New Mechanistic Insights on Na-Ion Storage in Non-Graphitizable Carbon Clement Bommier, Todd Wesley Surta, Michelle Dolgos, and Xiulei Ji* Department of Chemistry, Oregon State University, Corvallis, Oregon 97331-4003, United States Abstract: Non-graphitizable carbon, also known as hard carbon, is considered one of the most promising anodes for the emerging Na-ion batteries. The current mechanistic understanding of Na-ion storage in hard carbon is based on the ‘card-house’ model first raised in the early 2000’s. This model describes that Na-ion insertion occurs first through intercalation between graphene sheets in turbostratic nanodomains, followed by Na filling of the pores in the carbon structure. We testified this model by tuning the sizes of turbostratic nanodomains but found the opposite trend. Moreover, the results revealed a correlation between the structural defects and Na-ion storage. Based on our experimental data, we propose an alternative three-phase model for sodiation of hard carbon that consists of Na-ion storage at defect sites, through intercalation and lastly by pore-filling. KEYWORDS Na-ion batteries • intercalation • hard carbon • reaction mechanism • defect sites Abstract Graphic
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In today’s world, there is an ever-increasing need to move beyond Li-Ion Batteries (LIBs) for more sustainable and affordable electrochemical energy storage (EES) solutions.1 While LIBs represent the state-of-the-art technology, the relative scarcity and uneven global distribution of lithium reserves greatly hinders their use in large-scale applications such as grid-level EES.2 To date, Na-ion batteries (NIBs) represent one of the most promising alternatives due to the similar chemical and electrochemical properties between Na and Li and the vast abundance of Na resources.3-6 However, some differences between the two types of atoms have presented considerable challenges in the development of NIBs, especially in the case of anode materials.7, 8 LIBs utilize a graphite anode, as Li ions can reversibly form the LiC6 stage-one graphite intercalation compound (GIC). However, the Na-ion’s larger radii of 102 pm vs. 76 pm of Liion, along with its relatively high ionization potential only allows for the reversible formation of the NaC64 GIC.9-11 In pursuing functional NIB anodes, much attention has been devoted to non-graphitizable carbon, often referred to as hard carbon. Unlike highly ordered graphite, which forms ABAB stacked graphene sheets with an interlayer spacing of 0.335 nm, hard carbon is composed of disordered turbostratic nanodomains (TNs) and empty space (referred to as pores hereafter) between these domains. This leads to three distinct chemical environments for the storage of Naions: edge/defect sites on pore surfaces, e.g., carbenes, vacancies and dangling bonds on the edges of TNs, the interlayer space enclosed inside the TNs, and lastly, the empty pores.12 This specific structure allows for stoichiometry of NaC7.4 or 300 mAh g-1, which is nearly 10 fold that of graphite.
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In the literature, most studies on hard carbon and non-graphitic carbons13-18 as NIB anodes rationalize their electrochemical performances in accordance with the card-house model reported in the early 2000’s.19-22 Under the this model, the Na-ion storage is attributed to two distinct mechanisms: 1) storage of ions between the graphene sheets inside the TNs and 2) Naion/atom filling of the ‘pores’ between the domains. This model further assigns the Na-ion storage between the graphene sheets to the sloping voltage region of a potentiogram while ascribing the storage in the pores to the plateau region, as illustrated in Figure 1. In this study, we tuned the atomic structure of a model hard carbon in order to confirm the electrochemical structure-property relationships given in the card-house model.
We
hypothesized that annealing sucrose-derived hard carbon at progressively increasing temperatures would yield the desired ‘tunable’ hard carbon structures by controlling the domain dimensions of the all-important TNs along the axial axis, Lc, and along the ab planes, La. Sucrose-derived carbons were compared to a commercial glassy carbon (Aldrich) that is essentially the ‘hardest’ available hard carbon, as it is typically obtained through pyrolysis at temperatures of over 2000℃.23 In order to examine the Lc and La parameters of the TNs in different samples, we conducted X-ray diffraction (XRD), Raman spectroscopy and total scattering via neutron diffraction measurements. Surface area and porosity measurements, alongside TEM images, are also included in the supplementary information (SI Table 1 & SI Figure 1-5). In the XRD patterns, we observed the (002) peaks, indicative of the d-spacing between the graphene sheets in the TNs, to be contracting from ~0.375 nm for pyrolyzed sucrose annealed at 1100℃ (S-1100), to ~0.371 nm for S-1400, ~0.370 nm for S-1600 and ~0.352 nm for glassy carbon (Figure 2a).
