Electrochemical and Structural Investigation on Ultra-Thin ALD ZnO

Research Article Cite This: ACS Sustainable Chem. Eng. XXXX, XXX, XXX−XXX

pubs.acs.org/journal/ascecg

Electrochemical and Structural Investigation on Ultrathin ALD ZnO and TiO2 Coated Lithium-Rich Layered Oxide Cathodes Chih-Chieh Wang,*,† Jie-Wei Lin,† Yu-Hsuan Yu,‡ Kuo-Hsiang Lai,† Kuo-Feng Chiu,† and Chi-Chung Kei‡ †

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Department of Materials Science and Engineering, Feng Chia University, No. 100 Wenhwa Road, Seatwen, Taichung 40724, Taiwan ‡ Instrument Technology Research Center, National Applied Research Laboratories, 20 R&D Road VI, Hsinchu Science Park, Hsinchu 30076, Taiwan S Supporting Information *

ABSTRACT: Ultrathin coatings (1.5 ± 0.3 nm) of titanium dioxide and zinc oxide were deposited on lithium-rich layered oxide cathodes (Li1.2Mn0.6Ni0.2O2, LLO) by atomic layer deposition (ALD). The structures, electrochemical performances, and thermal stabilities of these coatings were investigated. An ultrathin uniform coating was obtained for TiO2 but not for ZnO because of differences in the layer growth mechanism. Regarding the initial charge−discharge curves under a current density of 0.04 C rate, the TiO2 coated samples exhibited a higher discharge capacity, 242 mAhg−1, compared with the ZnO coated samples, 220 mAhg−1, or the pristine samples, 228 mAhg−1. Both coated samples exhibited more stable cycling performance and thermal stability than the pristine samples. After 80 cycles under 0.5 C rate, the TiO2 and ZnO coated samples were found to have higher capacity retention (∼94% and 78%, respectively) than the pristine samples (68%). The reaction temperature of the exothermic peak of the TiO2 and ZnO coated samples at 4.8 V shifted to 280 °C with heat release of 88.7 J/g for TiO2 and 270 °C with heat release of 154.6 J/g for ZnO. This is compared with an exothermic peak at 258 °C with heat release of 253.5 J/g for the pristine sample. In particular, an enhanced rate capability was only observed for the TiO2 coated samples. When the current densities were higher than 2 C rate, the TiO2 coated samples exhibited superior capacities than the pristine and ZnO coated samples. At a current density of 5 and 10 C rate, the capacities were found to be 120 and 95 mAhg−1. The improved electrochemical performances were mainly attributed to lower resistance of the charge transfer, which resulted from the layer morphology of the TiO2 film. This feature lead to more preactivation of LLO, smoother electron transport, and suppression of more side reactions, when compared with the island structure of the ZnO film. KEYWORDS: Atomic layer deposition, Cathode, Lithium-rich layered oxides, Lithium ion battery, ZnO, TiO2

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efficiency. In addition, at high potentials (>4 V vs Li/Li+) decomposition of LiPF6 salt leads to large amounts of HF which attacks LLO and dissolves transition metals.5,6 Consequently, poor cycle stability is observed. Another issue is that the layered oxides would transform to spinel-like structures, leading to a voltage drop in the extended cycles.7,8 Many efforts have been made to overcome these drawbacks. These efforts include cation substitution7−10 and surface modification of the electrode with spinel structure,11,12 metal oxides,13−15 fluorides,16,17 phosphates,18,19 and carbon.20−22 In particular, surface modification is widely employed to improve the electrochemical performance and structural stability. Among the coating materials, TiO2 and ZnO prepared by wet chemistry processes have been demonstrated not only to

