Designing Hierarchical Assembly of Carbon-Coated TiO2

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Designing Hierarchical Assembly of Carbon-Coated TiO2 Nanocrystals and Unraveling the Role of TiO2/Carbon Interface in Lithium-Ion Storage in TiO2 Je Uk Ha, Jeongmin Lee, Muhammad A Abbas, Moo Dong Lee, Junghyun Lee, and Jin Ho Bang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b21705 • Publication Date (Web): 04 Mar 2019 Downloaded from http://pubs.acs.org on March 4, 2019

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ACS Applied Materials & Interfaces

Designing Hierarchical Assembly of Carbon-Coated TiO2 Nanocrystals and Unraveling the Role of TiO2/Carbon Interface in Lithium-Ion Storage in TiO2 Je Uk Ha,†,# Jeongmin Lee,†,# Muhammad A. Abbas,‡ Moo Dong Lee,† Junghyun Lee,† and Jin Ho Bang†, ‡,$,* Department of Bionano Technology, Nanosensor Research Institute, and Department of Chemical and Molecular Engineering, Hanyang University, 55 Hanyangdaehak-ro, Sangnok-gu, Ansan, Gyeonggi-do 15588, Republic of Korea AUTHOR INFORMATION Corresponding Author: Jin Ho Bang *

Email: [email protected]

†

Department of Bionano Technology, Hanyang University, 55 Hanyangdaehak-ro, Sangnok-gu,

Ansan, Gyeonggi-do 15588, Republic of Korea ‡

Nanosensor Research Institute, Hanyang University, 55 Hanyangdaehak-ro, Sangnok-gu, Ansan,

Gyeonggi-do 15588, Republic of Korea

ACS Paragon Plus Environment

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Department of Chemical and Molecular Engineering, Hanyang University, 55 Hanyangdaehak-

ro, Sangnok-gu, Ansan, Gyeonggi-do 15588, Republic of Korea #

These authors contributed equally to this work.

KEYWORDS: titanium dioxide, hierarchical assembly, interface, carbon layer, oxygen vacancy

ABSTRACT. Despite the many benefits of hierarchical nanostructures of oxide-based electrode materials for lithium-ion batteries (LIBs), it remains a challenging task to fully exploit the advantages of such materials partly because of their intrinsically poor electrical conductivities. The resulting limited electron supply to primary particles inside secondary microparticles gives rise to significant variation in lithium-ion (Li+) storage capability within the nanostructured particles. To address this, facile annealing, where in situ generated carbon-coated primary particles were assembled into porous micro-agglomerates, is demonstrated to prepare nanostructured titanium dioxide (TiO2). A systematic study on the effect of the carbon coating reveals that it is exclusively governed by the characteristics of the TiO2/carbon interface rather than by the nature of the carbon coating. Depending on their number, oxygen vacancies (VO) created by carbothermal reduction on the TiO2 surface are detrimental to Li+ diffusion in the TiO2 lattice, and structural distortion at the interface profoundly influences the Li+ (de-)intercalation mechanism. This new insight serves as a stepping stone toward understanding an important, yet often overlooked effect of the oxide/carbon interface on Li+ storage kinetics, thereby demanding more investigation to establish a new design principle for carbon-coated oxide electrode materials.


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1. Introduction Owing to its high chemical stability, structural robustness, abundance, and high working voltage (~1.75 V vs. Li+/Li), anatase TiO2 has extensively been investigated as an anode material for LIBs since the first report of the electrochemical insertion of Li+ into TiO2 by Ozhuku et al.1 However, these merits are largely compromised by its intrinsically poor electrical conductivity (~10 −9 Ω−1 m−1)2 and lithium diffusivity (~10−15 cm2 s−1),3 which are major hurdles for the widespread application of TiO2 to LIBs. Numerous studies to overcome this limitation have been carried out. Among them, strategies based on preparing nanostructured TiO2 stand out, and the nanostructuring is critical for the application of TiO2 to LIBs from two perspectives. First, it facilities Li+ transport by shortening diffusion length (τeq=L2/2DLi, τeq: equilibrium time for Li+ diffusion, L: particle size, and DLi: diffusion coefficient of Li+).4 Second, it alters the thermodynamics and kinetics of Li+ insertion into TiO2 (TiO2 + xLi+ + xe- → LixTiO2, x: mole fraction of Li+ in TiO2) much more favorably as Mulder et al. revealed earlier.5,6 Since the maximum Li+ insertion (typically x=0.5 for bulk TiO2) is size-dependent, decreasing the size of TiO2 particles to less than 7 nm can render them fully able to accommodate Li+.5 Although a wide variety of TiO2 nanostructures have proven their prowess, hierarchically structured TiO2––where nanoparticles are assembled into porous micro-agglomerates––has been considered as the ideal nanostructure. Hierarchically assembled TiO2 has often been prepared by burning sacrificial porogens off of organic/inorganic hybrid precursors while converting them into mesoporous TiO2.7-9 Another popular route is utilizing selfassembly by Ostwald ripening under solvothermal conditions.10-13 This nanoarchitecture can mitigate the disadvantages of TiO2 while not significantly compromising electrode packing density. However, the hierarchical structure promotes uneven electron mobility in each primary particle depending on the location in the agglomerate.14 Nanoparticles inside the micro-agglomerates are


