy barrier eventually leads to diffusion of Na+ ions into the parallel carbon layers of graphite-like microcrystallites (0.1 – 0.03 V). Finally, after graphite-like microcrystallites reach saturation point at a lower voltage of ~0.03 V, Na ions begin to enter into many of the confined spaces to resist the higher energy barrier for electron transfer and electrochemical reaction, causing DC begin to rise at a lower voltage of ~0.03 V.

In order to further clarify sodium storage mechanism, ex situ atomic pair distribution function (PDF) analysis was carried out at the different discharge states (see Figure 9b), and the PDF profiles are shown in Figure 9c. Significantly, Figure 9d shows the differential PDFs (DPDFs) between every two adjacent points, which can exhibit fine structural changes of HC anodes during the different discharge states. The D-PDF of 0.6 V and 0.1 V does not show more obvious features, indicating that the Na-ion adsorption didn’t significantly change structure in the HC anode. However, the D-PDF of 0.1 V and 0.03 V shows a tiny change at a low r range (r ﹤5 Å), combining the in situ XRD results, which could come from the distracted interlayer dspacing caused by the insertion of Na ions into the grapheme sheets of graphite-like microcrystallites. Interestingly, with further discharging, the D-PDF of 0.03 V and 0.01 V


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demonstrates a significant characteristic change at a larger range r (r﹤10 Å), which is similar to that of the D-PDF of 0.01 V and -0.02 V at a deeper discharge state, resulting in a plating of metallic Na. Therefore, the results suggest possible existing Na clusters in the discharge state of 0.03 V and 0.01 V, which proves the reliability of our three-phase structural model for HC materials to interpret sodium storage mechanism (shown in Figure 8b), and the phenomenon is also supported by previous studies.40, 42, 43 A large amount of sodium ions enter in the pore space (assigned as filling-pore). However, this process will cause further disorder of the curve graphene sheets, which may build two-dimensional sodium metallic clusters close enough to share one conducting electron, and in turn causes sodium cluster to exhibit quasi-metallicity obtained by D-PDF, which shows very similar characteristics to metallic sodium. Conclusion In summary, we have designed and synthesized two series of HC materials not only with perfect spherical morphology but also with tailored microstructures by simple Stöber method and carbonization route. The morphologies, microstructures, and electrochemical performance of HC materials were investigated. Besides the correlation between the electrochemical performance and structures, we investigated the sodium storage mechanism of HC anodes using GITT, ex situ XRD, and ex situ PDF. We propose a new “three-phase” structure model for HC materials, and prove its reliability to explain the evolution of HC structures under different carbonization conditions, thereby illuminating the relationship between their properties and structures, and revealing the sodium storage mechanism that SRC owes to the adsorption of Na ions in active sites on the surface of HC spheres, PRC can be divided into two parts according to the variation law of Na+ ion DCs. The first part is related to IC of Na ions among parallel graphene sheets of


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graphite-like microcrystallites, and the end part of PRC is assigned to FPC, forming Na clusters in profuse confined spaces in HC anodes at the voltage range of 0.03 – 0.01 V. We hope the findings give new insight for carbonaceous materials of SIBs and provide new perspectives to further develop the controversial sodium storage mechanism. Experimental Section Synthesis of HC-Ts The synthesis methods of RF reference the deformation of the Stöber Method, the molar ratio of resorcinol and formaldehyde (37wt%) is 1:2, using ammonia aqueous (25wt%) as catalyst for the polymerization process, while the mixture was added in water and was continuously stirred for 20 h at room temperature. Then the mixture was transferred into a Teflon-lined autoclave at 100 °C for 30 h to fabricate polymer spheres. After that using a centrifuge to separate and washing the resin spheres by water and ethanol for three times. Then the precursor was dried by vacuum freeze-drying for 48 hours. The obtained brown powder was carbonized at processing temperatures, from 800 to 1600 °C, signed as HC-T (T is temperature). Synthesis of HC-Ss The HC-Ss were obtained similar to that of the HC-Ts, only the mixture was added in different solvents. The obtained brown powder was carbonized at 1300 °C for 2 h under argon atmosphere. Other synthesis steps were the same as HC-Ts. In control experiments, using water, methanol, ethanol, isopropanol, butanol as solvent respectively, a series of different products, respectively are marked as the HC-H2O, HC-CH3OH, HC-C2H5OH, HC-C3H7OH, HC-C4H9OH,