INTRODUCTION

Lithium ion batteries (LIB) are widely used for portable electronics and are believed to be a potential power source for transportation in the future. The requirements of LIB for these applications are high energy density, safety, fast charge− discharge time, and long cycle life. Among the advanced cathode materials, lithium-rich layered oxide (LLO) cathodes, solid solutions of Li2MnO3 and Li[CoxMnyNiz]O2, have attracted attention because they possess higher capacity (>250 mAhg−1).1−3 However, the lithium rich layered oxides suffer from some issues. For instance, the high capacity of the LLO is associated with a long potential plateau at ∼4.5 V which is caused by delithiation of the Li2MnO3 domains accompanied by oxygen release from the lattice.1−3 The released oxygen is very reactive and is believed to react with the electrolyte and form a product at the interface of the electrode and electrolyte.4 This behavior is detrimental to the rate capability, thermal stability, and initial Coulombic © XXXX American Chemical Society

Received: August 28, 2018 Revised: October 31, 2018 Published: November 7, 2018 A

DOI: 10.1021/acssuschemeng.8b04285 ACS Sustainable Chem. Eng. XXXX, XXX, XXX−XXX

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ACS Sustainable Chemistry & Engineering

Figure 1. (a) XRD patterns of all the samples and SEM images of (b) the pristine, (c) P@TiO2, and (d) P@ZnO samples.

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prevent the surface from HF attack but also to facilitate electron transfer and hence improve the cycle life and rate capability for various cathode materials.23−26 Nevertheless, the widely used wet chemical processes have difficulty in controlling the uniformity and conformality of ultrathin films. On the contrary, the self-limiting ALD process enables deposition of precisely controllable thin films with excellent conformality and uniformity on complicated nanostructures.27−29 ALD has been applied to coat anode and cathode materials of LIB.13,30,31 The electrochemical performances of ALD ZnO-coated cathode materials are controversial. For instance, improved cycling performances and rate capabilities are obtained for LiMn2O432 and Li1.2Mn0.54Ni0.13Co0.13O2.33 However, ALD ZnO coated-LiCoO2 exhibits no enhancement.34 Very few studies have been reported for ALD TiO2 coating. Besides, the morphologies of the various coating layers would be different due to the surface reaction controlled process of ALD. The effects of the morphologies of the coating layers on the structure and thermal and electrochemical performances have not been investigated. In this regard, we present here an investigation of the structure, electrochemical performance, and thermal stability of ALD ZnO and TiO2 coated lithium-rich layered oxide cathodes. The results indicate that the ALD TiO2 coating has more advantages in the electrochemical characteristics of LLO than ALD ZnO. This is due to differences in the morphologies. In particular, the rate capability, cycling performance, and thermal stability of the ultrathin TiO2 coated LLO cathodes are shown to be remarkably enhanced. These comparative studies of surface modification will provide new insight into other cathode or anode materials for Li-ion batteries.

EXPERIMENTAL SECTION

Preparation of Lithium-Rich Layered Oxides (Li1.2Mn0.6Ni0.2O2). Li1.2Mn0.6Ni0.2O2 was synthesized by a sol−gel method. The required amounts of lithium acetate (98%, extra pure, ACROS Organics), manganese acetate (99+%, for analysis, ACROS Organics) and nickel acetate (99%, extra pure, ACROS Organics) were dissolved in deionized water. Ethylenediaminetetraacetic acid (EDTA, 99%, pure, ACROS Organics) and citric acid (99%, Pure, ACROS Organics) were mixed with water in a separate beaker to serve as complexing agents. Subsequently, NH4OH (35%, Extra Pure, Fisher Chemicals) was added to the EDTA/citric acid solution. A homogeneous solution was formed as a pH of ∼8 was reached. The metal ion solution was then added dropwise into the EDTA/citric acid solution. The molar ratio of EDTA: citric acid: metal ion was 1:1.5:1. The resulting mixture was then heated at 90 °C for 12 h until the gel phase was formed. Subsequently, the gel was fired at 450 °C for 3 h and then at 900 °C for 24 h to obtain the final cathode powder. The samples without TiO2 and ZnO ALD coating were named as pristine. Fabrication of TiO2-Coated Lithium-Rich Layered Oxide Cathodes by ALD. TiO2 was deposited on the Al foil with a lithiumrich layered oxide cathode by ALD at 100 °C. TiCl4 (≥99.0%, SigmaAldrich) and H2O were used as the precursors. The pulse time was 0.5 and 0.5 s. Between each pulse 10 s of N2 purge was applied. The growth rate of TiO2 was ∼1.5 Å/cycle. Ten cycles of ALD TiO2 was applied to compare with ALD ZnO. The samples were named as P@ TiO2. Fabrication of ZnO-Coated Lithium-Rich Layered Oxide Cathodes by ALD. For ZnO ALD coating, DEZ (diethyl zinc, Sigma-Aldrich) and H2O were used as the precursors. The substrate temperature was kept at 150 °C. The pulse time of DEZ and H2O were 0.2 and 0.1 s, respectively. A N2 purge was applied for 5 s after each precursor pulse. The growth rate of ZnO was 3 Å/cycle. In order to compare with TiO2, 5 cycles of ALD ZnO was applied. The samples were named as P@ZnO. Coin Cell Assemble. Electrochemical performance was evaluated using 2032 type coin cells. The coin cells were assembled with lithium metal as the anode, 1 M LiPF6 (Sigma-Aldrich) in ethylene carbonate (EC, Sigma-Aldrich)/ diethyl carbonate (DEC, Sigma-Aldrich) (v/v B


