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Probing enhanced lithium-ion transport kinetics in 2D holey nanoarchitectured electrodes
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Nano Futures 2 (2018) 035008
https://doi.org/10.1088/2399-1984/aada90
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14 June 2018
Probing enhanced lithium-ion transport kinetics in 2D holey nanoarchitectured electrodes
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7 August 2018 ACCEPTED FOR PUBLICATION
16 August 2018 PUBLISHED
Xiao Zhang1,7, Andrea M Bruck2,7, Yue Zhu1, Lele Peng1, Jing Li3, Eric Stach4, Yimei Zhu5, Kenneth J Takeuchi2,3, Esther S Takeuchi2,3,6, Amy C Marschilok2,3,8 and Guihua Yu1 ,8 1
31 August 2018 2 3
4 5 6
7 8
Materials Science and Engineering Program and Department of Mechanical Engineering, The University of Texas at Austin, Austin, TX 78712, United States of America Department of Chemistry, SUNY-Stony Brook University, Stony Brook, NY 11794, United States of America Department of Materials Science and Chemical Engineering, SUNY-Stony Brook University, Stony Brook, NY 11720, United States of America Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia, PA 19104, United States of America Condensed Matter Physics and Materials Science, Brookhaven National Laboratory, Upton, NY 11973, United States of America Energy Sciences Directorate, Brookhaven National Laboratory, Interdisciplinary Sciences Building, Building 734, Upton, NY 11973, United States of America These authors contributed equally. Authors to whom any correspondence should be addressed.
E-mail:
[email protected] and
[email protected] Keywords: energy storage, nanostructuring, 2D holey nanosheets, transport kinetics, Li-ion battery Supplementary material for this article is available online
Abstract Nanostructuring has been proved effective towards improving many energy storage and conversion devices, and is feasible for a wide range of materials. In particular, secondary nanoarchitectured materials exhibit collective advantages compared with nano-sized primary building blocks. Despite the manifold efforts in designed nanoarchitectures and synthetic routes, the underlying ion diffusion kinetics and phase transformation behaviors within nanoarchitectures still remain less explored. Herein, we probed enhanced lithium-ion transport behaviors using 2D holey zinc ferrite (ZFO) nanosheets as a model material, to demonstrate how self-assembled 2D holey nanoarchitectured electrodes can feature efficient ion diffusion channels, robust yet continuous electron transfer framework, and enlarged surface area, contributing to the superior performance over the ZFO nanoparticles without secondary structures. By revealing kinetic parameters through combined spectroscopic measurements and electrochemical techniques, our study manifests increased lithiumion diffusion coefficients, higher capacitive charge storage contribution and reduced charge transfer impedance in holey nanosheets compared to randomly aggregated nanoparticles. Our results promote deeper understanding of significantly enhanced electrochemical energy storage properties of these 2D holey nanoarchitectured electrodes resulted from more uniform and complete phase transformation and better active material utilization.