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likely to show higher resistance than those outside them because they are not directly in contact with conductive carbon additive; hence, electron transport within the micro-agglomerates primarily proceeds by hopping (Figure 1A).15 Given the importance of well-balanced transport between Li+ and electrons (e-) for effective Li+ migration into the TiO2 lattice (i.e., the concurrent hopping of Li+/e- pairs),16 this limited electron supply to the interior particles hampers Li+ intercalation and, hence, inevitably leads to high overpotentials for Li+ transport.4,17 To resolve this issue, we synthesized hierarchically assembled carbon-coated TiO2 nanocrystals (HAC TiO2), where all the primary particles were coated with a thin carbon layer. Our work demonstrates that the carbon network formed throughout the hierarchical agglomerate can significantly alleviate the spatially unbalanced electron supply, hence enhancing the storage and rate capability of TiO2. Another important contribution of this work is that it provides new insights into the influence of the TiO2/carbon interface on the Li+ storage of TiO2, which has long been overlooked. Our extensive literature survey led us to affirm the common belief in the carbon coating effect whereby a thin carbon layer overcoated on TiO2 particles enhances the electrical conductivity of TiO2, thus improving its LIB performance.18-25 However, we determined that this benefit intimately depended on the nature of the TiO2/carbon interface. In-depth surface analysis revealed that the thermal reduction of TiO2 by the in situ generated carbon removed some oxygen atoms, thereby introducing surface oxygen vacancies (VO) (Figure 1B). The degree of VO at the interface profoundly impacted the LIB performance of TiO2, thereby kindling our interest to understand the nature of the TiO2/carbon interface and its relationship with Li+ storage kinetics. Our systematic investigation on this undisclosed aspect of the interface is discussed in this report.


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2. Experimental Section 2.1 Synthesis of HAC TiO2. TiO2 nanocrystals were synthesized by a nonhydrolytic sol–gel reaction as follows: 10.0 mL of titanium (IV) isopropoxide and 50 mL of benzyl alcohol were poured into a 100 mL C-276 liner. After stirring the solution at 500 rpm for 30 min, the liner was slid into a bench-type autoclave and tightly sealed. The autoclave was heated to 250 °C, held there for 24 h, and then naturally cooled. A white precipitate was collected by centrifugation, washed with 60 mL of tetrahydrofuran, and dried in a vacuum oven at 60 °C for 24 h. To produce the carbon-coated TiO2 agglomerate, the as-obtained TiO2 powder (1.0 g), denoted as TiO2-as made, was placed in a ceramic crucible and annealed in N2 for 1 h. Each sample was heated to 500, 600, or 700 °C at 10 °C min−1. The obtained HAC TiO2 will hereafter be referred to as TiO2-500, TiO2600, and TiO2-700, respectively. For comparison, the as-obtained TiO2 powder was heated under the same conditions except in O2. 2.2 Characterization. X-ray diffraction (XRD) patterns of the HAC TiO2 were obtained using an X-ray diffractometer (Bruker D2 PHASER) with Cu Kα irradiation and operating at 10 mA and 30 kV. Nitrogen adsorption/desorption isotherms were measured at 77 K, using a surface area analyzer (BELSORP-mini II, BEL Japan), and the specific surface area and pore size distributions were determined by the Brunauer–Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) methods, respectively. Raman spectroscopy (Renishaw RM-1000) was performed to examine the carbon layer on the HAC TiO2. Surface analysis was performed by X-ray photoelectron spectroscopy (XPS, PHI Versa Probe system), and the oxygen vacancies of TiO2 surface were examined using an electron paramagnetic resonance (EPR) spectrometer (Bruker EMXplus9.5/2.7). Thermogravimetric analysis (TGA) were performed with a thermal analyzer (SDT-Q600, TA Instrument) ramped at 5 °C min−1 in air. Transmission electron microscopy (TEM, JEOL-


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2100F) and scanning electron microscopy (SEM, Hitachi S-4800) were performed to acquire information about morphology and local crystal structure. 2.3 Electrode Fabrication and Evaluation. For electrochemical characterization, CR-2032-type coin cells were assembled in an argon-filled glove box and were evaluated using various electrochemical techniques. Each TiO2 sample was thoroughly mixed with carbon black (Super P) and polyvinylidene difluoride (PVDF) and was dissolved in N-methyl-2-pyrrolidone (NMP) in a weight ratio of 70:15:15, using a mixer (AR-100, THINKY Corp.). The prepared slurry was spread uniformly thick onto a current collector (Cu foil), using a micrometer adjustable film applicator, and the electrode film was dried at 120 °C for 12 h in a vacuum oven. The active material loading was set to ~1.0 mg cm−1. For the coin cell assembly, lithium metal foil, 1 M LiPF6 dissolved in ethylene carbonate/dimethyl carbonate (1:1 v:v), and Celgard 2320 membrane were employed for the counter electrode, electrolyte, and separator, respectively, and the working electrode was made of TiO2 film. Galvanostatic charge–discharge measurements and galvanostatic intermittent titration technique (GITT) analysis were carried out using battery cyclers (WBCS3000, WonATech Co., Ltd.). In situ Electrochemical impedance spectroscopy (EIS) was measured during charge/discharge cycles at various potentials using ZIVE BP2A electrochemical workstation WonATech Co., Ltd.).

3. Results and Discussion In our approach to preparing the HAC TiO2, crystalline anatase TiO2 nanoparticles were first produced by a nonhydrolytic sol–gel reaction,26 where Ti–O–Ti bonds were formed through the C–C coupling of titanium tetrabutoxide and benzyl alcohol, and the resulting nanoparticles were then annealed at 500, 600, or 700 °C under an inert atmosphere. Annealing the TiO2


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nanocrystals formed porous agglomerates while converting surface species (i.e., residues of benzyl alcohol and other decomposed organic species) into thin carbon layers. This conversion of an organic stabilizer into carbon layer was quite successful in other reports as well.27,28 The annealing temperature was limited to up to 700 °C because the rutile phase (which is less efficient for Li + storage than the anatase phase) began to form at elevated temperatures (Figure S1). Figure 2 shows the XRD patterns of the TiO2-as made, TiO2-500, TiO2-600, and TiO2-700. All the XRD peaks matched the typical diffraction peaks of anatase TiO2 (JCPDS 21-1272), and none of the peaks were attributable to any impurities in the specimens. The TiO2 cell parameters and grain sizes were determined by Rietveld refinement and are summarized in Table 1. Although the grain size of the TiO2 annealed at 500 °C increased from 9.9 to 13 nm, further increase in the grain size was substantially suppressed even when the annealing temperature was raised to 700 °C. Unlike the TiO2 annealed in N2, however, that annealed in air showed a dramatic increase in grain size (14.9, 29.2, and 63.2 nm for TiO2-500, TiO2-600, and TiO2-700, respectively), highlighting another benefit of the in situ generated carbon layer over the TiO2 particles: preventing grain growth even during high-temperature annealing (Figures S2-3). Since there is always a compromise between crystallinity and surface area during high-temperature annealing,29,30 this simple, scalable process can be of great benefit for the preparation of high-performance electrode materials. The lattice parameters of TiO2-as made (a=3.7868 Å, c=9.4974 Å), on the other hand, slightly deviated from those of standard anatase TiO2 (a=3.783 Å, c=9.507 Å; JCPDS 21-1272), particularly in the [001] direction. However, the shorter lattice along the [001] direction of TiO2-as made lengthened upon annealing, and the lattice parameters of all the HAC TiO2 became closely matched to those of standard anatase TiO2.