Research Article

ACS Sustainable Chemistry & Engineering 1:1) was the electrolyte, and Celgard polypropylene was the separator. The cathode was fabricated using 80 wt % active material, 10 wt % Super P carbon (Conductive, 99+%, Alfa Aesar) and 10 wt % polyvinylidene fluoride (PVDF, Sigma-Aldrich) binder in n-methyl-2pyrrolidone (NMP, Sigma-Aldrich) solvent. The loading of the active material was controlled to be 3 mgcm−2. Characterization. The structure of the samples was characterized by a Philips diffractometer using Cu Kα radiation in the range of 10− 80°. The morphology of the samples was examined by a field emission scanning electron microscope (JEOL JSM-7800F). Microstructure analysis was conducted with high-resolution transmission electron microscope (JEOL, JEM-2100F). Surface composition was analyzed by X-ray photoelectron spectroscopy with Al Kα radiation (ULVACPHI, PHI 5000 VersaProbe). The thermal stabilities were examined by differential scanning calorimetry (TA Instrument, TGA2950). The coin cell was initially charged to 4.8 V at 0.04 C rate and was then disassembled in an Ar-filled glovebox. After evaporation of the electrolyte under vacuum, the active material was removed from the electrode. One mg powder was sealed in an aluminum pan. The measurements were conducted from 100 °C to 350 °C with 5 °C min−1 heating rate. A Solartron 1260 impedance analyzer applied an AC amplitude of 10 mV in the frequency range of 100 kHz to 1 mHz to obtain electrochemical impedance spectroscopy (EIS) data. Lithium foil served as the counter and reference electrodes.

