Effect of Ni substitution on structural stability, micromorphology, and electrochemical performance of Li₂MnSiO₄/C cathode materials

Abstract

The chief drawback of Li₂MnSiO₄ cathode materials is structural instability deriving from the transformation of [MnO₄] tetrahedra to [MnO₆] octahedra during oxidation from Mn(II) to Mn(III) and the Jahn–Teller effect especially from Mn(III) to Mn(IV). Based on the theory that Ni²⁺ could keep the [NiO₄] ligand unchanged and stable during oxidation/reduction, [NiO₄] tetrahedra may behave as a regional structural framework to support the total crystal lattice and alleviate structural distortion partially. In this paper we synthesize nano-Li₂Mn_1−xNi_xSiO₄/C samples with a small amount of Ni as the dopant via a hydrothermal route. Firstly, we discuss the existential state of Ni through XPS spectra, and demonstrate that both Ni²⁺ and metallic Ni coexist, especially Ni²⁺ could improve structural stability. The influence of Ni doping on morphology, structural stability, and electrochemical performance is discussed through Raman spectroscopy and electrochemical tests below. It is confirmed that there exists an optimal Ni content under which the optimal performance is displayed. In our work, the optimum Ni content is 5.0%, and the maximal discharge capacity achieved is 235 mA h g⁻¹ for the 5.0% Ni doped sample in the first cycle.

---

1. Introduction

The urgency of the energy crisis and environmental pollution makes lithium-ion batteries increasingly important, especially in providing power for energy storage systems such as hybrid electric vehicles and portable electronic devices, so developing superior cathode materials is crucial in future research.¹ Compared with traditional lithium-ion intercalated transition metal oxides (e.g. LiCoO₂, LiMn₂O₄), polyanionic compounds have remarkable advantages such as high capacity, safety and thermal stability because of strong covalent bonds between nonmetallic atoms.^2–4 Among the series of polyanionic compounds (Li_xM_y(XO_n)_z, M = Fe, Mn, Co, X = P, Si), Li₂MnSiO₄ has been proposed as the most potential material because it theoretically extracts two lithium ions totally and has a reversible capacity above 330 mA h g⁻¹.^5–7 It is hard for Li₂FeSiO₄ and Li₂CoSiO₄ to extract two lithium ions,^8,9 because the oxidation voltage is predicted to be outside the voltage stability window of common electrolytes.¹⁰ Furthermore, abundant resource, low cost and non-toxic precursors are also essential factors for the possible application of Li₂MnSiO₄. Besides, Li₂MnSiO₄ exhibits four types of structural forms, including two orthorhombic phases (Pmnb, Pmn2₁) and two monoclinic phases (P2₁/n, Pn), separately related to either β or γ Li₃PO₄ polymorphs.^11–14 The current research focuses more on orthorhombic Li₂MnSiO₄ polymorphs because of their higher potential and commercial value, and the object of our study is the Pmn2₁ polymorph.

Despite the promising prospect in theory, a lot of theoretical and technical problems have not been resolved still such as, poor cycling stability, reversibility and actual low discharge capacity etc. The most crucial drawback of this material is its structural instability during charge/discharge process.^11–16 During the oxidation of Mn²⁺ to Mn³⁺, the structural collapse is mainly derived from the transformation of [MnO₄] tetrahedron into [MnO₆] octahedron. Furthermore the dominant variation during the next oxidation of Mn³⁺ to Mn⁴⁺ is the Jahn–Teller effect, under which the [MnO₆] octahedron would be lengthened along the center axis, causing more obvious volume distortion and worse symmetry, even though the coordination number remains unchanged. That's why the material probably turns into amorphous state because of collapsing structure, the lithium-ion can't be extracted totally and the actual capacity descends promptly in initial cycles. Another dominating drawback of this material is the extremely low electronic conductivity (only 10⁻¹⁶ S cm⁻¹ at room temperature) since [MnO₄] tetrahedrons are surrounded by insulating [SiO₄] tetrahedrons. So carbon coating is a usual method to improve electronic conductivity.^17–20 In addition, the polarization in charge/discharge process, Li–Mn anti-site defect and dissolution in electrolyte are also negative factors to influence the performance of Li₂MnSiO₄.