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The microtextural properties of the HAC TiO2 were characterized by N2 physisorption isotherms (Figure 3), and the obtained microtextural parameters are summarized in Table 1. The overall trend in the isotherms was notably different from that observed for the TiO2 samples annealed in air. Although the mesoporosity of the agglomerated TiO2 annealed in air abruptly decreased with temperature (Figure S4), the HAC TiO2 inherited the original porosity of TiO2-as made by and large. It is noteworthy that the thermal treatment offered sufficient robustness to the porous structure, which played a crucial role in stability during battery operation that will be discussed later. The hysteresis loop became shifted toward a high-pressure region after the annealing, implying that a mesoporous structure with larger pores had been formed at the expense of smaller mesopores during the annealing. The specific surface area of TiO2-as made was 70.92 m2 g−1, and those of TiO2-500, TiO2-600, and TiO2-700 were as high as 93.78, 88.69, and 83.69 m2 g−1, respectively. The pore size distribution (in the insets of Figure 3) showed that although annealing had slightly narrowed the multimodal mesopore distribution, the similar total pore volumes (~0.29 cm3 g−1) of the HAC TiO2 suggested that there was no significant difference in the porous structures of the three TiO2 agglomerates. This trend was strikingly different from that found in the TiO2 annealed in air; the surface area of TiO2 heated at 700 °C was merely 3.73 m2 g−1, and the mesoporosity completely disappeared when TiO2 was annealed at 700 °C. The nature of the carbon layers in the HAC TiO2 was examined by Raman spectroscopy. Figure 4A-C display the Raman spectra of TiO2-500, TiO2-600, and TiO2-700, where two characteristic Raman peaks designated as D and G bands appeared at ~1343 and 1589 cm−1, respectively. This spectral feature was different from that observed for TiO2-as made, where the distinctive peaks were not well resolved (Figure S5). These bands were carefully deconvoluted to gain more insight into the carbon nature, the results of which are presented in Figure 4. It is


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generally accepted that D1 and G1 peaks appearing at ~1344 and ~1593 cm−1 are analogous to those of the D and G bands of crystalline graphite, and D2 and G2 peaks at ~1412 and ~1525 cm−1 are ascribed to those of the D and G bands of amorphous carbon.31,32 Note that additional peaks labeled D3 (~1275 cm−1) and S (~1131 cm−1) were included in the D band fitting as previously reported,33 which presumably stemmed from the very disordered nature of the carbon layer formed by the carbonization of organic residues at low temperatures.21 The D′ peak at ~1617 cm−1 has been known to be associated with disordered carbon edges.34-36 Shimodaira et al. proposed the G2/G1 ratio as a useful gauge to identify the relative disorderness in the sp2 network of activated carbons (a smaller ratio is associated with less disorderness),31 based on which we concluded that a more-disordered carbon layer had been formed on TiO2-700 than on TiO2-500 (Figure 4D and Table 1). The amount of carbon in each TiO2 sample was estimated using TGA, according to which the carbon contents were 4.5, 3.8, and 3.5 wt.% for TiO2-500, TiO2-600, and TiO2-700, respectively. The surface characteristics of the HAC TiO2 were investigated using XPS. Figure 5A exhibits the XPS Ti 2p spectra of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700. The raw spectral data were deconvoluted into two doublet peaks (2p3/2 and 2p1/2), each of which was assigned to Ti4+ at 458.7 and 464.4 eV and to Ti3+ at 457.3 and 462.5 eV, respectively (See Table S1 for detailed fitting results). Interestingly, although there was no Ti3+ species in TiO2-as made, Ti3+ evolved upon annealing and, in particular, a remarkable amount of it was found in TiO2-500. With increasing annealing temperature, the intensity of the peaks attributed to Ti 3+ diminished. Since the formation of Ti3+ accompanies the creation of VO,37-39 this implied that carbonization had introduced defects (VO) to the TiO2 surface as follows: TiO2 (s) + C (s) → TiO2-δ (s) + CO2 (g)

(1)


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In fact, the reducing power of carbon at elevated temperatures (i.e., carbothermal reduction) has long been utilized to extract metals from ores. The interfacial oxygen defects could survive with the aid of the carbon layer wrapping individual TiO2 nanocrystals.40-42 This could be part of the reason for the gradual decrease in the intensity of the peaks attributed to Ti 3+ with temperature because the carbon content decreased with increasing annealing temperature. In addition, the more disordered carbon layer formed in TiO2-700 may ease the diffusion of O2 through the carbon layer, leading to the decrease in VO. The XPS O 1s spectra of the four samples (Figure 5B, Table S2) showed the same pattern as the Ti 2p spectra. The peak at 530.32 eV, corresponding to VO, appeared upon annealing and faded with increasing annealing temperature. The appearance of a peak attributed to O2- at 527.7 eV along with the VO peak was in accordance with this observation because the adsorbed anionic oxygen could exist only when excess electrons provided by VO were available.43 The XPS spectrum of TiO2 annealed in air (Figure S7) did not show any VO signature peaks, confirming the role of the carbon layer as the reducing agent. The VO in the HAC TiO2 was affirmed by EPR spectroscopy. Figure 6 shows the EPR spectra of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700, where although a characteristic peak at g=2.001 (corresponding to VO) was absent in the EPR spectrum of TiO2-as made, it appeared upon annealing, and its intensity diminished with increasing annealing temperature. The morphology and microstructural information about the TiO2 samples were investigated using electron microscopy. The TEM analysis of TiO2-as made (Figure S8A) showed that the TiO2 nanocrystals consisted of a mixture of rhombic and rectangular particles whose average sizes were in the range 15–25 nm. A high-resolution TEM (HRTEM) image combined with a fast Fourier transform (FFT) pattern, where an ordered dotted pattern was observed, revealed the single


