Understanding De-protonation Induced Formation of Spinel Phase in Li-rich Layered Oxides for Improved Rate Performance①

LI Bo-Yun LI Gung-She ZHANG Dn FAN Jin-Ming FENG To LI Li-Ping

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Understanding De-protonation Induced Formation of Spinel Phase in Li-rich Layered Oxides for Improved Rate Performance①

LI Bao-YunaLI Guang-SheaZHANG DanaFAN Jian-MingbFENG TaoaLI Li-Pinga②

a(130012)b(350002)

Constructing layered-spinel composites is important to improve the rate perfor- mance of lithium-rich layered oxides. However, up to now, the effect of microstructure of composites on the rate performance has not been well investigated. In this study, a series of samples were prepared by a simple protonation and de-protonation for the pristine layered material (Li1.2Mn0.52Ni0.2Co0.08O2) obtained by sol-gel method. The characterizations of XRD, Raman and oxidation-reduction potentials of charge-discharge curves demonstrated that these samples after de-protonation are layered-spinel composites. When these composites were tested as a cathode of lithium-ion batteries, the sample treated with 0.1 M of nitric acid exhibited higher discharge capacities at each current density than that of other composites. The outstanding rate performance is attributed to the high concentration of conduction electron resulting from the low average valence state (44.2% of Ni3+) as confirmed by its high conductivity (1.124 × 10-2Ω-1×m-1at 39800Hz) and ambient temperature magnetic susceptibility (8.40 × 10-3emu/Oe×mol). This work has a guiding significance for the synthesis of high rate performance of lithium battery cathode materials.

protonation and de-protonation, layered-spinel composites, rate performance, conduction electron;

1 INTRODUCTION

Li-rich Mn-based layered oxides (LMROs) are well deemed as one of the most promising cathode materials because of their high energy density (as high as 900 Wh×kg-1) and low cost[1-4]. However, the electronically insulated Li2MnO3-like component, which is integrated with LiMO2matrix, weakens the conductivity of Li-rich Mn-based layered com- pounds and thus results in a poor rate performance of the cathodes[4-6]. Most of researchers[7-11]chose to construct layered-spinel composites to improve the rate performance of lithium-rich layered materials owing to 3pathway channels in spinel-related structures that are conducive to Li ion and electron transport.

Until now, many methods have been developed to synthesize layered-spinel composites. For example, composites can be prepared by coating spinel phase on the surface of Li-rich layered materials[12-14]. Certainly, composites could also been directly syn- thesized via a one-step strategy of solvothermal[15], sol-gel[11, 16], partial reduction reaction[17]and self- combustion reaction[18]. Recently, Zhao.[19]reported ion-exchange (Li+–H+exchange in Li1.2Mn0.54Ni0.13Co0.13O2in 2 M HCl solution) strategy and calcination (500oC for 6 h in air) process to obtain layered-spinel composites, which is helpful to understand the formation of layered- spinel composites. In these reports, layered-spinel composites greatly improve the rate performance compared with the originally layered materials. However, the relationship between microstructure and rate performance of layered-spinel composites is not still well studied.

Herein, a series of layered-spinel composites are designed to prepare by the processes of protonation and de-protonation. Scheme 1 illustrates our stra- tegy for generating layered-spinel composites, involving an ion-exchange step (protonation) and a subsequent post-heat treatment (de-protonation). In the pristine material obtained by sol-gel method, most of Li ions occupy the 2and 4sites, and part of Li ions locate at 2and 4sites of TM layer. Phase conversion from the initial layered structure to the spinel phase is first carried out by protonation of pristine powder in various concentration acidic solution (HNO3),, Li+–H+exchange (extraction of Li+and partial O species, intercalation of H+). Subsequently, the complete phase transition is reali- zed by de-protonation of the embedded H+via calcinations in O2atmosphere,, release of H+from Li site, complement of oxygen and formation of Li vacancies, leading Li and TM to migrate from octahedral site through tetrahedral site and rear- range to form composites with layered and Li4Mn5O12-type spinel structure. The structural changes are characterized by XRD, Raman and oxidation-reduction potentials of charge-discharge curves. Profiting from the proper nitric acid con- centration and 3channel of spinel structure, the sample treated with 0.1 M of nitric acid concentra- tion exhibits higher capacity and better rate perfor- mance. Furthermore, the reasons for better rate performance were systematically researched.