In most cases, the modification has been completed through cationic doping based on the theory of inhibiting the Jahn–Teller effect via adjusting electronic distribution of transition metal ionic d-orbitals and optimizing the geometry configuration, or increasing the lithium-ion content via defect reactions so as to raise initial discharge capacity.^21–27 However the intrinsic problem of structural collapse and drastic capacity fading still exists, which is closely related with the reversibility and capacity retention, so it's key to raise the capacity retention rather than increase the discharge capacity of initial cycles, because the actual discharge capacity of modified Li₂MnSiO₄ in first few cycles has already been higher than some commercial cathodes. Therefore searching an approach to stabilize the preliminary high capacity is important in prospective research.

Based on crystal field theory, some researchers try to introduce some kind of cation that could keep ligand unchanged and stable during oxidation, these stable ligands as structural framework to support the crystal lattice and prevent or ameliorate the regional structural degradation. Based on this, some research tried to use ferric salt or nickel salt as precursors. Ruiyong Chen et al. synthesized Li₂Fe_1−yMn_ySiO₄ nanocrystalline and the results showed better retention than pure Li₂MnSiO₄ but still poor reversibility to date.²⁸ However, the Fe²⁺ dopant also leads to theoretical capacity reduction because the redox reaction (Fe³⁺/Fe⁴⁺) is impractical.^29–31 While nickel salt is relatively better choice because the redox reaction (Ni²⁺/Ni⁴⁺) can take place theoretically, and substituting Ni²⁺ having close radius with Mn²⁺ would not cause the crystalline distorted excessively. The DFT calculations in A. Saracibar's research also suggest that Mn substituted by Ni might help to prevent the structural collapse because when Li₂MnSiO₄ is doped with transition metal ion (Mg²⁺, Fe²⁺, Co²⁺, Ni²⁺), Ni²⁺ is investigated as the only one which does not produce octahedral Mn⁴⁺ in any delithiated Li_1−yMn_1−xNi_xSiO₄ configuration.³² But the current study about Ni²⁺ doping is rare and the previous results show it's not easy to prepare Ni²⁺ doping sample because the reducing conditions of the carbon coating process (high temperature heating in Ar) easily form Ni nanoparticles.^32,33 In this work, we choose Ni²⁺ as dopant and synthesize the samples with different content of Ni element. Not only the existential valence state of Ni element has been firstly discussed, the influence of Ni element on the crystalline, and whether the expected stability effect of Ni²⁺ really works in structural transformation and capacity degradation has been also discussed.

2. Experimental

2.1 Synthesis of Li₂MnSiO₄/C nanoparticles

Samples of Li₂MnSiO₄/C nanoparticles were synthesized through a hydrothermal route. Tetraethyl orthosilicate (TEOS) (Aladdin, 99%) and Mn(Ac)₂·H₂O (Aladdin, 99%) were mixed in solvents of ethanol, distilled water and NH₃·H₂O (volume ratio of 8 : 2 : 1) for several hours in order to be fully hydrolyzed. Then the admixture was kept at 150 °C for 24 h in Teflon-lined stainless steel autoclave and the intermediate compounds were mixed homogeneously with Li₂CO₃ (Aladdin, AR) through the ball milling method. After the mixture was calcined at 700 °C for 10 h under Ar/H₂ (95/5 vol%) atmosphere in tube furnace, the product obtained as pristine Li₂MnSiO₄.

Starch ((C₆H₁₀O₅)_n, Aladdin, AR) as carbon source was divided into two parts to add. The first one was homogeneously mixed with admixture and stirred for several hours before hydrothermal reaction. Because a small quantity of carbon added during hydrothermal process can inhibit grain growth, particle agglomeration and accelerating reaction uniformity. Then the second part of starch was heated in about 70 °C water bath for several hours to be fully gelatinized, through which the starch would swell to many times of common volume, so that the gelatinized starch could coat on the surface of particles more homogeneously.