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Average Lc and La values were estimated through the PDXL software by use of the Scherrer equation on (002) and (100) peaks, respectively.24 The results reveal that the S-1100 is likely to have the smallest TN dimensions with Lc and La values of 1.15 and 2.54 nm, respectively, while these values increase to 1.17 and 3.18 nm for S-1400, 1.19 and 3.50 nm for S-1600 and 1.25 and 6.13 nm for the glassy carbon (SI Table 2). While the X-ray trend is difficult to unambiguously distinguish between the different sucrose carbons, Raman spectroscopy was used as a local probe for more reliable measurements of amorphous materials. The increase in La is also observed through Raman spectroscopy by using the equation below (Figure 2b).25
𝐿! 𝑛𝑚 = 2.4 ⋅ 10!!" ⋅ 𝜆!!"
𝐼! 𝐼!
where λ4nm is from the wavelength of the laser in nanometers, ID is the integrated D band at 1350 cm-1 attributed to sp2 carbon atoms with defects, and IG is the integrated G band at 1580 cm-1 which arises from planar sp2 configured carbon atoms and are typical in pristine graphene. Results of the Raman spectra show a similar trend as the XRD results, with La values of ~9.2 nm, 10.5 nm, 13.3 nm and 19.4 nm for S-1100, S-1400, S-1600 and glassy carbon, respectively (SI Table 2). The scale discrepancy between the XRD and Raman La values has been noted before, as Raman measurements tend to overestimate the La values.26
However, when performing a
linear regression of the Raman La values vs. the XRD La values, we find an R2 of 0.95, thus confirming the increasing trend in La in the carbon materials (SI Figure 6).
Additional
information from the Raman data includes an ID/IG ratio, which can be used to quantify the concentration of defects along the graphene sheets.27, 28 These ratios indicate that the S-1100
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material has the greatest concentration of defects, while the glassy carbon contains the smallest concentration. While these traditional measurements are insightful, they fall short of revealing local atomic structures. Thus, to probe the short and mid-range order of the carbon materials, we performed total neutron scattering measurements and the associated Pair Distribution Function (PDF) analysis. The peaks up to 5 Å represent the short-range correlations of the C-C bonds of one hexagon unit within a sheet of graphene (Figure 2c & SI Figure 7). Regardless of the annealing temperatures, all the peaks are well aligned, and are consistent with the graphene lattice.29 There is an absence of a C-C correlation at 3.35-3.45 Å, which is found in highly ordered, graphitic carbon. This lack ordering between the graphene sheets supports the existence of turbostratic disorder. The size of the graphene nanodomains can be estimated in the mid/long range where well-resolved peaks vanish at certain r values. The S-1100 has notable peak features until ~15-17 Å, the S-1400 exhibits features diminishing around ~18-21 Å, while the glassy carbon has features that persist past 30 Å (Figure 2c, SI figure 8-10). These data are consistent with XRD and Raman results, as the glassy carbon exhibits the largest domains, while the S-1100 has the smallest. In addition to the size of the graphene sheets in TNs, neutron PDF results also provide information on the sp2 defect concentration present in the TNs. Importantly, we note that the intensity of the PDF peaks differs from sample to sample. In PDF patterns, the area underneath a peak is proportional to the coordination number of that pair correlation.30,
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At 1.42 Å the
intensity of S-1400 and glassy carbon is virtually identical and only S-1100 has a slightly lower intensity. However, as r increases, the difference in peak intensity becomes readily apparent (Figure 2c). In particular, the first three peaks in the S-1100 pattern have lower intensities, 5
meaning that it possesses a significant number of non-hexagonal carbon rings. These defects likely result in a bending of the graphene sheets, which diminishes the long range ordering and therefore the size of the domains.32
The PDF data show S-1100 has the greatest defect
concentration, the glassy carbon has the least while S-1400 is somewhere in the middle. The trend regarding the defect concentration corroborates the Raman spectrum results. Using several techniques, we have demonstrated that TN sizes increase upon annealing and are the largest for glassy carbon. The larger domain size increases the total volume of the TNs in hard carbon; therefore, based on the card-house model, the capacity in the sloping part of the potentiogram curve should increase. The sloping region is defined as the capacity obtained above 0.115V vs. Na+/Na. Our data shows that the capacity in this region becomes smaller in the samples with the larger TNs, which is the opposite of what is predicted by the card-house model (Figure 3a). In fact, there appears to be an inverse relationship between the capacity obtained from the slope and the sizes of TNs. Instead of being dependent on the size of TNs, the capacity in the sloping region seems to be directly correlated to the defect concentration present. When plotting the slope capacity versus defect concentration in the carbon samples, as expressed by the integrated ID/IG ratio, we observe a linear relationship with an R2 value of 0.90 (Figure 3b & SI Figure 11). This leads to our hypothesis that the defected carbon sites in hard carbon rather than the TNs are responsible for the capacity in the sloping region. The obtained PDF data also supports this hypothesis as it shows that the S-1100 material has the greatest defect concentration and the largest slope capacity while the glassy carbon has the lowest defect concentration and the smallest slope capacity. This hypothesis is also supported by recent computational studies, where binding
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energy between Na-ion and carbon is calculated to be the highest at defect sites and vacancy sites, which corresponds to the high sodiation potentials in the sloping region.33-36 We further investigated the sodiation mechanism through the evaluation of kinetic properties using galvanostatic intermittent titration (GITT) measurements (SI Figure 12).37,
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The Na-ion diffusivity in hard carbon calculated as a function of potential (Figure 3c) shows that diffusion associated with the sloping potentials is much faster than that with the plateau. This suggests that initial sodiation happens on easily accessible sites in the carbon structure. It is reasonable to say that the surface sites of TNs are more accessible than the interlayer space in the TNs. As these sites are progressively sodiated, Na-ions should then diffuse inside the TNs. However, in order to do so, the Na-ions have to overcome a repulsive charge gradient from the previously bound Na-ions on defect sites in order to diffuse inside the TNs. This explains the steep drop in diffusivity in the plateau region of the potentiogram. These kinetic results not only support our hypothesis that the sloping capacity is due to defect sites on the TN surfaces but also give further insights into the sodiation mechanism of the plateau region of the potentiogram curve.