Scheme 1. Illustration of protonation and de-protonation for the formation of layered-spinel composites

2 EXPERIMENTAL

2.1 Sample preparation

Li1.2Mn0.52Ni0.2Co0.08O2oxide was successfully synthesized via a sol-gel method using a citric acid as the chelating agent. The detailed preparation procedure could be described as follows. First of all, given amounts of lithium acetate (5% excess), transition metal acetates and citric acid (with a molar ratio of 1:1 for transition metal ion to citric acid) were dissolved in 200 mL distilled water to get a mixed solution. The pH value of this mixed solution was adjusted to 8～9 using ammonium hydroxides. Secondly, the solution was dried in a water bath at 80 ℃ for 8 h to obtain a gel. Then, the gel was dried in an oven at 180 ℃ for 12 h to evaporate the residual water. The obtained powder was calcined in air at 500 ℃ for 5 h and then at 900 ℃ for 12 h. Finally, the calcined powder was cooled down to room temperature in furnace to obtain pristine sample (named as sample P).

Protonation of Li1.2Mn0.52Ni0.2Co0.08O2in acidic solution (HNO3) and de-protonation of its ion- exchanged derivatives: pristine Li1.2Mn0.52Ni0.2Co0.08O2powder (1 g) was put in 150 mL aqueous HNO3with various concentrations (0.1, 0.2, 0.3, 0.4 M), and then the mixture kept stirring at 55 ℃ for 48 h. The obtained powder was washed with DI water for several times and dried at 70 ℃ in an oven. The corresponding pro- tonation derivatives were named as A1, A2, A3 and A4, respectively. Finally, these derivatives were calcined at 500 ℃ for 6 h in O2. The resulted de-protonation samples were layered-spinel com- posites, named as C1, C2, C3 and C4, respectively.

2.2 Sample characterization

Crystallographic structures of all samples were identified by powder X-ray diffraction (XRD) on a Rigaku miniflex apparatus with a Cu(= 1.5418 ?) radiation source. XRD patterns were collected in the 2range from 10oto 70o. The chemical com- positions of samples were analyzed by inductively coupled plasma atomic emission spectrometry (ICP-AES). X-ray photoelectron spectroscopy (XPS) analysis was carried out with an ESCALAB250 X-ray Photoelectron. Charging shift was calibrated using C1photoemission line at a binding energy of 284.8 eV. Raman spectra were collected on a Raman spectrophotometer (INVIA) with the laser beam of 532 nm. The magnetic susceptibility was measured using the vibrating sample magnetometer option of a Quantum Design MPMS SQUID-VSM in the temperature range from 2 to 300 K. Magnetic field is 1000 Oe. Conductivity was measured using AC impedance method. All calcined pellets were polished by sand paper until smooth and flat enough. The thickness of the pellets is from 1.1 to 2 mm, and the diameter falling in the 9.2～9.5 mm range. A thin layer of silver paint was smeared on both sides of the pellet as electrodes. The measurement was carried out via impedance spectroscopy on a Solartron 1260 impedance analyzer over the fre- quency range from 10 to 1 MHz and an applied AC voltage of 100 or 200 mV from room temperature to 140 ℃.