For discussing the influence of different Ni content, Ni(Ac)₂·H₂O (Aladdin, 99%) was added in calculated quantity to offer corresponding partial Ni²⁺. As a comparison, Li₂Mn_1−xNi_xSiO₄/C (x = 0, 2.5%, 5.0%, 7.5%) composites were synthesized in our experiment. They are named respectively as pristine Li₂MnSiO₄/C, 2.5% Ni doped sample, 5.0% Ni doped sample and 7.5% Ni doped sample.

2.2 Structural characterizations

X-ray diffraction patterns of the synthesized powders were collected by X-ray diffraction (XRD, Rint-2000V/PC, Rigaku, Japan) using Cu-Kα radiation (l = 0.154 nm) and θ/2θ Bragg Brentano geometry. The data were recorded in the range of 10–60° at a scanning rate of 5° per min. The particle morphology was observed with scanning electron microscope (SEM, Carl Zeiss, ZEISS SUPRA®55) and high resolution transmission electron microscope (HRTEM, JEOL-2011). Laser Raman Tester (HORIBA, LabRam HR800) was used to reflect the structural changes of active substance during charge/discharge process, and X-ray photoelectron spectra (Thermo Fisher, ESCALAB 250X) was applied to measure existential state of elements. The carbon content of the products was tested by thermal gravimetric-differential scanning calorimetry analysis (NETZSCH STA 449F3 thermal analyzer) in air from room temperature to 900 °C at a heating rate of 10 °C min⁻¹.

2.3 Cells assembly and electrochemical tests

To fabricate the cathode of a cell, the 80 wt% active material need to be mixed with 10 wt% carbon black and 10 wt% polyvinylidene fluoride (PVDF) dispersed in N-methylpyrrolidone (NMP) and stirred for several hours to obtain the slurry. The slurry was coated onto aluminum foil and then dried at 120 °C for 10 h in a vacuum oven. The coated foil was about 120 μm thick and cut into circular discs, 15 mm in diameter. The lithium foil was used as anode, and 1 M LiPF₆ in ethylene carbonate (EC)–dimethyl carbonate (DMC) mixture (volume ratio of 1 : 1) was used as the electrolyte solution, and Celgard 2300 polypropylene film as the separator. All of them were assembled as coin cells in an argon-filled glove box (H₂O, O₂ ≤ 0.5 ppm).

The charge/discharge cycles were carried out using a LAND CT2001A tester at a rate of 0.05C in the range of 1.5–4.8 V at 25 °C. And the electrochemical impedance spectroscopy was examined by applying an AC voltage of 5 mV over the frequency range of 0.01 Hz to 100 kHz in the same workstation.

3. Results and discussion

3.1 Structural analysis

In this work, Pmn2₁ polymorph is prepared, and Ni²⁺ substitute for partial Mn²⁺ in modified materials. The possible crystal structures of the prepared Li₂MnSiO₄ and Li₂Mn_1−xNi_xSiO₄ are as shown in Fig. 1, wherein Ni²⁺ occupies and replaces Mn²⁺ partially, forming [NiO₄] tetrahedrons, which substitute for a few of [MnO₄] tetrahedrons.

Fig. 1 The crystal structures (a) Li₂MnSiO₄ (b) Li₂Mn_1−xNi_xSiO₄.

The fact that phase structure of Li₂Mn_1−xNi_xSiO₄ keeps unchanged is proved by XRD patterns, which of the pristine and modified Li₂MnSiO₄/C samples are as shown in Fig. 2. The peaks are all in good agreement with orthorhombic phase of Pmn2₁ space group, and the indices of typical crystal plane have been marked.^34,35 It means Ni element doping has no influence on main crystalline structure, but the diffraction peaks of modified samples become broader, the possible explanation is lower crystallinity of the sample, or smaller crystallite size, and the latter has been proved by the corresponding SEM and TEM figures. It's inevitable to form some impurities such as MnO and Mn₂SiO₄ in as-prepared samples via hydrothermal route.¹¹ As the increase of Ni element content, the relative impurities decrease, so we assume Ni²⁺ doping is favorable in inhibiting the formation of impurities, but the new MnO peak emerges in 7.5% Ni doped sample. Referring to the standard Pmn2₁-Li₂MnSiO₄ XRD pattern, the lattice parameters are determined by GASA refinement, as listed in Table 1. The substitution of Ni element leads to the continuous increasing of the lattice constant b and c, which kind of variation probably derives from larger radius of doped Ni²⁺(0.69 a) compared to Mn²⁺(0.67 Å).