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RESULTS AND DISCUSSION Structure and Surface Morphology of ALD TiO2 and ZnO Coated LLO. Figure 1(a) indicates an XRD pattern of the pristine, P@TiO2 and P@ZnO samples. All of the samples possess similar diffraction peaks; meaning there were no structure changes after ALD TiO2 and ZnO coating. The main diffraction peaks at 18.72°, 36.89°, 37.91°, 38.65°, 44.63°, 48.79°, 58.75°, 64.38°, 65.30°, and 68.61° can be assigned as R3̅m structure. An extra peak at ∼20−25° corresponding to the characteristic of the lithium-rich layered oxides was obtained. (108) and (110) diffraction peaks are well separated. This indicates that the ordering arrangement of the transition metal ions in the layered oxides. The layered oxides are well formed because the ratio of the (003) and (104) peaks is larger than 1.2. However, no diffraction peaks of TiO2 and ZnO are observed because the oxide layer is too thin and amorphous. The layered oxides are polyhedral and 200−300 nm in size, as shown in Figure 1(b−d). In order to investigate the surface morphology and element composition of the ultrathin layers, HR-TEM, and EDS-mapping are employed, as shown in Figures 2 and 3. The pristine samples possess a clean and smooth surface. A well-defined boundary and no lattice fringes are observed for the surface of TiO2 and ZnO ALD coated samples. However, the surface morphologies are slightly different. TiO2 coated samples possess more even and uniform layers. The thickness is found to be 1.5 nm, corresponding to the growth rate of TiO2 ALD, 1.5 Å/cycle. On the contrary, aggregated particles are obtained for the ZnO-coated samples (blue arrows in Figure 2(b)) and the thickness varies between 1.2 and 1.8 nm, with average value 1.5 nm. The reasons for that are related to the differences in the nucleation and growth modes. For ALD TiO2, full coverage of an amorphous film appears in the nucleation stage, followed by particle aggregation to grow the film.35 However, ALD ZnO prefers island modes in the nucleation and growth stages, which tends to merge particles and leads to an uneven surface.36 The composition of the ALD TiO2 and ZnO coated samples were analyzed by TEM-EDX image mapping, as shown in Figure 3. The identical signals of Mn, Ni and O corresponding to the lithium-rich layered oxides are detected over the entire surface

Figure 2. TEM images of (a) the pristine, (b) P@ZnO, and (c) P@ TiO2 samples.

of the particles for both samples. The Ti and Zn signals are found and uniformly distributed for TiO2 and ZnO coated samples. The EDX spectra indicate that Ti and Zn signals appear on the edge of both samples, confirming the presence of a TiO2 and ZnO coating on the surface. Initial Charge−Discharge Capacities and Cycling Performances of the Pristine, P@TiO2, and P@ZnO. The initial charge−discharge capacities of the pristine P@TiO2 and P@ZnO are investigated, as shown in Figure 4(a, b). All of the samples possess a slope and plateau region in the initial charge from 2.0 to 4.8 V under C/25 rate. The capacity in the slope and plateau region are related to the oxidation of Ni from +2 to +4 and the oxygen extraction from the lattice, respectively.1−3 Interestingly, the P@TiO2 samples exhibite a higher discharge capacity (242 mAhg−1) than the pristine samples (228 mAhg−1). However, as ALD ZnO is deposited on the pristine samples, the discharge capacity (220 mAhg−1) is inferior to the pristine samples. This implies that the TiO2 layer may possess higher electronic and ion conductivities than ZnO because the TiO2 coating is more uniform. Figure 4(c) shows the cycling performances of the pristine, P@TiO2, and C


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Figure 3. Image of element mappings and EDX spectra of (a) P@TiO2 and (b) P@ZnO samples.

P@ZnO samples under a current density of 0.5 C rate with 80 charge−discharge cycles. Better capacity retention is observed for both samples. The capacity retention of the P@TiO2 and P@ZnO samples are found to be 94 and 78%, compared with

the pristine samples, 68%. Furthermore, the P@TiO2 samples possess a higher Coulombic efficiency (∼99%) than the P@ ZnO (∼98%) and pristine samples (∼97%) after 80 cycles at 0.5 C rate, as shown in Figure S1 of the Supporting D


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Figure 4. (a) Initial charge−discharge curves, (b) the corresponding histograms of the first charge, discharge and irreversible capacity loss at the current density of 0.04 C rate, (c) cycling performances at the current density of 0.5 C rate, and (d) rate performances of the pristine, P@ZnO, and P@TiO2 samples at 0.04 C to 1 C rate during 2.0 to 4.8 V.