2.3 Electrochemical test

Electrochemical performance of samples C1, C2, C3 and C4 was tested using CR2025-type coin cells. The cathode slurry was made with active material, carbon black and binder (polyvinylidene fluoride (PVDF), Alfa Aesar) in a weight ratio of 8:1:1 in NMP. The mixed slurry was coated on Al foil and dried at 80 ℃ overnight. The coin cells were assembled in an argon-filled glove box (both H2O and O2< 1 ppm). The metallic lithium foil was used as the counter electrode and Cellgard 2400 as the membrane. The commercial electrolyte comprised 1 M LiPF6in a mixed solvent of ethylene carbonate (EC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) (1: 1: 1 in volume). The coin cells were tested on a Neware Battery Test System in the potential window between 2.0 and 4.8 V at 30oC. The rate performance was measured at different current densities of 20, 100, 200, 400, 600 and 800 mA/g, and then returned to cycling at 20 mA/g. Electrochemical impedance spectroscopy (EIS) was investigated using an electrochemical workstation (CHI660C) in a frequency from 100 kHz to 0.01 Hz with AC voltage amplitude of 5 mV.

3 RESULTS

3.1 Structure

The phase structures of all samples were firstly determined by XRD. Fig. 1 shows the XRD patterns of all samples. All strong diffraction peaks of sample P can be indexed to a hexagonal layered structure (space group-3). While the weak diffraction peaks at 20～25oare associated with the LiM6cation ordering that occurs in the transition metal layers of Li2MnO3, which can be indexed using a monoclinic (/2) model[20, 21]. The strong and sharp diffraction peaks indicate good crys- tallinity of the sample synthesized by sol-gel me- thod. The clear separations of (006)/(012) and (108)/(110) doublets suggest the well-ordered layered structure of material[22, 23]. In the process of protonation, Li+–H+exchange, accompanied with the release of Li2O[24]from Li2MnO3, which is equivalent to the process of pre-activation. Thus, the XRD of the protonation derivatives obtained by treating the pristine powder with different con- centrations of nitric acid was slightly different. Almost all of the diffraction peaks are the same as those in the sample P except for the presence of a spinel-like phase near (003), which is consistent with the results reported in literatures[19, 25]. The protonation process provides the prerequisite for the subsequent phase transition. In the process of de-protonation, H+ions with smaller diameter are easily released under high temperature, which results in the formation of Li vacancies. Then, Li and TM ions migrate from octahedral site through tetrahedral site, resulting in structural rearrangement and the formation of composites with layered and Li4Mn5O12-type spinel structure. Such a phase transition is highly feasible due to structural compatibility of cubic close-packed oxygen arrays in layered and spinel structures. XRD patterns of layered-spinel composites in comparison with those of pristine are shown in Fig. 1b. The main diffrac- tion peaks of resultant products can be indexed as a spinel Li4Mn5O12structure with F-3space group. The peaks at 18.8°, 36.5°, 44.4°, 48.6°, 58.7° and 67.9° are assigned to (111), (311), (400), (331), (511)and (531), respectively. However, the splitting XRD doublets with selected enlarged portions at 38°～40° and 64°～66° indicate some layered features are preserved. This is mainly due to the lower concentration of nitric acid that only gave partial protonation for sample P. After de-protona- tion, partial occurrence of phase transition from layer to spinel produced layered-spinel composites.

Fig. 1. XRD patterns of the samples. a) Intermediates after the process of protonation (A1, A2, A3, A4) compared to sample P. The asterisk indicates the spinel-like phase. b) Layered-spinel composites after de-protonation (C1, C2, C3, C4) compared to sample P. The dashed rectangle indicates the preserved layered feature

Because the vibration models of Mn–O bonds are different in the layered and spinel structures, the Raman spectroscopy was applied to study the structural changes of electrode materials. Raman spectra of pristine Li1.2Mn0.52Ni0.2Co0.08O2and layered-spinel composites are shown in Fig. 2. For pristine sample, three peaks at 300～450 cm-1belong to the typical monoclinic Li2MnO3com- ponent, and two strong peaks at 490 and 607 cm-1are assigned to the bending Egand stretching A1gmodes of LiMO2[26, 27]. Compared to 607 cm-1for pristine Li1.2Mn0.52Ni0.2Co0.08O2, high wavenumber Raman band of layered-spinel composites shifts to around 625 cm-1. This band is associated with the symmetric Mn–O stretching mode in the spinel structure[28, 29], suggesting the formation of spinel phase. Similar shift for Raman peaks was also reported for the generation of spinel structure during cell cycling[30]. The Raman peaks at 300～450 cm-1belonging to the monoclinic Li2MnO3component in pristine disappeared in samples C1–C4. This is because the Li2MnO3component releases Li2O in the process of protonation, resulting in a change of the coordination environ- ment around the Mn ions, which is in agreement with the XRD analysis (the disappeared charac- teristic diffraction peaks of Li2MnO3for samples C1–C4). Similar to pristine sample, C1–C4 samples exhibit the Raman band at 490 cm-1, angmode of LiMO2structure, thus indicating the samples after protonation and de-protonation preserve partial layered feature. Therefore, Raman data also verify that samples C1–C4 obtained through the process of protonation and de-protonation are layered-spinel composi tes.