Fig. 2 The XRD patterns of the pristine Li₂MnSiO₄/C and modified samples.

Table 1 Lattice parameters of as-prepared Li₂Mn_1−xNi_xSiO₄ samples

x	a (Å)	b (Å)	c (Å)	Rwp (%)	Rp (%)
0	6.3000(1)	5.3737(5)	4.9697(5)	12.8	10.2
2.5%	6.2994(1)	5.3798(9)	4.9721(5)	11.0	8.6
5.0%	6.2980(2)	5.3798(7)	4.9734(3)	12.6	9.3
7.5%	6.2965(7)	5.3864(4)	4.9751(6)	14.3	10.1

---

Besides, the right part of Fig. 2 is enlarged part of XRD pattern in the range of 43° to 53°, it proves the existence of metallic Ni unavoidably, the peaks of which are easily covered up by the peaks of (220) and (122), and the precipitation as Ni nanoparticles could explain the increase of MnO in high Ni element content doped sample.^32,33 To exclude the potential reducing effect of trace H₂ in sintering atmosphere, the high pure N₂ is applied, which still causes the production of metallic Ni. The reduction conditions are most likely to be derived from the carbonization of the cladding layer. We preliminarily conclude the coexistence of Ni²⁺ and metallic Ni, which is also verified by the following tests. Finally no peaks corresponding to carbon have been observed here, which indicates the embedding carbon is amorphous.

3.2 Existential state of Ni atom

To fully certify the existence of bivalent nickel ion, the X-ray photoelectron spectroscopy is applied to analyze the valence state of nickel element. Fig. 3(a) and (b) display the atom hybridization state of Mn and Si in pure sample and modified samples separately, where the binding energy of Mn 2p_3/2 (641.7 eV) is corresponding to Mn²⁺, showing that the divalent state of manganese in our samples.^17,36 And it is almost unchanged with Ni element substitution content, and accords well with past results.³² According to the basically conformity of fitting consequences, we can conclude that the doping of Ni element has no much effect on hybridization state of manganese element and silicon element. Fig. 3(c) illustrates that the atom hybridization state of nickel ion is 2p orbital hybridization, and it has three peaks, with the binding energy in 854.5 eV, 860.4 eV and 879.1 eV respectively through accurate fitting. A. Saracibar's work shows Ni 2p spectra do not change at all with charge voltage most likely due to its precipitation as metallic Ni nanoparticles, which only exhibits peak in 854.5 eV.³² This peak is in good agreement with our results, and the extra 879.1 eV and 860.4 eV peaks correspond to different hybridization state of divalent nickel ion, which verify the coexistence of Ni²⁺ and metallic Ni. And along with the increase of Ni element content, the ratio of two forms is almost same. Above all, the existence of Ni²⁺ and the redox of Ni²⁺ have been proved by XPS.

Fig. 3 Ex situ XPS patterns of the samples in pristine state (a) The Mn 2p spectra (b) The Si 2p spectra (c) The Ni 2p spectra in modified materials.

3.3 Effect of grain refinement

Except for doping and carbon coating, nanocrystallization of grains is also a main task of material modification, the most remarkable advantage of which is increasing contact area between active materials and conductive material, in order to shorten the diffusion distance of lithium ion and further facilitate transport of lithium-ion via those densely packed nano-sized grains. Besides the conductive carbon coating connects the nanoparticles in close crystal domains, providing highly conductive channels for electron mobility between neighboring Li₂MnSiO₄ nanoparticles.²³ So the electrochemical performance of samples with nanoparticles micromorphology would be improved. In this paper, we focus on the grain refinement effect, the possible explanation of which is the Ni element doping.