1. Introduction Energy storage systems are of vital importance for the development of a cleaner and more energy efficient society. Due to increasing demands for long life time, safety, high capacity, and high energy density, there is still room for improvements based on current material systems. It has been well-acknowledged that nanostructuring is a powerful tool to improve electrochemical performance of many energy storage and conversion devices, and is applicable for a wide spectrum of materials [1, 2]. Despite great achievements in synthetic approaches towards electrode fabrication with carefully designed nanoarchitectures, a deeper understanding at the atomic and mesoscale levels of the underlying physical/chemical processes governing macroscopic performance is still lacking [3]. Even though efforts on theoretical modeling have led to insights from a structural perspective into © 2018 IOP Publishing Ltd
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the reaction mechanism, complexity and non-uniformity in real battery devices prohibit revealing corresponding kinetic ionic and electron transport. The link between atomic scale and mesoscale is still weak due to a lacking of thorough understanding of reaction kinetics [4]. Therefore, investigations on the charge carrier transport behavior by electrochemical measurements combined with advanced in situ characterization are critically needed and can provide useful information for multivariate battery systems. Herein mixed transition-metal oxides (MTMOs) are explored as model materials for probing ion transport study. As promising anode materials for lithium-ion batteries (LIBs), MTMOs usually possess high theoretical capacities due to unique conversion reactions. With two types of transition metal cations in the crystal structure, MTMOs offer higher electrical conductivity and richer redox reactions than simply transition-metal oxides (TMOs) [5, 6]. Moreover, advantages including low cost, low toxicity and facile synthesis render MTMOs great potential for commercialization [7]. However, their practical applications are often hindered by poor transport kinetics of charge carriers and significant volume change during electrochemical processes [8–10]. Towards the above-mentioned issues, nanostructuring can be applied in order to obtain substantial improvements [11]. Especially, porous nanoarchitectured MTMO electrodes could further extend the surface area, featuring more open structure buffering volume change, adjustable frameworks and efficient ion transports [12–14]. For example, two-dimensional holey nanoarchitectures have been recently shown to exhibit greatly enhanced alkaliion storage properties compared to non-holey counterparts [15–17]. Therefore, it is of fundamental interests to study lithium-ion transport kinetics in 2D holey MTMO electrodes. In this study, we aim to develop deeper understanding of lithium-ion transport behaviors by looking into differences between ZFO nanoparticles without secondary nanostructures (ZFO NP) and self-assembled holey nanosheet (ZFO HNS) architectures with appealing structural features. By assembling nano-sized building blocks into secondary nanoarchitectures, nanocrystal properties are brought to larger length scales. In addition, collective advantages could be induced due to artificially designed superstructures [18–20]. The enhanced rate capability exhibited in ZFO HNS suggests possible morphology-induced superior ion transport kinetics. Further experimental evidence is provided by applying various electrochemical techniques to extract diffusion coefficients under different depths of discharge, capacitive/diffusion-controlled capacity contributions, and charge transfer impedances. In addition, in situ TEM and XRD are employed to reveal a holistic picture of the nature and rates of phase transformations accompanying the conversion reactions. The integrated electrochemical characterization and spectroscopic measurements promote better understanding of architecture-performance nexus that has not been well studied in prior research.