Fig. 2. Raman spectra of samples. The large shift of high wavenumber mode compared with pristine and the presence of Egmode suggest that samples C1～C4 are layered-spinel composites

3.2 Electrochemical performance

Electrochemical performances of samples P, C1, C2, C3 and C4 were tested by a galvanostatic charge-discharge technique between 2.0 and 4.8 V at ambient temperature. Fig. 3a and 3b show the first charge-discharge curves and the corresponding differential capacity (dQ/dV) plots at a current density of 200 mA×g-1. We can see that sample P exhibits a typical charge-discharge feature of Li-rich layered oxide: (1) the first sloping stage around 3.5～4.4 V in the charge data is assigned to lithium extraction from lithium layer of LiMO2accompanied by the oxidation of Ni2+to Ni4+and Co3+to Co4+[31], and another relatively longer pla- teau region at 4.5 V corresponds to the activation of Li2MnO3-like structure accompanied with a con- comitant oxygen loss from the lattice[27, 32]; (2) in the discharge process, the voltage continually decreases to show a discharge capacity of 176.1 mAh/g in the voltage window of 4.8～2.0 V and coulombic efficiency is only 67.16%. Similar to sample P, samples C1, C2, C3 and C4 also show two step charge processes: a sloping stage around 3.5～4.4 V and plateau region above 4.5 V. For C3 and C4, this plateau voltage is as high as 4.65 V. In their discharge curves, long discharge plateau region at 2.72 V is observed, which belongs to the characteristic plateau of spinel[33, 34], very different from the plateau of sample P. The above analyses further confirm that the samples after protonation and de-protonation are layered-spinel composites. Moreover, the dQ/dV curves also show the dif- ference of charge-discharge between pristine sample and those protonation and de-protonation ones. The corresponding redox potentials for the first dis- charge plateau region are shown in region III in Fig. 3b. The layered-spinel composites have obvious reduction peak at 2.72 V, while the pristine layered material does not. Redox potentials shown in regions I and II of Fig. 3b are the characteristic of layered structure, corresponding to slope step around 3.5～4.4 V and plateau at 4.5 V in the first charge-discharge plot for all samples. It should be noted that in region II, the oxidation potential is 4.55 V for samples C1 and C2, while 4.65 V for samples C3 and C4. The higher oxidation potential in the electrode of samples C3 and C4 suggests that those samples are more difficult to activate[35]. These post-treated samples have more vacancies and 3channel, and Li ions are easily embedded in the structure, so that the first discharge capacities are much higher than the pristine. These discharge capacities of the first cycle at 1.0C are 176.1, 217.5, 227.4, 212.1 and 223.2 mA×h/g for P, C1, C2, C3 and C4, respectively. Fig. 3c displays a comparison of rate performance for all samples. The electrodes are cycled five times at different current density of 0.1C, 0.5C, 1C, 2C, 3C and 4C (1C = 200 mA/g) and then returned to cycling at 0.1C. Among layered-spinel composites, the sample C1 dealt with 0.1 M of nitric acid exhibits the best rate per- formance. The discharge capacities are 259.1, 226.5, 211.4, 193.1, 173.2 and 150.0 mA h/g for 0.1C, 0.5C, 1C, 2C, 3C and 4C, respectively, much higher than those of the other samples. The normalized discharge capacities are shown in Fig. 3d. It is clear that sample C1 owns the best rate performance among all samples, while sample C3 shows the worst rate performance. We also study the rate performance and EIS of the sample treated with 0.05 M of nitric acid and find that its rate per- formance and resistance values are close to P. That is to say, the lower concentration of nitric acid has less effect on the sample. To investigate the rela- tionship between microstructure and rate perfor- mance of layered-spinel composites, we will discuss the reason of the better rate performance for sample C1 through the characterizations of SEM, XPS, AC impedance and magnetism measurement.