Fig. 4(a)–(d) are the SEM images of the pristine Li₂MnSiO₄/C, Li₂Mn_97.5%Ni_2.5%SiO₄/C, Li₂Mn_95.0%Ni_5.0%SiO₄/C and Li₂Mn_92.5%Ni_7.5%SiO₄/C respectively, where the grain refinement phenomenon emerges apparently. As shown in Fig. 4, Ni element substitution is effective in reducing the grain size, and grain size is homogeneous, no other abundant impurities are observed. The fact is much more obvious in TEM images. Fig. 4(e) is the TEM image of pristine Li₂MnSiO₄/C, Fig. 4(f)–(h) correspond to the TEM images of the modified sample with 2.5%, 5.0% and 7.5% Ni element substituted. The particles in Fig. 4(e) are inhomogeneous, the size range covers from about 30 nm to over 100 nm and the carbon layer is thick. However Fig. 4(f) distinctly exhibits the homogeneity of arrayed particles in visual field, the average size of which is estimated about 20 nm, and the carbon layer is thinner and more uniform. The possible explanation of grain refinement effect is the decrease of surface energy, so as to prevent grain growth because of doping.³⁷ The similar morphology appears in Fig. 4(g) and (h), but particles in Fig. 4(h) occur agglomeration.

Fig. 4 The SEM images of the nanocomposites (a) Li₂MnSiO₄/C (b) Li₂Mn_97.5%Ni_2.5%SiO₄/C (c) Li₂Mn_95.0%Ni_5.0%SiO₄/C (d) Li₂Mn_92.5%Ni_7.5%SiO₄/C and the TEM images of the materials (e) Li₂MnSiO₄/C and (f) Li₂Mn_97.5%Ni_2.5%SiO₄ (the insert is enlarged figure of one particle), (g) Li₂Mn_95.0%Ni_5.0%SiO₄ and (h) Li₂Mn_92.5%Ni_7.5%SiO₄.

Besides, nanoparticles with less than 20 nm in diameter would provide shorter pathways for quick lithium-ion and electronic conduction within the nanoparticles.³⁸ The inserted picture in Fig. 4(f) also suggests the crystal fringes and coating carbon layer of an intact grain, the fringe spacing corresponding to (210) plane. The improvement in transport of lithium-ion and electric conductivity originated from the grain refinement may be one of the most important factors to influence the electrochemical performance.

3.4 Improvement in structural stabilization

Raman tester is used to characterize structural stability of active substance to a certain degree, which reflects the capacity retention during charge/discharge process. According to the previous reports, the peak around 615 cm⁻¹ is assigned to the symmetric Mn–O stretching vibration of [MnO₄] group and the strong band around 635 cm⁻¹ is assigned to the symmetric Mn–O stretching mode of [MnO₆] octahedron (A_1g).^39–42 When the oxidation reaction occurs, the characteristic peak of Mn–O bonds in laser Raman spectrograms would shift right. The previous results show even only after first cycle, the Raman shift has shifted right obviously from 619 cm⁻¹ to 650 cm⁻¹, this transformation can reflect the rapid variation of [MnO_x] ligand in Li₂MnSiO₄.⁴⁰ On the other hand, the peak around 490 cm⁻¹ is associated with the Ni²⁺–O stretching mode in the structure and the peak located in 370 cm⁻¹ could be assigned to the first-order transverse optical mode of Ni²⁺–O.^43,44 And no Raman peak corresponding to pure metallic Ni would appear in figures.