2. Methods 2.1. Synthesis of ZFO NP and ZFO HNS ZnFe2O4 nanoparticles were synthesized via a previously reported method [21]. Using deionized water as solvent, stoichiometric solutions of Zn2+ and Fe3+ precursors were added simultaneously to a basic aqueous solution. The resulting precipitate, after washing and drying, was subjected to a hydrothermal treatment at 220 °C for 12 h. The final product was collected, washed and vacuum-dried. Holey ZFO nanosheets were synthesized using a confined self-assembled method, where graphene oxide (GO) and Pluronic copolymer served as sacrificial template and surfactant, respectively [22]. GO was prepared from purified natural graphite using a modified Hummers method [23]. In a typical process, 20 mg Pluronic copolymer with a molecular weight of 4400 was first mixed with 20 mg ZFO nanoparticles in 5 ml ethylene glycol (EG). Then 3 mg GO dispersed in 5 ml EG was added dropwise to the ZFO-Pluronic mixture. After magnetic stirring for 24 h, the precipitate was centrifuged and washed with ethanol for 4 times and vacuum-dried under 50 °C. Finally, the as-obtained precursor was annealed in a Lindberg/Blue box furnace at 400 °C for 2 h at a ramp rate of 0.5 °C min−1 in air. 2.2. Cell assembly All electrochemical measurements were performed in a half-cell configuration. CR2032 coin cell assembly was carried out in an argon-filled glovebox with lithium metal as anode and reference electrode. The cathode was composed of a mixture of ZFO active materials (either nanoparticles or holey nanosheets), super P carbon and polyvinylidene difluoride binder with a weight ratio of 7:2:1. The mixture was made into a slurry and was spread onto a copper foil. After drying at 120 °C under vacuum for 24 h, the copper foil was cut into circular electrodes with an average mass loading at ∼1.0 mg cm−2. Celgard 2320 was used as separator between cathode and anode. The electrolyte was 1.0 M LiPF6 dissolved in a mixture of ethylene carbonate and dimethylene carbonate with 1:1 ratio. All electrochemical tests were conducted at room temperature (25 °C) with a voltage range 0.1–3.1 V. 2
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2.3. Morphology TEM characterization, including imaging and electron diffraction, was performed on a JEOL JEM-2100F analytical TEM equipped with a field-emission electron gun operated at 200 kV. 2.4. In situ XRD Specially designed cells were assembled based on a previously published procedure [24]. The cell stack was constructed with a Li metal anode, a polyethylene separator, and the corresponding ZFO electrode on a Cu foil current collector in a dryroom environment (0.99). Therefore, the diffusion coefficient calculation could be expressed as the following equation:
D=
2 2 4 ⎛ IVm ⎞ ⎛ DEs ⎞ ⎛ L2 ⎞ ⎜ ⎟⎜ ⎟ , ⎜t ⎟ , D⎠ pt ⎝ ZA FS ⎠ ⎝ DEt ⎠ ⎝
(1)
where D is the diffusion coefficient of lithium-ion (cm2 s−1); τ is the duration of the current pulse (s); I is the applied current (A); Vm is the molar volume of the electrode (cm3 mol−1); ZA is the charge number of the ion (1 for lithium-ion); F is the Faraday constant; S is the surface area of the electrode (cm2); L is the thickness of the electrode (cm); ΔEs is the steady-state voltage change between two adjacent current pulses and ΔEt is the voltage change excluding iR drop in each current pulse. To clearly illustrate diffusion coefficient trends, the log diffusion coefficients (D) calculated according to GITT data were plotted versus ees and the two nanostructures are compared in figure 2(c). In general, ZFO NP and ZFO HNS show a similar trend, where D gradually increases before ∼2.0 ee, followed by a sharp decrease to ∼4.0 ee, and then increases again when the batteries are further discharged to ∼7.0 ee. As concluded by previous studies, lithiation of spinel ZFO crystallites is based on an intercalation mechanism within the first 2 ee [7]. More specifically, lithium-ions will first insert into octahedral 16c sites accompanied by the displacement of Zn2+ from tetrahedral (8a) to an octahedral (16c) site. The migration of the Zn2+ ions activates adjacent 16c sites for additional Li insertion resulting in a stable LixZnFe2O4 configuration [36]. Therefore, a general D increase from 0 to 2.0 ee can be attributed to the accelerated lithiation reaction facilitated by vacant sites in the spinel structure followed by the Zn2+ migration to an octahedral site, yielding fast lithium ion storage kinetics before ∼2.0 ee. Subsequent D decrease can also be explained through previous DFT calculations, which conclude that once all 5
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Figure 3. (a) Summarized intensity profiles of in situ SAED patterns recorded in ZFO HNS during lithiation. Morphological evolution observed by (b)–(d) in situ TEM and (e)–(g) corresponding FFT patterns.