4 DISCUSSION

4.1 Morphologies of the samples

It is reported that the morphology of cathode material has an important effect on the rate per- formance. Treating the sample surface[36]or reducing the particle size[37]could improve the rate performance of cathode material. Therefore, it is highly necessary to examine the morphology of all samples. Fig. 4 shows the SEM of all samples. The surface of pristine Li1.2Mn0.52Ni0.2Co0.08O2powders is smooth compared with the layered-spinel com- posites. The surface of all integrated samples is rough and porous due to the activation of acid treatment and the high temperature calcinations. The particle size of all samples is in a range of 800～1200 nm, and there is no difference in magnitude. Therefore, the effect of size could be ruled out. On the other hand, the roughness and porousness of the surface could slightly improve their rate performance, but couldn’t be dominated factors to explain the best rate performance observed for sample C1.

Fig. 3. Electrochemical performances. a, b) Charge-discharge curves of the initial cycle at 1C and corresponding derivative (dQ/dV) plots, c, d) Comparison of the rate performance and normalized rate performance for all samples

Fig. 4. SEM images of all samples

4.2 X-ray photoelectron spectroscopy

In order to determine the influence of valence state of the transition metal ions on the rate per- formance, XPS data are analyzed in Fig. 5. Generally, the 3orbital of Mn ions is more sensi- tive to characterize their valence state. The splitting of the double peaks in Mn3core levels for P and C1–C4 are all at approximate 4.3～4.6 eV, demons- trating the presence of Mn4+primarily[38, 39]. Co 2core level shows four representative peaks for P and C1–C4, including two spin-orbital splitting peaks and two weak satellite peaks. The strong peak at approximate 780 eV is assigned to Co 23/2, and that at 795 eV is attributed to the Co 21/2. The spin- orbital splitting between 23/2and 21/2is about 15.0 eV, implying the valence state of Co ions is +3 for all samples[40, 41]. Because splitting of Mn3double peak and Co2p3/2-21/2of the sample C1–C4 are essentially the same as sample P, the valence state of Mn and Co ions remain unchanged after protonation and de-protonation. For Ni ions, it has been previously reported that the binding energy of Ni2+in Li1.080Mn0.503Ni0.387Co0.030O2is 854.5 eV[42]and that of Ni3+in LiNiO2is 856.0 eV[43]. In our work, the binding energy of Ni 2p3/2is about 855.18 eV, locating in the range of 854.5～856.0 eV, which implies that the chemical valence state of Ni ions comprises of +2 and +3[6]. From the fitting results, the content of Ni3+is about 40% in P, consistent with the predict value from its chemical formula. Due to the exchange of Li+–H+in protonation process and release of H+at 500oC during de-pro- tonation, the partial divalent nickel ions would be oxidized to trivalent to maintain charge balance for samples C1–C4. As displayed in Fig. 4, the percentage of Ni3+is 44.2, 45.8, 47.5 and 46.5 for C1, C2, C3 and C4, respectively. The average oxidation state of the transition metal ions increased in the order of P, C1, C2, C4 and C3. Generally speaking, the lower the valence of Ni, the greater the contribution to capacity[44]. Among the layered- spinel composites, the given discharge capacities at each current density (Fig. 3d) decreased from C1, C2, C4, to C3, the same as the order of the increased Ni3+content, i.e. C1 exhibits the highest discharge capacity among composites. In addition, the lower average valence of nickel indicates that more unpaired electrons are present, which may be conducive to its increase in conductivity. Although the average oxidation state of nickel ions in sample P is lower than that in sample C1, the sample P might contain small amount of inert component Li2MnO3(the charge capacity of plateau at 4.5 V is only about 101.58 mAh/g, smaller than the predict value of 229.47 mA×h/g from the formula of Li1.2Mn0.52Ni0.2Co0.08O2). Therefore, Li ions in sample P are not easily embedded in the structure during the discharge process, resulting in a lower discharge capacity compared to the sample C1 at each current density.