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signed as HC-S (S is solvent). Meanwhile, we prepared RF by controllable the Stöber time, PH value and hydrothermal temperatures to get RF resin with different crosslinking degrees. Characterization X-ray diffraction (XRD) patterns were measured by using Panalytical X’pert PRO MRD (Holland) with Cu Kα radiation. Raman spectra were obtained by a LabRAM HR800 (Horiba JobinYvon) using an Ar ion laser with a wavelength of 532 nm. JSM 7600F field-emission scanning electron microscope (FE-SEM, JEOL, Japan) was used to check the morphology and size of product. Transmission electron microscopy (TEM) observations were carried out with JEM-2100 electron microscope. Fourier-transform infrared spectroscopy (FT-IR) of KBr powder pressed pellets was recorded on a Bruker Vector 22 spectrometer. Thermogravimetry-Differential Scanning Calorimetry (TG-DSC) measurement was performed on a Netzsch STA 449F3 analyzer from 30 to 1000 °C at a rate of 10 °C min-1 in Ar. Reciprocal space data for PDF analysis were obtained using a PANalytical Empyrean outfitted with Ag-Kα radiation and GaliPIX3D detector. Electrochemical measurements The composition of hard carbon electrode is prepared by a mixture of hard carbon active mass 80%, 10 wt% Super-P 10% (as a conductive agent), 10 wt% polymer binder (NaCMC: SBR = 1: 1) into a thin film. Electrode slurry is coated on copper foil and dried at 120 °C under vacuum for 12 h. CR2032 coin cells are assembled in a glove box full of argon gas, 1 mol L-1 NaClO4 in ethylene carbonate/propylene carbonate (1:1 by volume) as electrolyte, using porous glass fiber from Whatman as separator and a sodium sheet with an 8 mm diameter as counter electrode. Cyclic voltammetry (CV) was measured on a Princeton electrochemical workstation between 0


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V to 2 V at a scan rate of 0.5 mV s-1. The charge/discharge tests were performed on a battery testing system (Land Electronics, China) at room temperature between 0 V and 2 V versus Na+/Na. PDF tests Batteries for PDF analysis were cycled at a rate of C/15 (based on 1C= 300 mAh g-1) to the test point on the curve. The batteries were disassembled inside an argon atmosphere glovebox, and the electrodes were washed with propylene carbonate and immediately dried in vacuum. The active materials were scraped from the electrodes and sealed into kapton tubes. Structural models were refined against PDF data using PDFGui. Difference PDFs were calculated by subtracting the PDF for the former electrode from the PDFs of the electrodes at various discharge stages. An instrument schematic shows in Figure S12.


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Figure 1 (a) The Na+ ion diffusion coefficients calculated from the GITT curve of HC electrode. (b) Schematic illustration of the mechanisms for Na-ion storage in hard carbon: absorption capacity (AC), interlayer intercalation capacity (IC), and filling-pore capacity (FPC).


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Figure 2 (a) X-ray diffraction patterns of HC-Ts carbonized at different temperatures. (b) Raman spectrum of HC-Ts. (c) X-ray diffraction patterns of HC-Ss assembled in different solvents. (d) Raman spectrum of HC-Ss.


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Figure 3 FE-TEM images of HC-Ts carbonized at different temperatures of (a – e) 800, 1100, 1300, 1500, and 1600 °C, respectively. Graphitic structures gradually grew with increasing temperatures. SEM image of (f) RF-H2O and (g) HC-1300 (HC-H2O).


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Figure 4 FE-TEM images of (a) HC-H2O. (b) HC-CH3OH. (c) HC-C2H5OH. (d) HC-C3H7OH. (e) HC-C4H9OH. SEM image of as-prepared RF assembled in different solvents of (f) RFCH3OH. (g) RF-C2H5OH. (h) RF-C3H7OH. (i) RF-C4H9OH.


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Table 1. Structural parameters and electrochemical properties of HC-Ss assembled in different solvents, and HC-Ts carbonized at different temperatures.