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crystalline nature of the TiO2 nanoparticles (Figure S8B). The TEM images along with the SEM images of the HAC TiO2 (Figure 7A-C and Figures S9-10) show that the TiO2 nanocrystals coated with the thin carbon layer were packed into dense micro-agglomerates. To acquire the detailed atomic structures of the HAC TiO2, the local crystal structures of the TiO2 surfaces were analyzed in depth. The HRTEM images in Figure 7D-F were taken along the [1̅1̅1] zone axis (as affirmed by analyzing the FFT patterns), hence allowing the inspection of the (101) plane of each HAC TiO2. The Ti–Ti interlayer distance (d) along the [101] direction (as depicted in the modeled crystal structure) was carefully measured using TEM intensity profiles (Figure 7G-I), where the spacings of five exterior atomic layers and five interior ones were examined for each sample. For TiO2-500 (Figure 7G), although the average interior interlayer distance was 3.36 Å, the average exterior one was 3.50 Å, implying structural distortion on the TiO2-500 surface.44-46 In contrast, the spatial difference in the interlayer distance became much less significant for TiO2-600 (Figure 7H), and there was nearly no variation in the interlayer distance for all the examined areas of TiO 2-700 (Figure 7I), which corroborated the XPS and EPR analysis results in that VO were preferentially created at the interface between the TiO2 surface and the carbon layer. The Li+ storage capability of the HAC TiO2 was assessed by galvanostatic charge/discharge measurements at various C rates in the range 1.0–3.0 V vs. Li+/Li. Figure 8A displays the voltage profiles of the four TiO2 electrodes measured at 1 C (168 mA g−1). The voltage profile of TiO2 has been well established by three characteristic regions (often referred to as Regions A, B, and C).47,48 In Region A (above 1.75 V), Li+ is initially stored in the TiO2 lattice by the formation of a lithiumpoor solid solution (i.e., no structural change occurs during Li+ insertion) denoted as LiδTiO2.49,50 Region B corresponds to the flat potential region at ~1.75 V, where biphasic domains of lithiumpoor tetragonal titanate (LiδTiO2) and lithium-rich orthorhombic titanate (Li0.55TiO2) coexist.51-53


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The sloped region below 1.75 V is designated as Region C, where a diffusion-limited biphasic transition between the orthorhombic Li0.55TiO2 and fully lithiated TiO2 (Li1TiO2) occurs. Although there is a second pseudo-plateau at ~1.5 V, it is often indistinguishable at fast charge/discharge rates (Figure S11) because of the more sluggish diffusion kinetics for inserting Li+ into Li0.55TiO2 whose octahedral sites are more than half-filled.47,54-57 The specific capacities increased with increasing annealing temperature, and the capacity of TiO2-700 at 1 C reached as much as 225 mAhg−1. Figure 8B shows that the capacity of TiO2-700 can be sustained at as high as 188 and 162 mAhg−1 at 5 and 10 C, respectively, (and even 91 mAhg−1 at 50 C), which are comparable to the best sustained capacities reported in the literature (Table S3). On the contrary, with decreasing annealing temperature, the capacity declined more noticeably; for instance, the capacities of TiO2500 at 5 and 10 C were merely 85 and 56 mAhg−1, respectively. The LIB performance of TiO2 annealed in air (uncoated TiO2), on the other hand, was inferior to that of HAC TiO2 (Figure S12), highlighting the merits of the in situ generated thin carbon layer: preserving nanostructures that facilitated Li+ diffusion while providing better electrical conductivity. The electrochemical behavior of TiO2 is dictated by several factors including crystallite size, porous structure, surface area, and morphology. Although there were slight dissimilarities in these structural parameters, our analyses revealed that all three TiO2 agglomerates were nearly identical in their structural and textural characteristics. It is noteworthy that the lattice parameters of TiO2-as made were different from those of standard TiO2, implying structural disordering (i.e., lattice distortion) in TiO2-as made, but not preferentially on the surface. Strictly speaking, TiO2500 seemed to possess more favorable structural characteristics for LIB application than the other annealed TiO2 samples: smaller grain size, larger pore volume and surface area, and larger lattice spacing along the c axis (i.e., the principal Li+ diffusion channel),58,59 albeit slightly different


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among the annealed samples. Despite these benefits of TiO2-500, its performance was inferior to that of TiO2-700, which led us to explore the influence of the TiO2/carbon interface on Li+ storage. Aside from the aforementioned TiO2 characteristics, the thin carbon layer could be an extrinsic factor that affected the LIB performance of TiO2. Since Li+ intercalates into graphitic layers at potentials below 1.0 V vs. Li+/Li,60 the major diffusion pathways for Li+ were likely to be an interconnected pore network created by various defects in the carbon layer. For this reason, the porous structure rather than the thickness and graphitic nature of the carbon layer, as noted previously, would be the most critical factor that determined the effect of the carbon layer.61,62 The TEM images of the carbon layers revealed that the layers were thin, rough, porous, and very similar in appearance, regardless of the annealing temperature (Figure S13). Although we speculated that the observed disparities in carbon content and the degree of graphitization may have impacted Li+ diffusion, the TEM analysis suggests that the impact would unlikely be so great to entirely dictate the LIB performance of TiO2 (as will be evidenced by electrochemical investigation). A noticeable difference in the discharge curves (Figure 8C) was the more elongated plateau region for TiO 2700 than for its other three counterparts. The capacity contributions from Region C were similar for all four electrodes (70–80 mAhg−1), and the slight differences in contribution (TiO2-500 > TiO2-600 > TiO2-700 > TiO2-as made) followed the trend in their specific surface areas. Given that the pseudocapacitive reaction is a dominant charge storage mechanism in Region C, this consistency sounded plausible because interfacial storage tended to be proportional to surface area. Hence, the similar capacity in Region C could be attributed to similar surface area and porous structure. It is noteworthy that if there were a significant difference in Li+/e- transport through the carbon layers, this would never be the case. Therefore, this observation partly supported our speculation on the effect of the carbon layer. In contrast, the capacity contributions from Regions