Fig. 5. XPS spectra of samples P, C1, C2, C3 and C4

4.3 AC conductivity and activation energy of AC conduction

The electrode material for lithium ion battery is a hybrid conductor of ions and electrons. The rate performance of lithium ion battery is directly related to the conductivity of electrode material[45-47]. Fig. 6a showsacof all samples varies with fre- quency at 100 ℃. From the figure, for all samples, it is clear that the conductivity increases as the frequency increases. Sample C1 exhibits the highest conductivity, more than 100 times higher than other samples, which indicates it has an enhanced Li ion and electronic transport, resulting in its better rate performance. Then, conductivity decreased in the sequence of P, C2, C4 and C3. The conductivity of samples C2, C4 and C3 is very close. This trend is consistent with the above results of rate per- formance,sample C3 with the lowest con- ductivity shows the poorest rate performance. The conductivities of all samples at different frequencies are shown in Table 1. The Arrhenius equation is used to determine the activation energy. The behavior of all samples with a linear relationship between lnσacand the reciprocal of temperature in a temperature range from 100 to 140 ℃ is plotted as shown in Fig. 6b. As the temperature rises, a higher AC conductivity value has been observed. This occurs because of the charge carrier short range translational hopping between the localized states[48, 49]. The activation energies are 5.9952, 5.1637, 6.3635, 7.0877, 6.8598 eV for P, C1, C2, C3 and C4, respectively. Activation energy increased in the sequence of C1, P, C2, C4 and C3. Sample C1 shows the lowest activation energy. In other words, it is more easily activated during the process of charge-discharge, thus contributing to a higher capacity at each current density.

Fig. 6. (a) Variation of AC conductivity of all samples with respect to frequency, (b) Temperature dependence of lnσacfor all samples

Table 1. Conductivities and Magnetic Parameters for All Samples

* The chemical formulas are determined by ICP

4.4 Magnetism

The difference in conductivity for fives samples can be explained by the mole magnetic suscepti- bility. The mole magnetic susceptibility was measured in the temperature range from 2 to 300 K. Fig. 7 shows the temperature dependence of the mole magnetic susceptibility() of samples. The paramagnetic data at the temperature higher than 150 ℃were fitted by Curie-Weiss formula=0+/(–), where0is the temperature-independent term,is the Curie constant, andis the Curie- Weiss temperature (> 0 for ferromagnets and< 0 for antiferromagents).0could come from the contribution of free electrons (Pauli paramagnetism), excited states of the system (Van Vleck paramagne- tism) and diamagnetic species. For our sample, the main difference of0should be highly related with Pauli paramagnetism[50, 51]. The linear fitting 1/–0()1/is shown in the inset of Fig. 7, and the obtainedvalues are –34.9, 64.8, 88.4, 95.0, 96.9 K for P, C1, C2, C3, C4, respectively. It indicates that layered-spinel composites undergo a transition from antiferromagnetic to ferromagnetic after protonation and de-protonation. The fitted magnetic parameters of all samples are listed in Table 1. Based on the Curie constant, the effective magnetic momenteffcould be determined. For the sample P, theeffvalue is 2.89. Theeffvalue was reduced by protonation and de-protonation. The lower effective magnetic moment indicates the lower spin magnetic moment. XPS measurement demonstrated that in the layered-spinel composites, the valence states of Mn and Co remain unchanged at +4 and +3 after protonation and de-protonation process. Compared to sample P, the decrease of spin magnetic moment in the C1～C4 samples is related to the oxidation of some Ni2+(2g6e2) to Ni3+(2g6e1). Among the four samples after protonation and de-protonation treatment, sample C1 has the largesteffvalue, suggesting the lowest Ni3+content, which is consistent with XPS analysis. On the other hand, antimagnetic pristine material (sample P with negative) gave a zero value of0, indicating that the contribution of free electrons could be neglected. Therefore, the relative low conductivity and poor rate performance could be explained by the low concentration of free electrons. Among four treated samples with positive, i.e. showing ferromagnetic interaction, sample C1 exhibits the0value as large as 8.40 × 10-3emu/Oe×mol, 3 or 4 time higher than other three samples. This result demonstrates that sample C1 has the highest concentration of con- duction electron, which gives a reasonable explana- tion for its high conductivity and good rate per- formance.