RC (mA ICE L ab a sample IG/IDa Sample I – d G/ID hg (%) (nm) 1 c ) HC-H2O 0.323 6.2 310 84 HC-800 0.304 5.86 HC-CH3OH 0.330 6.34 288 78 HC-1100 0.307 5.90 HC-C2H5OH 0.339 6.51 278 76.1 HC-1300 0.323 6.20 HC-C3H7OH 0.350 6.75 266 72.5 HC-1500 0.341 7.88 HC-C4H9OH 0.359 6.91 140 40 HC-1600 0.538 10.34 a ID and IG are the integrated intensities of the D and G band. b L (nm) = (2.4 ×10−10 )λ L ab (nm)

a

RC (mAh g–1)c

nm

ICE (%)d

215 288 310 275 264

71 78 84 80.1 80.7 c I 4 The ( G) ID

d

reversible capacity (RC). The initial coulombic efficiency (ICE)


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Figure 5 (a) The second discharge/charge profiles of HC-Ts carbonized at different temperatures with the current density of 20 mA g-1. (b) Cycling performance at current density of 20 mA g1

. (c) The second discharge/charge profiles of HC-Ss at the current density of 20 mA g-1. (d)

Cycling performance at current density of 20 mA g-1.


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Figure 6 (a – e) GITT curves and the inserts of Na+ ion diffusion coefficients calculated from GITT curves of HC-Ts electrodes (carbonized at different temperatures) as a function of cell voltage during discharge processes. The blue blocks represent the intercalation region and the red blocks represent the plating region. (f) Discharge capacity from adsorption, intercalation, and filling pores.


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Figure 7 (a), (b), (c), (d), (e) GITT curves and the Na+ ion diffusion coefficients calculated from the GITT curves of different HC electrodes (assembled in different solvents) as a function of cell voltage during discharge processes (inset). The blue blocks represent the intercalation region and the red blocks represent the filling pores region. (f) Discharge capacity from adsorption, intercalation, and filling pores.


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Figure 8 Schematic illustration of the mechanisms for Na-ion storage in hard carbon: (a) Schematic illustration of HC-Ts at different temperatures. (b) Na-ion storage mechanism in HC.


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Figure 9 (a) The ex situ XRD mapping of the HC-H2O electrode during the first dischargecharge at 20 mA g-1. (b) Electrodes discharge at various stages. (c) PDF data for electrodes discharging at the different stages. (d) D-PDFs of HC-H2O anodes at various states of discharge.


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ASSOCIATED CONTENT

Supporting Information Available: SEM, FTIR, TG and electrochemistry performance for materials mentioned in the article, including RF resin assembled with different processing time, with different PH and hydrothermal temperatures; Details for fitted Raman curves; Details for GITT and PDF tests; AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] ORCID Jiantao Han: 0000-0002-9509-3785

Author Contributions The manuscript was written through contributions of all authors. Y.J. and J.H designed the experiment, Y.J., C.F., M.O., prepared the series of HC-Ts and HC-Ss. Y.J., S.S., Y.L. designed electrochemical and physical measurements. Y.J., S.S., J.P., P.W., Z.D.,Y.Q. performed and analyzed electrochemical measurements and physical characterization. Y.J., Y.L and Y.X. carried out ex situ XRD. Y.J., M.O. and X.S. designed the PDF tests and analyzed the structure characterization. Y.J. , J.H. and Y H wrote this paper. ACKNOWLEDGMENT


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This work was supported by the National Key R&D Program of China (Grant Nos. 2016YFB010030X and 2016YFB0700600), National Natural Science Foundation of China (Grant Nos. 51772117 and 51732005). The authors also thank the Analytical and Testing Centre of HUST and the State Key Laboratory of Materials Processing and Die & Mould Technology of HUST for XRD, SEM, TEM, Raman, TGA and other measurements.

ABBREVIATIONS HC, Hard carbon; SIBs, sodium-ion batteries; RF, resorcinol formaldehyde; PDF, pair distribution function; D-PDF, Difference PDFs LIBs, Lithium-ion batteries; SRC, sloping region capacity; PRC, plateau region capacity; DC, diffusion coefficient; AC, absorption capacity; IC, interlayer intercalation capacity; FPC, filling-pose capacity;