Figure 5. Charge−discharge curves of (a) the pristine, (b) P@TiO2, and (c) P@ZnO samples and differential capacity (dQ/dV) plot of (d) the pristine, (e) P@TiO2, and (f) P@ZnO samples with various cycles.

results suggest that the ALD TiO2 coating can prevent more structural transformations and side reactions from occurring on the surface of the active material. Interestingly, Figure 4(d) indicates that the P@TiO2 samples exhibit a relatively high

Information (SI). The capacity loss and lower Coulombic efficiency after the cycling may be attributed to the transformation of the layer structure to spinel phases and to side reactions between the electrode and electrolyte. The E


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Figure 6. XPS spectra of oxygen (1s) and F (1s) (a, c) before and (b, d) after cycling for the pristine, P@TiO2, and P@ZnO samples at the current density of 0.5 C rate from 2.0 to 4.8 V.

surface Ni ions diffuse into the bulk to occupy the lithium vacancies and this leads to the structure rearrangement and the decrease of the electrochemical activities of Ni ions. Hence, the results indicate that this phenomenon can be better alleviated by a TiO2 ALD coating than ZnO. At the initial charge, all of the samples exhibit another reduction peak at ∼3.3 V corresponding to the reduction of Mn4+ to Mn3+ in the layered oxide structures.39 In the extended cycles, the intensity of the reduction peak decreases and shifts to lower voltage for all the samples. In the case of the P@TiO2 samples, the peak shift is less pronounced and stops at ∼3.1 V, which corresponds to the reduction of Mn4+ to Mn3+ in layered structures with spinel-like characteristics.39 The cathodic peak between 2.0 to 2.4 V corresponding to the lithiation of TiO2 to LixTiO2 were absent.40,41 This implies that no capacity comes from TiO2. However, other reduction peaks at ∼2.5 V and ∼2.8 V, which are associated with the reduction of Mn4+ to Mn3+ in the spinel structures, appear for the pristine and P@ ZnO samples.39 This means that the structure transformation of the layered to spinel phases during the cycling is less favorable for the P@TiO2 samples. Both of these features indicate that the P@TiO2 samples are more stable than the pristine and P@ZnO samples. Consequently, there is less polarization and loss of capacity. Surface Composition of the Pristine, P@TiO2, and P@ ZnO before and after Cycling. In order to understand the role of the various surface modifications in the enhanced stability of the lithium-rich layered oxide cathodes, XPS spectra of O (1s) and F (1s) for the pristine, P@TiO2, and P@ZnO samples before and after 80 cycles at 0.5 C rate are investigated in Figure 6. The O (1s) spectra of the samples before cycling could be well fitted with two peaks as shown in Figure 6(a).

capacity at 0.5 C rate than Figure 4(c). The reason for that is due to the differences of the charge−discharge steps. No electrochemical activation process at low C rate is used before the actual cycling at 0.5 C rate for Figure 4(c). However, the P@TiO2 samples possess 13 cycles with lower C rate that can act as the electrochemical activation process before the 0.5 C rate reaches. The activation process at the initial charge for the lithium-rich layered oxides leads to the structure rearrangement to reduce the diffusion impedance and increase the diffusion coefficient of the lithium ion.37 Therefore, the higher capacity of the P@TiO2 samples is observed for Figure 4 (d). Charge−Discharge Curves and Differential Capacity (dQ/dV) Plots of the Pristine, P@TiO2, and P@ZnO with Various Cycles. Figure 5 (a−c) indicates the charge− discharge curves of all the samples with various cycles under 0.5 C rate. When the charge−discharge cycles increase to 80, the polarization loss is more pronounced. The P@TiO2 samples exhibit less polarization loss (ΔE = 0.70 V) when compared with the pristine samples (ΔE = 1.67 V) and P@ ZnO (ΔE = 1.50 V). To investigate the polarization and capacity loss in the extended cycles, dQ/dV plots of the pristine, P@TiO2 and P@ZnO samples are shown in Figure 5(d−f). At the first discharge, the reduction peaks of the pristine samples between 3.70−4.40 V, corresponding to the reduction of Ni ion from 4+ to 3+/2+, shift to a higher voltage when TiO2 and ZnO ALD are used. When the charge− discharge cycle increases to 80, the reduction peaks become weak and shift to lower voltage for the P@TiO2 samples but are absent for the pristine and P@ZnO samples. According to the study reported by Armstrong et al., at the initial charge, the excess lithium ions would be extracted from the transition metal ion layer followed by oxygen loss from the lattice.38 The F