Fig. 7. Temperature-dependent magnetic susceptibilities for all samples

4.5 Electrochemical impedance spectroscopy

Electrochemical impedance spectroscopy (EIS) is an important technique for evaluating the interfacial electrochemistry as well as the reaction kinetics in lithium ion battery materials[52, 53]. Impedance spectra were measured on the half-cells in fre- quency region from 100 kHz to 0.01 Hz to further understand the difference of rate performance for all samples. The EIS spectra of the five samples are shown in Fig. 8. As shown in Fig. 8a, the Nyquist plot of sample P comprises of two semicircles at the high and medium-to-low frequencies, and a slope in the low frequency region. However, for the layered-spinel composites C1～C4, two semicircles at the high and medium frequencies are superim- posed together, making it difficult to distinguish. Therefore, the different equivalent circuits are adopted for P and composites C1～C4 (Fig. 8e), and the resultant fitting parameters are listed in Table 2. The sum of solid-electrolyte interface (SEI) films resistance Rsfand charge transfer resistance Rctvalues for P, C1, C2, C3 and C4 are 135.7, 106.0, 200.0, 244.8 and 202.2 Ω, respectively. The smaller value of Rsf+ Rctfor C1 is positive for the transfer of electron and Li+during the lithium insertion and extraction reactions, resulting in a better rate performance. This result is consistent well with its high concentration of conduction electron as con- firmed by the results of AC and magnetism.

Fig. 8. (a, b, c, d, f ) EIS plots of samples P, C1, C2, C3 and C4 charged to 4.3 V during the 5th cycle, (e) Equivalent circuits for EIS fitting

Table2. Values of Re, Rsf, Rct, and Rtotal for All Samples. Re Represents the Internal Ohmic Resistance, Rsf and Rct Corresponding to the Solid-electrolyte Interface (SEI) Films Resistance and Charge Transfer Resistance

5 CONCLUSION

A series of layered-spinel composites are pre- pared by the facile processes of protonation and de-protonation for pristine layered material. These treated compounds are demonstrated to be layered- spinel composites by the characterizations of XRD, Raman and oxidation-reduction potentials in charge-discharge curves. Compared to pristine sample, the treated samples exhibited higher initial discharge capacity at current density of 200 mA/g. Moreover, sample C1 dealt with 0.1 M nitric acid gives excellent rate performance. In addition to the 3D channel related to the spinel structure in the integrated material, we have further revealed that C1 sample with low average valence state (44.2% of Ni3+), high conductivity (1.124 × 10-2Ω-1×m-1at 39800 Hz) and ambient temperature magnetic susceptibility (8.40 × 10-3emu/Oe×mol) has high concentration of conduction electron by the analysis of XPS, AC conductivity and magnetism data. The high concentration of conduction electron for sample C1 can bring about an enhanced velocity of lithium-ion diffusion and electronic transport and a decreased total resistance, which results in its better rate performance as well as higher specific capacity.

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9 March 2018;

28 June 2018

① This work was financially supported by NSFC (No. 21571176, 21611530688, 21771171, 21671077 and 21025104)

.E-mail: lipingli@jlu.edu.cn

10.14102/j.cnki.0254-5861.2011-2001