octahedral sites are occupied, further insertion of lithium-ion is accommodated by less active tetrahedral 8a/ 48f/8b sites. Once the tetrahedral Li sites are fully occupied, the structure becomes unstable with no additional sites for further lithiation available the structure begins to convert to the corresponding metal species and lithium oxide, which could limit lithium diffusion between 2 and 6 ee. Finally (between 6 and 8 ee), conversion reactions progress where the metallic species become the majority of the phase fraction, since the metal species and corresponding oxides are finely divided, the increase in D observed around 7.0 ee could be explained by the expansion of the material creating more conversion sites for lithium-ions. Despite the similar trend in D change, ZFO NP electrodes exhibit an earlier and less drastic D decrease beginning at ∼1.7 ee, while a rapid D decrease in ZFO HNS electrode happens later (∼2.4 ee). Based on the lithiation mechanism mentioned before, D decrease is caused by further lithiation into energetically unfavorable vacancy sites. Therefore, this delayed D decrease in ZFO HNS can be interpreted as more lithium-ions stored before conversion begins, indicating higher utilization of ZFO material in the HNS. In contrast, less lithium storage is achieved in ZFO NP, most likely due to the aggregation nature of the nanoparticles causing heterogeneous lithiation between particles [37]. This is also supported by the facts that, reduction peaks of ZFO NP in CV scans are at higher potentials, reaffirming a hasty phase transformation (figure S7). Furthermore, the slope of decreasing D can be considered to represent the rate of phase transformation from spinel structure to rock salt structure in the overall electrode. Hence the faster D decrease in ZFO HNS electrode also stands as clear evidence for more rapid and facile phase transformation in this case. Besides the initial lithiation, it can be easily seen that ZFO HNS exhibits larger D than ZFO NP at all the other ees, indicating overall faster lithium-ion storage in ZFO HNS, which is consistent with the its superior rate performance. The enhanced lithium-ion diffusion in ZFO HNS is mainly attributed to shortened diffusion lengths due to highly interconnected nanoparticle network and larger contact area with electrolyte and conductive additives. 3.3. Phase evolution probed by in situ SAED and XRD To study the phase transformation directly, the integrated in situ SAED shows a gradual phase transformation in ZFO HNS upon lithiation, where spinel structured ZFO are converted into Zn, Fe and Li2O as final products (figure 3(a)). This result matches well with previous study, which can be expressed as the equation: ZnFe2 O4 + 8Li+ + 8e- Zn + 2Fe + 4Li2O.
(2)
Localized phase evolution at the atomic scale is visualized with in situ TEM (figures 3(b)–(d)) and corresponding phase components are determined by fast Fourier transfer (FFT) (figures 3(e)–(g)). It appears that the conversion reaction takes place from surface to bulk, forming Li2O layer at the surface. Such Li2O surface layer might limit further lithiation into the bulk by hindering charge transfer process. Given the high dispersity and interconnectivity in ZFO HNS, we believe that the Li2O surface layer would have less influence to charge transfer process, therefore, leading to enhanced lithium transport kinetics. In situ XRD provides the opportunity to analyze phase evolution in ZFO materials with kinetic differences between nanoparticle aggregations and uniform holey nanosheets morphologies. It also provided sensitivity toward the ZnFe2O4 to LixZnFe2O4 (LZFO) transition. Figures 4(a) and S8 present the in situ XRD patterns of ZFO HNS and ZFO NP electrodes. Before discharge, both ZFO NP and ZFO HNS exhibit broad Bragg reflections at ca. 34.7° and 56.1°, which correspond to (311) and (511) lattice planes respectively in the spinel 6
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Figure 4. (a) In situ XRD data of ZFO HNS showing detailed phase transformation and component evolution during the first discharge process. (b) Comparison of phase evolutions in ZFO NP versus ZFO HNS.
Fd3m ZnFe2O4 crystal structure. The LixZnFe2O4 phase is identified with Bragg reflections at ca. 41.9° and 61.1°, representing the (400) and (044) lattice planes. All other unchanging peaks are attributed to Cu foil current collector, Ni current collector tabs, Li metal, separator, and plastic housing directly from the in situ experimental set up, as shown by the red, green and blue traces in figure 4(a). Note, the (044) peak presented in ZFO material shows a 0.7° shift (indicating a lattice expansion from 1.50 to 1.52 Å along the (044) plane) up to 2.0 ee. This expansion is associated with the change from a spinel to a rock salt-like structure to accommodate more lithiumions. The right side of figure 4(a) shows the corresponding discharge profile with respect to ees. Comparing galvanostatic discharge curves of the two, it is found that ZFO HNS reaches ∼11.9 ee, and ZFO NP reaches only ∼11.1 ee when fully discharged under the same condition (figure S9). Peak intensity evolution curves for ZFO NP and ZFO HNS are summarized and compared in figure 4(b). In the initial lithiation process (