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The transformation of the layer to spinel structure is induced by the following reaction:

The major peak at 529.4 eV is related to the lattice oxygen in the cathode material (red dash line). Another peak at 531.2 eV (blue dash line) corresponds to Li2CO3 that results from the spontaneous reaction of the sample surface with CO2 in the atmosphere.42 The area ratio of the Li2CO3 peak to the lattice oxygen for pristine, P@TiO2, and P@ZnO is estimated to be 42.5%, 29.8%, and 35.2%, respectively. As is well-known, an SEI layer (ROCO2Li, polyether, Li2CO3 and LiOH) with a considerable thickness would be formed through the decomposition of the electrolyte during the charge−discharge cycles. The passivation layer is believed to be detrimental to the capacity, kinetics, and polarization, and it contributes to the resistance. After 80 cycles, the peak at ∼531 eV associated with the SEI layer gets strengthened and the lattice oxygen content is reduced. This indicates that the SEI layer becomes thicker because there is more electrolyte decomposition on the surface of the active material after cycling. The area ratio of the peak of the SEI layer to the lattice oxygen for the pristine, P@ TiO2, and P@ZnO samples increased to 91.4%, 77%, and 88.2%, respectively. The interaction of the electrolyte with active material to form metal fluoride byproducts also contributes to the growth of the SEI layer. Hence, the XPS spectra of F (1s) of the pristine, P@TiO2, and P@ZnO samples, before and after cycling, are shown in Figure 6(c, d). All of the samples possessed similar features of the peaks. The peak at 688.2 eV (blue dash line) corresponding to C−F originating from the binder (PVDF) was observed before the cycling.43 After 80 charge−discharge cycles at 0.5 C, the intensity of the peak gets weaker and an additional peak at 684.5 eV (red-dash line) associated with F-anion resulting from MFx byproducts appears.43 The area ratio of the peak at 684.5 to 688.2 eV for the pristine, P@TiO2, and P@ZnO samples are found to be 90.1%, 76.8%, and 81.6%. The formation of the SEI layers begins with the decomposition of the LiPF6 electrolyte. In fact, Li salt always possesses small amounts of water that can facilitate breakdown of the electrolyte and generate HF by the following reactions:44,45 LiPF6(sol) ↔ LiF(s) + PF5(sol)

(1)

PF5(sol) + H 2O → POF3(sol) + 2HF(sol)

(2)