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REFERENCES (1) Bruce, D.; Kamath, H.; Tarascon J. M. Electrical Energy Storage for the Grid: A Battery of Choices. Science 2011, 334, 928-934. (2) Larcher, D.; Tarascon, J. M. Towards greener and more sustainable batteries for electrical energy storage. Nat. chem. 2015, 7, 19-29. (3) Choi, N. S.; Chen, Z.; Freunberger, S. A.; Ji, X.; Sun, Y. K.; Amine, K.; Yushin, G.; Nazar, L. F.; Choi, J.; Bruce, P. G. Challenges facing lithium batteries and electrical double-layer capacitors. Angew. Chem. 2012, 51, 9994-10024. (4) Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research development on sodium-ion batteries. Chem. Rev. 2014, 114, 11636-11682. (5) Luo, W.; Shen, F.; Bommier, C.; Zhu, H.; Ji, X.; Hu, L. Na-Ion Battery Anodes: Materials and Electrochemistry. Acc. Chem. Res. 2016, 49, 231-240. (6) Slater, M. D.; Kim, D.; Lee, E.; Johnson, C. S. Sodium-Ion Batteries. Adv. Funct. Mater. 2013, 23, 947-958. (7) Kundu, D.; Talaie, E.; Duffort, V.; Nazar, L. F. The emerging chemistry of sodium ion batteries for electrochemical energy storage. Angew. Chem. 2015, 54, 3431-3448. (8) Wang, Y.; Xiao, R.; Hu, Y. S.; Avdeev, M.; Chen, L. P2-Na0.6[Cr0.6Ti0.4]O2 cation-disordered electrode for high-rate symmetric rechargeable sodium-ion batteries. Nat. Commun. 2015, 6, 6954. (9) Wang, Y.; Liu, J.; Lee, B.; Qiao, R.; Yang, Z.; Xu, S.; Yu, X.; Gu, L.; Hu, Y. S.; Yang, W.; Kang, K.; Li, H.; Yang, X. Q.; Chen, L.; Huang, X. Ti-substituted tunnel-type Na(0.44)MnO(2) oxide as a negative electrode for aqueous sodium-ion batteries. Nat. Commun. 2015, 6, 6401.


30

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


(10) Liu, Q.; Meng, X.; Wei, Z.; Wang, D.; Gao, Y.; Wei, Y.; Du, F.; Chen, G. Core/Double-Shell Structured Na3V2(PO4)2F3@C Nanocomposite as the High Power and Long Lifespan Cathode for Sodium-Ion Batteries. ACS Appl. Mater. interfaces 2016, 8, 31709-31715. (11) Zhao, Y.; Wei, Z.; Pang, Q.; Wei, Y.; Cai, Y.; Fu, Q.; Du, F.; Sarapulova, A.; Ehrenberg, H.; Liu, B.; Chen, G. NASICON-Type Mg0.5Ti2(PO4)3 Negative Electrode Material Exhibits Different Electrochemical Energy Storage Mechanisms in Na-Ion and Li-Ion Batteries. ACS Appl. Mater. interfaces 2017, 9, 4709-4718. (12) Wang, L.; Song, J.; Qiao, R.; Wray, L. A.; Hossain, M. A.; Chuang, Y. D.; Yang, W.; Lu, Y.; Evans, D.; Lee, J. J.; Vail, S.; Zhao, X.; Nishijima, M.; Kakimoto, S.; Goodenough, J. B. Rhombohedral prussian white as cathode for rechargeable sodium-ion batteries. J. Am. Chem. Soc. 2015, 137, 2548-2554. (13) You, Y.; Wu, X.-L.; Yin, Y.-X.; Guo, Y.-G. High-quality Prussian blue crystals as superior cathode materials for room-temperature sodium-ion batteries. Energy & Environ. Sci. 2014, 7, 1643-1647. (14) Wu, X.; Deng, W.; Qian, J.; Cao, Y.; Ai, X.; Yang, H. Single-crystal FeFe(CN)6 nanoparticles: a high capacity and high rate cathode for Na-ion batteries. J. Mater. Chem. A 2013, 1, 10130. (15) Nobuhara, K.; Nakayama, H.; Nose, M.; Nakanishi, S.; Iba, H. First-principles study of alkali metal-graphite intercalation compounds. J. Power Sources 2013, 243, 585-587. (16) Wen, Y.; He, K.; Zhu, Y.; Han, F.; Xu, Y.; Matsuda, I.; Ishii, Y.; Cumings, J.; Wang, C. Expanded graphite as superior anode for sodium-ion batteries. Nat. Commun. 2014, 5, 4033.