In this work, we test the Raman spectrograms of a series of samples doped with different Ni element content in the range of 100–2000 cm⁻¹, as shown in Fig. 5(a). Along with the increase of Ni element content, the characteristic peak of [MnO₄] ligand keeps unchanged basically, but the intensity of Ni²⁺–O peaks which are detected in 370 cm⁻¹ and 490 cm⁻¹ gradually increases, representing that the quantity of [NiO₄] tetrahedrons and Ni²⁺–O bonds has kept rising. Complementally, the peaks of D band and G band in diagrams reveal the existing forms of C–C bonds, which composed of both disordered carbon and ordered carbon.^29,38

Fig. 5 The Raman spectra of (a) as-prepared Li₂Mn_1−xNi_xSiO₄/C samples with different Ni content and samples before 1st cycle and after 1st cycle (b) pristine Li₂MnSiO₄/C (c) Li₂Mn_97.5%Ni_2.5%SiO₄/C (d) Li₂Mn_95.0%Ni_5.0%SiO₄/C (e) Li₂Mn_97.5%Ni_2.5%SiO₄/C (the inserted figures correspond to enlarged parts of [MnO_x] transformation).

To analyze the degree of structural distortion according to the Raman results. The variation of Mn–O bonds of the pristine sample and the Ni-doped samples were tested before cycling and after first cycle. Fig. 5(b) is the spectra of the pristine Li₂MnSiO₄/C, and Fig. 5(c)–(e) are the spectra of the Ni-doped samples. Refer to the immobile peaks of C–C bonds, the right shift of Mn–O bonds from 615 cm⁻¹ to 642 cm⁻¹ is obvious. In comparison, the Raman shift of Mn–O bonds in the 2.5% Ni doped sample only occurs in the range of 615 cm⁻¹ to 625 cm⁻¹ after first cycle, besides 620 cm⁻¹ for 5.0% Ni doped sample and 626 cm⁻¹ for 7.5% Ni doped sample respectively, which reflects smaller part of [MnO₄] tetrahedrons in modified samples falling into structural degradation, and the distortion extent of them is much lower than the structure of pristine sample. Especially the 5.0% Ni doped sample also displays lowest distortion extent. As presumed, Ni substitution does effectively lower the structural distortion extent, most likely because of Ni²⁺ acting as a pillar to prevent lattice collapsing. And for Ni–O bond, the peak position always keeps unchanged in 490 cm⁻¹, which means the coordination number of Ni²⁺ is unchanged.

3.5 Analysis of electrochemical performance

Previous DFT results predict that Li₂NiSiO₄ has high voltages platform, 4.5 V for the Ni²⁺/Ni³⁺ couple and 5.2 V for the Ni³⁺/Ni⁴⁺ couple.³² We can thus assume that Ni element substitution would increase the average voltage platform of Li₂MnSiO₄. Fig. 6(a) displays the galvanostatic charge/discharge curves of first cycle of pristine Li₂MnSiO₄/C and Li₂Mn_1−xNi_xSiO₄/C nanocomposites, and discharge platforms increase for the modified samples as predicted. Thus the modified materials reveal improved capacity and higher discharge voltage, especially 5.0% Ni doped sample owns highest capacity and voltage platform. On the other hand the reduced capacity in 7.5% Ni doped sample probably results from the influence of metallic Ni.

Fig. 6 (a) The charge and discharge curves of the as-prepared Li₂Mn_1−xNi_xSiO₄/C nanocomposites (b) the midpoint potential curves of the as-prepared Li₂Mn_1−xNi_xSiO₄/C nanocomposites (cycled at room temperature, the rate is 0.05C, the voltage range is 1.5–4.8 V).

The discharge midpoint potential curves of samples are obtained in Fig. 6(b), the same tendency of midpoint voltage accords with Fig. 6(a) to some extent. Moreover, it also conforms to the theoretical trend of increasing average voltage, suggesting a new redox reaction with higher oxidation voltage exists. We infer that partial Ni²⁺ enters into crystal lattice and is oxidized indeed, as mentioned above. However, along with the much higher Ni element content, a peak value of potential appears in 5.0% Ni doped sample. The possible reason is the limitation of Ni²⁺ substitution and oxidation in lattice, that's why 5.0% Ni doped sample displays best performance and there is no obvious midpoint potential difference between 5.0% doped sample and 7.5% doped sample as shown in Fig. 6(b). We think there is an optimal balance quantity between metallic Ni/Ni²⁺ that would has a best improvement effect on electrochemical performance, but once the metallic Ni content is over some limit, the performance may deteriorate instead.