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A and B to each electrode were distinct. For TiO2-700, it was as high as 150 mAhg−1, but it decreased with decreasing annealing temperature to ~90 mAhg−1 for TiO2-500 and TiO2-as made. In general, the relative contributions of Regions A and B to the overall capacity is intimately related to crystallite size.5,6,54 Given the similar crystallite sizes of all the TiO2 samples, however, the remarkably different capacity suggested that VO at the interface might play a critical role in dictating Li+ (de-)intercalation kinetics. The stability of HAC TiO2 was examined by a long-term cycling test (Figure S14), where all HAC TiO2 retained their original capacity over 500 cycles unlike TiO2-as made that showed gradual degradation. This revealed that the carbon layer could stabilize TiO2 upon repeated charge/discharge cycles. This was partly supported by the SEM images taken after the long-term cycling test (Figure S15). The majority of micrometer-sized particles were preserved in TiO2-500, TiO2-600, and TiO2-700. In the contrary, TiO2-as made were randomly agglomerated because of the lack of rigid interparticle connections and carbon coating as in TiO2-500, TiO2-600, and TiO2-700. The XRD analysis provided additional insight into the capacity degradation (Figure S16). While anatase TiO2 phase was preserved by and large in all the electrodes, there was a small peak at ~45° in the case of TiO2-as made and TiO2-500 that was determined to be bronze-type TiO2 phase (TiO2(B)). We speculated that the formation of this new crystal phase may be associated with the capacity fade. To gain insight into the influence of the TiO2/carbon interface, cyclic voltammetry was carried out, and the resulting cyclic voltammograms (CVs) are presented in Figure 8D-G. A salient feature in the CVs was a pair of new redox peaks (labeled as P2) appearing in the ranges 1.2–1.5 and 2.5–2.8 V in addition to the typical redox peaks corresponding to lithium insertion/extraction into the TiO2 lattice (labeled as P1). The new cathodic and anodic peaks appeared more prominent with increasing scan rate, and the peak currents associated with P2 became even higher than those


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associated with P1 at the high scan rates for TiO2-as made. Another important trend in the CVs was the disappearance of P2 with increasing annealing temperature. According to our extensive literature survey, although P2 has rarely been reported, it seemed to appear when Li+ diffusion was severely restricted.30,54,63,64 The recent report by Passerini et al. disclosed that the advent of this alternative Li+ insertion/extraction pathway was intimately related to structural disorder, which impeded Li+ diffusion in TiO2.63 Given this report, the distorted lattice of TiO2-as made (i.e., the shrunken lattice along the [001] direction) was presumably responsible for the relatively high P2 current. However, the lattice distortion was significantly alleviated upon annealing, as evidenced in the XRD analysis, on the basis of which we attributed P2 in the CVs of the HAC TiO2 to the remarkably contrasting surface characteristics at the TiO2/carbon interface. In general, the Li+ diffusion coefficient (DLiδ) of nonstoichiometric TiO2 (i.e., TiO2-δ) is a function of the ambipolar conductivity of lithium (σLiδ), as expressed by the following relationship:4,65 σLiδ ∝ σLi+ × σe- / (σLi+ + σe-),

(2)

where σLi+ is the ionic conductivity and σe-, the electrical conductivity. The dependencies of σLi+ and σe on the degree of Vo (δ) run in opposite directions;4 hence, there exists an optimal δ, where Li+/e- transport is well balanced for efficient Li+ diffusion. Therefore, more surface Vo have often been reported as deleterious because irreversibly trapped Li+ on defects could significantly hinder Li+ diffusion,4,66,67 which apparently was the case for TiO2-500, whose surface accommodated a considerable number of Vo. The relatively low initial coulombic efficiency of TiO2-500 compared to those of TiO2-600 and TiO2-700 reflected greater irreversible Li+ trapping in TiO2-500 (Figure S14). With decreasing Vo, however, unfavorable Li+ trapping at the TiO2/carbon interface was mitigated, allowing TiO2 to better utilize the benefits of the carbon layer. The more hindered Li+ diffusion in TiO2-500 was also supported by the overshooting peak in the


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potential profile: a peak appearing in the charging profile before the potential plateau begins.68,69 This peak is the signature of the TiO2 memory effect caused by sluggish Li+ diffusion;69 therefore, the overshooting peak in the charging profile for TiO2-700 was undeveloped compared to its counterparts, indicative of less hindered Li+ diffusion. The Li+ diffusion kinetics of each TiO2 electrode was further investigated using in situ EIS, where impedance spectra were measured during the charge/discharge process. The limitations of Li+ insertion/extraction originate from the hindrance of ionic (or electronic) charge transfer at the electrode/electrolyte interface and inside a Li+-hosting matrix (i.e., TiO2). These kinetic barriers were reflected in the Nyquist plots of the TiO2 electrodes measured during charge and discharge cycle at various potentials (Figure 9A,B), where distorted arcs associated with interfacial charge transfer processes appeared in the high frequency region, and the inclined linear responses related to Li+ diffusion inside TiO2 followed in the low frequency region. Several resistance components linked to the kinetics of Li+ insertion/extraction were extracted by fitting the spectra to an equivalent circuit (Figure S17) consisting of uncompensated ohmic resistance (Rs), mostly representing the electrolyte resistance, coupled with several resistance (R)/constant phase element (CPE) components. In the equivalent circuit, the resistance associated with charge transfer were represented by Rct, and the resistance of Li+ diffusion within the TiO2 lattice by RLi. The CPE— expressed by ZCPE = Q(jw)-α (0≤ α ≤1), where j is an imaginary unit, w is the angular frequency, and Q and α are frequency-independent CPE parameters—represented the nonideality associated with electrode roughness, which is common for battery electrodes.38 Figure 9C-F show the resistances derived from the EIS fitting results. Although the interfacial charge transfer resistance (Rct) were almost independent of the applied potentials, the bulk diffusion resistance (RLi) varied to some extent depending on the potentials in particular in TiO2-as made and TiO2-500, which was