31


Page 32 of 37

(17) Liu, Y.; Merinov, B. V.; Goddard, W. A. Origin of low sodium capacity in graphite and generally weak substrate binding of Na and Mg among alkali and alkaline earth metals. Proc. Natl. Acad. Sci. U. S. A 2016, 113, 3735-3739. (18) Hong, K.-l.; Qie, L.; Zeng, R.; Yi, Z.-q.; Zhang, W.; Wang, D.; Yin, W.; Wu, C.; Fan, Q.-j.; Zhang, W.-x.; Huang, Y.-h. Biomass derived hard carbon used as a high performance anode material for sodium ion batteries. J. Mater. Chem. A 2014, 2, 12733. (19) Lv, W.; Wen, F.; Xiang, J.; Zhao, J.; Li, L.; Wang, L.; Liu, Z.; Tian, Y. Peanut shell derived hard carbon as ultralong cycling anodes for lithium and sodium batteries. Electrochimi. Acta 2015, 176, 533-541. (20) Xu, D.; Chen, C.; Xie, J.; Zhang, B.; Miao, L.; Cai, J.; Huang, Y.; Zhang, L. A Hierarchical N/S-Codoped Carbon Anode Fabricated Facilely from Cellulose/Polyaniline Microspheres for High-Performance Sodium-Ion Batteries. Adv. Energy Mater. 2016, 6, 1501929. (21) Gu, J.; Z. D.; Zhang, C.; Ma, J.; Li, B. and Yang, S. Liquid-Phase Exfoliated Metallic Antimony Nanosheets toward High Volumetric Sodium Storage. Adv. Energy Mater. 2017, 7, 1700447. (22) Huang, Y.; Li. X., Luo, J.; Wang, K.; Zhang, Q.; Qiu, Y.; Sun, S.; Liu, S.; Han, J. and Huang, Y. Enhancing Sodium-Ion Storage Behaviors in TiNb2O7 by Mechanical Ball Milling. ACS Appl. Mater. interfaces 2017, 9, 8696-8703. (23) Fang, C.; Huang, Y.; Yuan, L.; Liu, Y.; Chen, W.; Huang, Y.; Chen, K.; Han J.; Liu, Q. and Huang, Y. A Metal-Organic Compound as Cathode Material with Superhigh Capacity Achieved by Reversible Cationic and Anionic Redox Chemistry for High-Energy Sodium-Ion Batteries. Angew. Chem. 2017, 56, 6793-6797.


32

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


(24) Li, W.; Zhou, M.; Li, H.; Wang, K.; Cheng, S.; Jiang, K. A high performance sulfur-doped disordered carbon anode for sodium ion batteries. Energy Environ. Sci. 2015, 8, 2916-2021. (25) Luo, W.; Bommier, C.; Jian, Z.; Li, X.; Carter, R.; Vail, S.; Lu, Y.; Lee, J. J.; Ji, X. Lowsurface-area hard carbon anode for Na-ion batteries via graphene oxide as a dehydration agent. ACS Appl. Mater. interfaces 2015, 7, 2626. (26) Dahn, J. R.; Stevens, D. A. High Capacity Anode Materials for Rechargeable Sodium-Ion Batteries. J. Electrochem. Soc. 2000, 147, 1271-1273. (27) Komaba, S.; W. M.; Ishikawa, T.; Yabuuchi, N.; Ozeki, T.; Nakayama, T.; Ogata, A.; Gotoh, K. and Fujiwara K. Electrochemical Na Insertion and Solid Electrolyte Interphase for HardCarbon Electrodes and Application to Na-Ion Batteries. Adv. Funct. Mater. 2011, 21, 3859-3867. (28) Xiao, L.; Cao, Y.; Henderson, W. A.; Sushko, M. L.; Shao, Y.; Xiao, J.; Wang, W.; Engelhard, M. H.; Nie, Z.; Liu, J. Hard carbon nanoparticles as high-capacity, high-stability anodic materials for Na-ion batteries. Nano Energy 2016, 19, 279-288. (29) Li, Y.; Hu, Y.; Titirici, M. M.; Chen, L.; Huang, X. Hard Carbon Microtubes Made from Renewable Cotton as High-Performance Anode Material for Sodium-Ion Batteries. Adv. Energy Mater. 2016, 6, 1600659. (30) Qiu, S.; Xiao, L.; Sushko, M. L.; Han, K. S.; Shao, Y.; Yan, M.; Liang, X.; Mai, L.; Feng, J.; Cao, Y.; Ai, X.; Yang, H.; Liu, J. Manipulating Adsorption-Insertion Mechanisms in Nanostructured Carbon Materials for High-Efficiency Sodium Ion Storage. Adv. Energy Mater. 2017, 7, 1700403. (31) Jian, Z.; Bommier, C.; Luo, L.; Li, Z.; Wang, W.; Wang, C.; Greaney, P. A.; Ji, X.; Insights on the Mechanism of Na-Ion Storage in Soft Carbon Anode. Chem. Mater. 2017, 29, 2314.