Fig. 7(a) exhibits the cycle curves of Li₂MnSiO₄/C and Li₂Mn_1−xNi_xSiO₄/C nanocomposites at room temperature, which reflects the discharge capacity results and the electrochemical retention performance. The raised capacity of modified samples is consistent with Fig. 6(a) especially in preliminary cycles, and the maximal discharge capacity in first cycle could achieve above 235 mA h g⁻¹ within 5.0% Ni doped sample. Although the capacity degradation phenomenon still exists, the descending trend has been weakened. The raised capacity mainly results from the stabilization of Ni²⁺ and the reduced resistivity caused by metallic Ni.

Fig. 7 (a) The discharge capacity curves of the as-prepared Li₂Mn_1−xNi_xSiO₄/C nanocomposites (cycled at room temperature, the rate is 0.05C, the voltage range is 1.5–4.8 V). And the calculated capacity retention is as shown in (b).

Fig. 7(b) displays the specific calculated capacity retention results based on the second discharge capacity as the initial capacity because of the possible inaccuracy of first cycle. It's noticeable the Ni element doped materials also show improved retention in initial several cycles. The curves show similar superiority among modified samples more visually, but the superiority disappears after 30 cycles. Associated with the Raman spectroscopy, the alleviated fading capacity exactly accords with the lower distortion extent of structure. Therefore substituted Ni²⁺ not only promotes the initial discharge capacity but also alleviates the capacity decay in preliminary stage of cycles via acting as a pillar to prevent regional lattice collapsing. Besides, among the samples with different Ni element content, the similar improvement effect in retention may imply that no matter the Ni element dopant amount, the effective Ni²⁺ as [NiO₄] is finite, as suggested from Fig. 6(b). But the promotion is still weak and transient, so further studies on improving the capacity retention are needed.

The cycling stability of samples in high charge/discharge rate 1C are shown in Fig. 8(a). In higher charge/discharge rate, the discharge capacity of samples has all been lower than those in Fig. 7(a) as a whole, and especially in third cycles all of them all descends promptly to a stable status, obviously which is caused by faster charge/discharge speed. Even though the rapid degradation, the same tendency still exists among the materials. The samples doped with Ni exhibit better performances than pristine sample, and 5.0% Ni doped sample shows optimal capacity and retention. That means no matter in low rate or high rate, the effect of Ni element on electrochemical performance both exists, and are similar on whole tendency. We also provide the rate capability of material as shown in Fig. 8(b). Along with the increase of rate, the discharge capacity descends promptly, and when the rate goes back into 0.05C, the capacity goes up relatively, but still indicating poor rate performance of this material.

Fig. 8 (a) The discharge capacity curves of the as-prepared Li₂Mn_1−xNi_xSiO₄/C nanocomposites (cycled at room temperature, the rate is 1C, the voltage range is 1.5–4.8 V). (b) The rate capability of pristine Li₂MnSiO₄/C nanocomposite.

To investigate the influence of Ni element substitution on the charge transfer resistance and lithium-ion diffusion coefficient, the samples are analyzed by AC impedance spectroscopy. Fig. 9(a) shows the EIS measurements of the materials at half of discharge state, and matched equivalent circuit is inserted. All the samples compose of a semi-circle in the high frequency and an inclined line in the low frequency. The semicircle can be ascribed to the combined action of charge transfer resistance (R_ct) and the complex impedance of a constant phase element (CPE), while the straight line is related to the lithium-ion diffusion.⁴⁵ The modified samples reveal much smaller charge transfer resistance than pristine Li₂MnSiO₄/C, which may behave as critical factor in promoting the electrochemical performance. Besides, among the modified samples the 5.0% Ni doped sample reveals lowest R_ct value. The R_ct is related to electronic conductivity and ion diffusion coefficient, so the possible explanation of the remarkable promotion is the synergistic effect of Ni and Ni²⁺ dopant, and there is optimal equilibrium state when 5.0% Ni doped.

Fig. 9 (a) EIS measurements of as-prepared Li₂Mn_1−xNi_xSiO₄/C samples at half of discharge state. Matched equivalent circuit is inserted. (b) The fitting results of the latter linear part of (a). The calculated data is recorded in Table 1.