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consistent with the EIS analysis results of TiO2 reported previously.70 The Rct values of TiO2-700 measured during the discharging were notably smaller than those of TiO2-600, TiO2-500, and TiO2-as made (Figure 9C), implying that the kinetic limitation associated with Li+ diffusion through the interfacial Li+ intake from the electrolyte was mitigated more effectively in TiO2-700 than in its counterparts. The same Rct trend was observed in the charging (Figure 9E), revealing that the Li+ extraction from TiO2 host at the interface was more favorable in TiO2-700 as well. Since the nature of the carbon layer was primarily responsible for the Li+ diffusion and electron transport (i.e., electrical conductivity) that dictated these surface processes, we speculated that more a kinetically favorable carbon layer had been formed in TiO2-700. Given the greater RLi values and their variation among the electrodes, however, we speculated that the kinetic limitation of TiO2 mainly stemmed from Li+ diffusion within TiO2 particles. The overall trend in RLi was consistent with LIB performance (Figure 9D, F), and this trend in solid-state Li+ diffusion was also verified by the diffusion coefficients measured by the GITT analysis (Figure S18). As GarciaBelmonte et al. noted,70 the difference in RLi most likely resulted from the strain developed at the TiO2/carbon interface given that other structural characteristics were very similar. The interfacial strain was obviously associated with lattice distortion driven by Vo in our case. It is noteworthy that there is a remarkable difference in the potential-dependency of RLi in TiO2-as made and TiO2500 during the charging and discharging cycles. Unlike the RLi trend in the charging (Figure 9D), RLi rose greatly during the charging (Figure 9F), indicating the presence of more sluggish Li + diffusion during the Li+ extraction process. This suggested that the structural disordering had a more influence on Li+ extraction rather than Li+ insertion. To investigate the carbon layer effect on the solid-electrolyte interface (SEI), the in situ EIS spectroscopy was performed again during the first and second cycles and the resulting Nyquist plots were compared (Figure S19). Since the


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resistance corresponding to the SEI layer (RSEI) was not resolved in the Nyquist plots as a separate arc, it was difficult to quantify the resistance. However, the difference in the diameter of the arcs measured in the first and second cycles can reflect the resistance resulting from the SEI formation. The RSEI values were in the order of 100–200 mΩ·g, and their variation among the TiO2 electrodes were very close, hinting that there was no significant influence on the SEI formation by the carbon layer. 4. Conclusion Carbon coating over metal oxides utilized as Li+ storage matrixes has been a common practice to mitigate their poor electrical conductivities. Despite the development of various carbon coating strategies, it has been a great challenge to carbon-coat individual nanoparticles, in particular when metal oxides form hierarchical nanostructures. To remedy this, TiO2 nanocrystals were prepared by a nonhydrolytic sol–gel reaction and then annealed. A thin carbon layer was formed in situ over primary particles while hierarchical agglomerates were produced. This approach enabled the formation of a uniform carbon network throughout the micro-agglomerates, thereby alleviating the spatial unevenness in electron supply often encountered in many hierarchically nanostructured oxides. Aside from devising this synthesis route, our study offered new insights into a critical, yet often overlooked effect of the TiO2/carbon interface on Li+ storage in TiO2. Carbothermal reduction introduced oxygen vacancies on the TiO2 surface, thereby profoundly affecting the Li+ (de-)intercalation mechanism and Li+ diffusion kinetics. Cyclic voltammetry unveiled that structural distortion caused by excess oxygen defects could significantly alter the Li+ (de-)insertion pathway, and in situ EIS analysis combined with galvanostatic charge/discharge measurements revealed sluggish Li+ diffusion due to Li+ trapping at the undesirably defected TiO2 surface. This work not only highlighted the importance of seeking


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a comprehensive understanding of the oxide/carbon interface, but also called urgent attention to devising a delicate approach that could manipulate the characteristics of the interface to fully exploit the benefits of carbon coating.

ASSOCIATED CONTENT Supporting Information. Additional characterizations and electrochemical analysis results. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author: Jin Ho Bang *Email: [email protected] Notes The authors declare no competing financial interests. ACKNOWLEDGMENTS. This work was supported by the Samsung Research Funding Center of Samsung Electronics, under Project Number SRFC-MA1601-03. It was also supported by a grant from the Basic Science Research Program through the National Research Foundation (NRF) of Korea funded by the Ministry of Science and ICT (NRF-2016R1A1A1A05005038, NRF2018M3A7B8061494) and by the Ministry of Education (NRF-2018R1A6A1A03024231).


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Figure 1. (A) Schematic of Li+/e- transport in the hierarchical assembly of TiO2 nanocrystals and the hierarchical assembly of carbon-coated TiO2 nanocrystals (blue arrows represent electron supply from carbon additive and red arrows, electron hopping through TiO2 particles). (B) Schematic of TiO2/carbon interface formed in the carbon-coated TiO2.


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TiO2-as made Raw data Calculated data Diffrence Bragg positions

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1.0

Relative Pressure (P/P0)

Figure 3. N2 physisorption isotherms of (A) TiO2-as made, (B) TiO2-500, (C) TiO2-600, and (D) TiO2-700. Insets show the pore size distributions.