2Mn+3 → Mn+2 + Mn+4

where is more favorable in the presence of Mn3+ and acid environment.46 Regarding the XPS of F(1s) spectra (Figure 6(d)) and dQ/dV plots of the P@TiO2 samples (Figure 5(e)), less fluoride compounds and reduction of Mn4+ to Mn3+ at the initial charge are observed for the P@TiO2 samples. This implies that the uniform layer of the ALD TiO2 can effectively scavenge HF and possess less Mn3+ to suppress the reaction. Figure S2 presents the SEM and TEM image of the P@TiO2 samples after 50 cycles at 0.5 C rate. Indeed, they clearly show that the composite material and smooth edge of the ALD TiO2 layer maintain. Meanwhile, the ex-situ XRD spectra after 50 cycles at 0.5 C rate indicate that the spinel phases are less pronounced for the P@TiO2 samples, compared with the pristine and P@ZnO samples as shown in Figure S3. The reasons and evidence demonstrate the improvement of the structure stability of the P@TiO2 samples during the cycling. Rate Capability of the Pristine, P@TiO2, and P@ZnO. The P@TiO2 samples exhibited a capacity comparable with the pristine and P@ZnO samples when the C rate is lower than 0.5, as shown in Figure 7(a). It can be clearly observed that the performances of the rate capabilities of the P@TiO2 samples are superior to the pristine and P@ZnO samples when the C rate increases from 0.5 to 10. For instance, at 5 and 10 C rate, the P@TiO2 samples exhibit a relatively high capacity, 120 and 95 mAhg−1, but the capacity is absent for the pristine and P@ZnO samples. Alternatively, the P@ZnO samples exhibit rate capability performance that is inferior to the pristine samples when C rate is higher than 0.5 C. Electrochemical Impedance Spectra (EIS) of the Pristine, P@TiO2, and P@ZnO. In order to understand the possible reasons for the enhanced electrochemical performance for the P@TiO2 samples, an EIS of the pristine, P@TiO2, and P@ZnO samples was performed after 1 and 50 cycles at 0.2 C rate, as shown in Figure 7(b, c). Two semicircles in the high and medium frequency region are observed for all of the samples. The first semicircle in the high frequency region represents lithium ion diffusion across the interphase of the electrode/electrolyte.47 The second semicircle in the medium frequency range is related to the charge transfer region between the surface film and the interface of the active material.47 An equivalent circuit is used to quantify the polarization behavior for the EIS measurement, as shown in the inset of Figure 7(b, c). Here, Rs and Rct refer to the surface and charge transfer resistance, which are shown in Table 1. It appears that the P@TiO2 samples possessed the minimum Rs and Rct values at the 1st and 50th cycle, indicating that lithium ion diffusion between the solid/liquid interface and electron transfer between the surface film and the interface of the active material becomes much easier. This also suggests that the P@ TiO2 samples possess more active sites for lithium insertion/ extraction and open channels for lithium ion diffusion, resulting in excellent high rate capability, low potential polarization, and higher discharge capacity. On the contrary, the maximum Rs and Rct values are obtained for the P@ZnO samples. This explains the poor rate capabilities and lower discharge capabilities. The higher resistances are attributed to the morphologies of the ALD ZnO film. Compared with the layer of TiO2 film, more pinhole defects are present for the

The HF attacks the cathode to form metal fluoride related byproducts. The reactions can be represented by the following: Li[Li 0.2Mn 0.6Ni 0.2]O2 + HF + Li+ + e− → Li1 − x[Li 0.2Mn 0.6Ni 0.2]O2 + LiF + H 2O

MnO + 2HF → MnF2 + H 2O

(3)

NiO + 2HF → NiF2 + H 2O

(4)

The byproducts then precipitate onto the cathode surface, leading to increased resistance. Regarding the XPS of O (1s) and F(1s) spectra, the aforementioned reactions would occur much easier for the pristine samples due to their bare surface. Alleviation of such a decomposition reaction would be more favorable for the P@TiO2 than the P@ZnO samples. The reason for that is attributed to the morphologies of the TiO2 and ZnO films. A layer of TiO2 film growth can provide pinhole free protection of the electrode. On the contrary, more defect pinholes are formed due to the island growth of the ZnO film. These defects result in ineffective HF scavenging and lead to incomplete protection for the bare samples. G


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Figure 8. Scheme of the (a) P@ZnO and (b) P@TiO2 samples during the charge−discharge process.

of the P@ZnO sample which decreases from 4598 Ω to 4352 Ω. The decrease of the Rct value can be also observed in Al2O3, AlF3 coated, and surface nitridated lithium-rich layered oxides.48−50 This phenomena results from the electrochemical activation of the lithium-rich layered oxides in the initial charge.48−51 The high rate capability of the layered oxides is difficult to reach because of the extremely small diffusion coefficient of the lithium ion in the initial charge (10−19 cm2s−1).52 After the initial charge, lithium extraction and oxygen loss from the lattice occurs, and this leads to the structure rearrangement. The paths for ion diffusion are then established to reduce the diffusion impedance and consequently increase the diffusion coefficient of the lithium ion.37 Hence, the results suggest that the TiO2 ALD coated surface can facilitate the electrochemical activation process because the pinhole free coverage and acidity of the TiO2 layer can efficiently extract lithium ion from the surface of the active material. The electrochemical activation in the initial charge process can be then ignored and the diffusion impedance can be significantly reduced. In the case of ALD ZnO coating, the lower acidity of ZnO leads to the less extraction of the lithium ion in the preactivation process.53 In addition, many pinholes are present in ALD ZnO coatings. So, these coatings are less able to prevent HF attack, and this results in a thicker SEI layer. A combination of both reasons explains the decrease in the Rct values that are observed. DSC Measurement of the Pristine, P@TiO2, and P@ ZnO. The effects of ALD TiO2 and ZnO coating on the thermal stability of the cathodes charged to 4.8 V are investigated as shown in Figure 9. The reaction temperature of the exothermic peak of the pristine and P@ZnO samples are