This has the subject of ambiguity, as recently
published experimental results have suggested that the plateau region is due to Na insertion into TNs as opposed to the pore filling mechanism suggested by the card-house model.39-41 Similarly, we also observed reversible expansion and contraction of d-spacing due to sodiation and desodiation at the plateau region, i.e., from 0.2 V to 0.01 V, while conducting ex-situ XRD. This suggests that intercalation occurs at lower voltages (Figure 4a & 4b). However, these newly obtained diffusion results, combined with the rest of our experimental data suggest that the plateau region can be a function of both intercalation into TNs 7
as well as the pore-filling suggested by the card house model. When observing the values at low voltages, we can observe that diffusion values reach a minimum at 0.05V—the same voltage where dQ/dV values reach a maximum. From 0.05V to the cutoff voltage, the diffusion values gradually increases while dQ/dV values decrease (Figure 3d). This suggests that the sodiation mechanism is changing in the penultimate steps before reaching the cutoff potential. At such low voltages, storage sites can be characterized by a weak binding energy and facile diffusion. We considered the hypothesis that the storage at low voltages is due to the Na-atom adsorption on the sp2 configured pore surfaces. This is in good agreement with the original one made by the card-house model, and is also supported by the computational results; which show that the binding energy between planar graphene sheets and Na is much smaller than binding on defect sites or in a graphene bilayer.[17] It is critical to determine whether such adsorption is of same potential as the plating. When running a half-cell at voltages below 0.0 V, we found that the onset of Na-metal plating actually occurs at -0.02 V before reaching a steady value of -0.015 V (Figure 4c). To summarize, we have revealed some discrepancies with the card-house model on the Na-ion storage in hard carbon. We have demonstrated that the storage mechanism in the sloping part of the potentiogram curve can be better explained through storage at defect sites, as opposed to intercalation between graphene sheets. Furthermore, our additional diffusion, ex-situ XRD and Na-plating experiments lead us to hypothesize that the storage mechanism in the low voltage plateau region is due to the intercalation between graphene sheets and minor phenomenon Naion adsorption on pore surfaces (Figure 4d).
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Figure 1. Visual representation of the card-house model on Na-ion storage in hard carbon. The two distinct phases: intercalation inside TNs and pore filling are seen.
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a
b
c
Figure 2. a,b) XRD patterns and Raman spectra of the different carbon materials. c) PDF results for total neutron scattering, with the inset showing the short-range order.
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a
c
b
d
Figure 3. a) Sodiation potentiograms for different carbons. Complete electrochemical data in SI Figure 13-19 & SI Table 3-5. b) Plot of the sloping capacity vs. the ID/IG ratio from Raman spectra. c) GITT profile and diffusivity as a function of states of charge (inset). d) dQ/dV plot from 0.12 V to 0.01 V with corresponding diffusivity values.
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a
b
c
d
Figure 4. a) Ex-situ XRD profiles at various stage of sodiation (S) and desodiation (D) with the (002) peak indicated. b) Corresponding d-spacing plots. c) Potentiogram of sodiation until Nametal plating is induced. d) Potentiogram and schematic of proposed Na-ion three part storage mechanism.
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ASSOCIATED CONTENT Supporting Information. Surface Area & Porosity, HRTEM, Additional PDF data, Complete electrochemical characterization and experimental methods are included. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *E-mail:
[email protected] ACKNOWLEDGMENT X. Ji acknowledges the financial supports from Advanced Research Projects Agency-Energy (ARPA-E), DOE of the United States, Award number: DE-AR0000297TDD.
M. Dolgos
gratefully acknowledges the financial support from Oregon State University.
Research at
ORNL’s Spallation Neutron Source was sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences of the U.S. DOE. X. Ji thanks Dr. Yuliang Cao from Wuhan University for discussion.
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