33


Page 34 of 37

(32) Li, Z.; Bommier, C.; Chong, Z.; Jian, Z.; Surta, T. W.; Wang, X.; Xing, Z.; Neuefeind, J. C.; Stickle, W. F.; Dolgos, M.; Greaney, P. A.; Ji, X. Mechanism of Na-Ion Storage in Hard Carbon Anodes Revealed by Heteroatom Doping. Adv. Energy Mater. 2017, 7, 1602894. (33) Bommier, C.; Surta, T. W.; Dolgos, M.; Ji, X. New Mechanistic Insights on Na-Ion Storage in Nongraphitizable Carbon. Nano Lett. 2015, 15, 5888-5892. (34) Wang, K.; Jin, Y.; Sun, S.; Huang, Y.; Peng, J.; Luo, J.; Zhang, Q.; Qiu, Y.; Fang, C.; Han, J. Low-Cost and High-Performance Hard Carbon Anode Materials for Sodium-Ion Batteries. ACS Omega 2017, 2, 1687−1695. (35) Sadezky, A.; Muckenhuber, H.; Grothe, H.; Niessner, R.; Pöschl, U. Raman microspectroscopy of soot and related carbonaceous materials: Spectral analysis and structural information. Carbon 2005, 43, 1731-1742. (36) Cançado, L. G.; Takai, K.; Enoki, T.; Endo, M.; Kim, Y. A.; Mizusaki, H.; Jorio, A.; Coelho, L. N.; Magalhães-Paniago, R.; Pimenta, M. A. General equation for the determination of the crystallite size La of nanographite by Raman spectroscopy. Appl. Phys. Lett. 2006, 88, 163106. (37) Uvarov, V.; Popov, I. Metrological characterization of X-ray diffraction methods at different acquisition geometries for determination of crystallite size in nano-scale materials. Mater. Charact. 2013, 85, 111-123. (38) Chen, Y.; Chen, Z.; Xiao, S.; Liu, H. A novel thermal degradation mechanism of phenol– formaldehyde type resins. Thermochim. Acta 2008, 476, 39-43. (39) Tsai, P.; Chung, S.; Lin, S.; Yamada, A. Ab initio study of sodium intercalation into disordered carbon. J. Mater. Chem. A 2015, 3, 9763-9768.


34

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


(40) Joshua, M. S.; Phoebe, K. A.; Oliver P.; Philip A. C. and Grey, C. P. Mechanistic insights into sodium storage in hard carbon anodes using local structure probes. Chem. Mater. 2013, 00, 1-3. (41) Forse, A. C.; Merlet, C.; Allan, P. K.; Humphreys, E. K.; Griffin, J. M.; Aslan, M.; Zeiger, M.; Presser, V.; Gogotsi, Y.; Grey, C. P. New Insights into the Structure of Nanoporous Carbons from NMR, Raman, and Pair Distribution Function Analysis. Chem. Mater. 2015, 27, 68486857. (42) He, H.; Huang, C.; Luo, C.; Liu, J.; Chao, Z. Dynamic study of Li intercalation into graphite by in situ high energy synchrotron XRD. Electrochimi. Acta 2013, 92, 148-152. (43) Hantel, M. M.; Nesper, R.; Wokaun, A.; Kötz, R. In-situ XRD and dilatometry investigation of the formation of pillared graphene via electrochemical activation of partially reduced graphite oxide. Electrochimi. Acta 2014, 134, 459-470.


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Table of Contents

High-performance Hard Carbon Anode: Tunable Local Structures and Sodium Storage Mechanism In this work, we proposed a new “three-phase” structural model for HC materials and found a 3region Na-ion storage behaviors, chemi-/physisorption, intercalation between carbon layers, and pore-filling. These results enable new understanding and insights into sodium storage mechanism in HC materials and improve carbon-based materials for anodes of SIBs.


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High-Performance Hard Carbon Anode: Tunable Local Structures and

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