Figure 7. (a) The rate capability and electrochemical impedance spectra of the pristine, P@ZnO, and P@TiO2 after (b) 1 and (c) 50 cycles of charge−discharge at the current density of 0.5 C rate from 2.0 to 4.8 V. The inset was the equivalent circuit.

island ZnO film to retard the electron transport and lithium ion diffusion as shown in Figure 8. Interestingly, when the charge−discharge cycles increase from 1 to 50, the Rct value of the P@TiO2 samples significantly decreases from 1984 Ω to 730 Ω, compared with the Rct value

Table 1. Simulated Values of Rs and Rct for the Pristine, P@TiO2, and P@ZnO Samples at 1st and 50th Cycles by the Equivalent Circuit ΔRs

Rs (Ω) sample

1 cycle

50 cycles

pristine P@TiO2 P@ZnO

520 37 696

842 39 814

+61.9% +5.4% +16.9%

ΔRct

Rct (Ω) 1 cycle

50 cycles

2207 1984 4598

4052 730 4352 H

+83.5% −63.2% −5.4%

ΔRtotal

Rtotal (Ω) 1 cycle

50 cycles

2727 2021 5294

4894 769 5166

+79.4% −61.9% −2.4%



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ACKNOWLEDGMENTS

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REFERENCES

The work was supported by Ministry of Science and Technology of Taiwan under contracts MOST 105-2221-E035-026 and MOST 107-2221-E-035−012 -MY3.

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Figure 9. DSC profiles of the pristine, P@ZnO, and P@TiO2 samples charged to 4.8 V.

found to be 258 °C with a heat release of 253.5 J/g and 270 °C with a heat release of 154.6 J/g, respectively. In particular, when ALD TiO2 is applied, the reaction temperature shifts to 280 °C and the heat is reduced to 88.7 J/g. This indicates that the ALD TiO2 coating rather than ALD ZnO could provide more improvement in the thermal stability of the lithium-rich layer oxide cathodes. The improved thermal stability is associated with the pinhole free film, which is better at suppressing the interface reaction between the electrode and electrolyte. Moreover, the coated TiO2 delithinated samples have less oxygen loss from the lattice due to the preactivation process, which improves the stability at the charge states.

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CONCLUSIONS The electrochemical and structure characteristics of the ALD TiO2 and ZnO coated lithium-rich layered oxide cathodes were investigated. Both coated samples exhibited a more stable cycling performance and thermal stability than the pristine samples. However, the initial charge−discharge capacities and polarization loss of the TiO2 coated samples are superior to the ZnO-coated ones. In addition, the much-enhanced rate capability can be only observed for the ALD TiO2 coated samples. This is attributed to the lower resistance of the charge transfer, which results from the pinhole free coating of the ALD TiO2 film.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acssuschemeng.8b04285. The Coulombic efficiencies under 0.5 C rate during 2.0 to 4.8 V and ex-situ TEM, SEM images and XRD spectra of the pristine, P@TiO2, and P@ZnO samples after 50 cycles at 0.5 C rate during 2.0 to 4.8 V (PDF)

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AUTHOR INFORMATION

Corresponding Author

*Tel: 886-4 245-17250. Fax: 886-4 245-10014. E-mail: [email protected]. ORCID

Chih-Chieh Wang: 0000-0002-8939-0871 Notes

The authors declare no competing financial interest. I


Research Article


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Electrochemical and Structural Investigation on Ultra-Thin ALD ZnO

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