30

Page 31 of 36

800

Raw Data D' G1 G2 D1 D2 D3 S Fitting Line Background

1000

1200

TiO2- 500

1400

Raman Shift Shift Raman

800

1600

1800

1200

800

Raman Shift Shift Raman

1600

1200

TiO2- 600

1400

Raman Shift Raman TiO2- 700

1400

1000

(cm (cm-1))


1000

2000


-1

(D) Relative Amount Amount (% (%)) Relative

(C)

(B) Intensity (a.u.) Intensity (a.u.)

Intensity (a.u.) Intensity (a.u.)

(A)

Intensity (a.u.) Intensity (a.u.)

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


1800

2000

1600

1800

2000

-1

(cm (cm-1)) D' G1 G2 D1 D2 D3 S

30 25 20 15 10 5 0

500

-1

(cm (cm-1))

600

700

Temperature ((oC) ) Temperature ℃

Figure 4. Raman spectra of (A) TiO2-500, (B) TiO2-600, and (C) TiO2-700. (D) Relative amount of each deconvoluted peak (G1: E2g mode associated with the in-plane stretching vibration from graphite-like basal planes, G2: E2g mode associated with the in-plane stretching vibration from the basal planes of disordered carbon, D1: A1g mode associated with the in-plane breathing vibration from the basal planes of disordered carbon, D2 and D3: A1g mode associated with the in-plane breathing vibration from the basal planes of disordered carbon, D′: signal from the disordered edges of carbon, S: signal associated with various sources of disorderness.


31


Ti 2p TiO2-pristine

Ti4+ 2p1/2 464.40 eV

Ti4+ 2p3/2 458.70 eV

2p1/2 464.46 eV

2p1/2 462.56 eV

Ti3+ 2p3/2 457.36 eV

5.7eV

Ti4+ 2p1/2 464.45 eV

Ti4+ 2p3/2 458.75 eV

5.2eV Ti3+ 2p1/2 462.55 eV

Ti3+ 2p3/2 457.35 eV

5.7eV

TiO2-700 2p1/2 464.41 eV

-OH 531.32 eV

464

462

460

Ti-O 529.52 eV

Ov 530.32 eV

O2527.77 eV

TiO2-600 -OH 531.31 eV

5.2eV

-OH 531.33 eV

Ti3+ 2p1/2 462.51 eV

466

Ti-O 529.63 eV

Ti-O 529.51 eV

Ov 530.32 eV

O2527.71 eV

TiO2-700

Ti4+ 2p3/2 458.72 eV

Ti4+

TiO2-pristine -OH 531.53 eV

Ti4+ 2p3/2 458.76 eV

Ti3+

TiO2-600

O 1s

TiO2-500

5.2eV

Ti4+

468

(B)

5.7eV

TiO2-500

Intensity (a.u.)

5.7eV

Intensity (a.u.)

(A)

Ti-O 529.56 eV

Ov 530.36 eV

O2527.90 eV

Ti3+ 2p3/2 457.31 eV

458

456

454

452

536

534

Binding Energy (eV)

532

530

528

526

524

Binding Energy (eV)

Figure 5. XPS spectra of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700: (A) Ti 2p spectra and (B) O 1s spectra.

TiO2-700

Intensity (a.u.)

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

Page 32 of 36

TiO2-600

TiO2-500 TiO2- as made

1.90

1.95

2.00

2.05

2.10

g value

Figure 6. EPR spectra of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700 measured at 110 K.


32

Page 33 of 36 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


Figure 7. TEM images of (A) TiO2-500, (B) TiO2-600, and (C) TiO2-700 (Insets are SEM images). HRTEM images of the (101) lattice plane of (D) TiO2-500, (E) TiO2-600, and (F) TiO2-700 taken along the [1̅1̅1] zone axis, FFT patterns, and anatase TiO2 crystal structure viewed along the [1̅1̅1] direction showing the Ti–Ti distance (d). Lattice distance along the yellow arrow ([101] direction) in the HRTEM images: (G) TiO2-500, (H) TiO2-600, and (I) TiO2-700.


33


(C)

(A)

3.0

TiO2-700 TiO2-600 TiO2-500 TiO2-as made

2.5

2.0

1.5

+

1.0 0

40

80

120

160

200

240

-1

Specific Capacity (mAhg )

(B)

TiO2-700 TiO2-600 TiO2-500 TiO2-as made

-1

Specific Capacity (mAhg )

300 0.5 C

250

1C 2C

200

0.5 C

3C 5C 10 C

150

30 C

100

50 C

50 0

0

10

20

30

40

50

Region C

Region C

0.8 0

Region C

P1

P2

P1

-2.4

Region C

0

40

80

(E)

4 3 2 1 0

2.0

2.5

3.0 1.5

(G) 10 -1

-1.5 -3.0 -4.5 2.5

1.5

2.0

2.5

3.0

+

0

2.0

P2

P2 1.0

8

1.5

TiO2-500

P1

Voltage (V vs. Li /Li) TiO2-600

0.1 mV 0.5 mV 1 mV 2 mV 3 mV 5 mV 10 mV

1.0

240

P1

-3

3.0

Current Density (Ag )

-1

4.5

200

-2

+

6.0

160

-1

Voltage (V vs. Li /Li)

(F)

120

0.1 mV 0.5 mV 1 mV 2 mV 3 mV 5 mV 10 mV

-4 1.5

A+B: 150.40 mAhg -1 C: 74.77 mAhg-1

Region B

P2 1.0

TiO2-700

Region A

-1

-0.8 -1.6

A+B: 110.32 mAhg -1 C: 76.34 mAhg -1

Region B

-1

1.6

TiO2-600

Region A

2.5 2.0 1.5 1.0

Specific Capacity (mAhg ) TiO2-as made

0.1 mV 0.5 mV 1 mV 2 mV 3 mV 5 mV 10 mV

A+B: 90.54 mAhg -1 C: 78.13 mAhg-1

Region B


-1


2.4

TiO2-500

Region A

2.5 2.0 1.5 1.0 3.0

Cycle Number

(D)

A+B: 89.46 mAhg -1 C: 70.86 mAhg-1

Region B

3.0 2.5 2.0 1.5 1.0

20 C

TiO2-as made

Region A

2.5 2.0 1.5 1.0 3.0


+


3.0


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

Page 34 of 36

3.0

6 4 2

TiO2-700

0.1 mV 0.5 mV 1 mV 2 mV 3 mV 5 mV 10 mV

0 -2 -4 -6 -8

1.0

+


1.5

2.0

2.5

3.0

+


Figure 8. (A) Galvanostatic charge/discharge voltage profiles of TiO2-as made, TiO2-500, TiO2600, and TiO2-700 measured at 1 C and (B) corresponding rate capabilities measured at various C


34

Page 35 of 36

rates. (C) Discharge curves of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700 measured at 1 C with relative capacities measured at different voltage regions: Region A (E > 1.75 V), Region B (E ≈ 1.75 V), Region C (E < 1.75 V). Cyclic voltammograms of (D) TiO2-as made, (E) TiO2-500, (F) TiO2-600, and (G) TiO2-700 measured at various scan rates.

(A) Nyquist plots during discharge (2.55 V to 1.53 V)

0.4

TiO2-as made TiO2-500

0.2

0.6 0.4


0.2

TiO2-600

0

TiO2-as made

0.5

TiO2-700

0.6 0.4

TiO2-as made

1.78 V

TiO2-600

TiO2-600

0.4

TiO2-700

0.3 0.2

0

0

0

0.2 0.4 0.6 0.8 1.0

Z (g)

0

0

0.2 0.4 0.6 0.8 1.0

Z (g)

TiO2-700

0.3 0.2 0.1

0.1

TiO2-700

0.2 0.4 0.6 0.8 1.0

1.53 V

TiO2-500

0.4

TiO2-500

TiO2-600

0.2

TiO2-600

TiO2-700

2.03 V

TiO2-500

0.8

-Z (g)

0.6

0

TiO2-as made

0.8

-Z (g)

-Z (g)

0.8

0.5

0.6

1.0 2.35 V

-Z (g)

1.0 2.55 V

-Z (g)

1.0

0

0.15 0.30 0.45 0.60

Z (g)

0

0

0.1 0.2 0.3 0.4 0.5

Z (g)

Z (g)

(B) Nyquist plots during charge (1.48 V to 2.55 V)

TiO2-500

TiO2-500

TiO2-500

0.30 0.15

TiO2-700

0.30 0.15

0

0.15 0.30 0.45 0.60

0

0

10

(D)

10

1

10

0

10

TiO2-as made

2.4

0

-1

TiO2-as made

1.6

Potential (V)

1.2

10

-2

2.8


0

10

2.4

2.0

10

TiO2-700

0.3 0.6 0.9 1.2 1.5


0

TiO2-600 TiO2-700

0

(F)

10

1

10

0

10

TiO2-as made

0.3

0.6

0.9

-1


TiO2-600

TiO2-600

TiO2-600

TiO2-700

TiO2-700

1.6

Potential (V)

1.2

1.2

1.6

2.0

Potential (V)

2.4

1.2

Z (g)

Z (g)

0

-1

0.4 0.2

TiO2-600

0

0.6

TiO2-500

TiO2-500

2.0

0.4

0.15 0.30 0.45 0.60

(E)

TiO2-600

2.8

0.6

Z (g)

TiO2-500 TiO2-700

0.8

0.2

Z (g)

0

-1

0

0.15 0.30 0.45 0.60

RLi (  g)

10

TiO2-700

0.15

Z (g)

(C)

TiO2-600

0.30

Rct (  g)

0

0.45

TiO2-600

-Z (g)

TiO2-700

2.55 V

0.8

-Z (g)

0.45

TiO2-600

1.0

2.10 V

1.95 V TiO2-as made

-Z (g)

-Z (g)

0.45

1.0

0.60 1.71 V TiO2-as made

-Z (g)

0.60 1.48 V TiO2-as made

RLi (  g)

0.60

Rct (  g)

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

TiO2-700

-2

1.2

1.6

2.0

2.4

Potential (V)

Figure 9. In situ EIS measurement results. (A) Nyquist plots of TiO2-as made, TiO2-500, TiO2600, and TiO2-700 measured during discharge at various potentials (2.55, 2.35, 2.03, 1.78, 1.53 V) (B) Nyquist plots of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700 measured during charge at various potentials (1.48, 1.71, 1.95, 2.10, 2.55 V). Various resistances obtained from the fitting at different voltages: (C) Rct and (D) RLi (during the discharge) and (E) Rct and (F) RLi (during the charge). All potentials in the plots are with respect to the standard potential of Li+/Li).


35


Page 36 of 36

Table 1. Lattice Parameters, Grain Size, Surface Area, Total Pore Volume, and ID/IG Ratio of TiO2-as made, TiO2-500, TiO2-600, and TiO2-700. Lattice Parameter (Å)

Grain Size

Total Pore Volume

a

c

(nm)a)

Surface Area (m2g-1)b)

TiO2-as made

3.7868

9.4974

9.9

70.93

0.198

-

TiO2-500

3.7873

9.5077

13.0

93.78

0.295

1.97

TiO2-600

3.7867

9.5067

13.6

88.69

0.290

1.35

TiO2-700

3.7865

9.5066

14.5

83.69

0.287

1.06

IG2/IG1 c)

3 -1

(cm g )

a)

Grain size was determined using the Scherrer equation; b) Surface area was determined using the BET method; c) Ratio of G2 band-to-G1 band peak intensities in the Raman spectra.

Table of Contents Graphic Li+

Carbon

TiO2/Carbon Interface

Li+ e

TiO2

Titanium Oxygen Oxygen Vacancy

Li+


36

Designing Hierarchical Assembly of Carbon-Coated TiO2

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