The diffusion coefficient results of the lithium-ion (D_Li⁺) can be calculated from the EIS curves of the straight line region according to the following equation.^45,46

(R is the gas constant, T is the absolute temperature, A is the area of electrode surface, n is the theoretical number of electrons per molecule, F is the Faraday's constant, C is the concentration of lithium-ion and σ_W is the Warburg coefficient which is slope parameter of the curves in Fig. 9(b)).

Table 2 shows the impedance results, Warburg coefficients and calculated diffusion coefficients of lithium-ion after fitting. The results display reduced charge transfer resistance and increased lithium-ion diffusion coefficient of modified samples compared to pristine sample. And the peak value also appears in 5.0% Ni doped sample, which is in good agreement with the test results above.

Table 2 Parameters of fitting impedance results and calculated diffusion coefficient of lithium-ion

x in Li2Mn1−xNixSiO4/C	Rct (Ω)	σW (Ω cm^2 s^−0.5)	DLi^+ (cm^2 s^−1)
0	573.0	162.5	2.94 × 10^−16
2.5%	177.3	86.29	1.05 × 10^−15
5.0%	59.88	30.04	8.64 × 10^−15
7.5%	92.92	48.89	3.26 × 10^−15

---

Summing up the above discussion, the role of Ni doping is observed, and the optimal effect is all displayed for 5.0% Ni doped sample. We suggest that no matter Ni²⁺ or metallic Ni, they both influence the materials performance. Firstly the leading effect of Ni²⁺ is substituting the Mn²⁺ site, stabilizing the regional crystal lattice, reducing the total structural distortion extent via keeping stable ligand, and promoting voltage platform, alleviating capacity degradation. To explain the specific improvement effect of Ni²⁺ more visually, the models in Fig. 10 are speculated to reveal the transformation of Mn–O ligand and Ni–O ligand in oxidation reaction of pristine sample and modified samples respectively. The variation of Mn–O ligand and the Jahn–Teller effect causes structural degradation and capacity fading. The most possible oxidation process has been simulated in Fig. 10(b), where the ligand number keeps unchanged and the distortion caused by the Jahn–Teller effect is weak, so that the doped [NiO₄] can sustain the whole crystal and prevent the crystal structure from collapsing so as to improve the capacity retention.

Fig. 10 The oxidation process of metal ligand (a) Mn–O oxide (b) Ni–O oxide.

On the other hand, the metallic Ni works as conductive additive, and thus improves the poor intrinsic conductivity. However, the formation of metallic Ni would result in producing Ni vacancy, which probably leads to MnO impurity and theoretical capacity fading, meanwhile the deficiency of [NiO₄] ligand would lead to regional structural distortion, thus blocking the transport of lithium-ion. In a conclusion, we think there is an optimal quantity between the two forms of Ni element to achieve best electrochemical performance, and in our work the 5.0% Ni doped sample displays optimal performance.

4. Conclusions

The concept that substitution with equivalent metal ions which can maintain unchanged ligand during oxidation/reduction process theoretically, working as structural pillars to improve electrochemical performance via alleviating the structural degradation, and then inhibiting the Jahn–Teller effect in structural variation model of Li₂MnSiO₄ material has been firstly hypothesized in our work. Besides, the Pmn2₁ samples with different Ni element content are successfully synthesized, and the existential state of Ni atom in crystalline has also been firstly discussed via XPS spectra, which proves to be coexisting metal and divalent ions.

The Raman results demonstrate that transformation of [MnO₄] tetrahedron to [MnO₆] octahedron has been decelerated in modified samples compared with pristine Li₂MnSiO₄/C material, corresponding to the improved retention in electrochemical tests especially in initial several cycles. Moreover, the Ni element substitution is also effective in preparing pure phase, refining crystalline grain, promoting discharge capacity, reducing electrical resistivity and improving diffusion coefficient of lithium-ion, reflected by XRD patterns, microscope pictures and electrochemical tests.

As a conclusion, we think there are two kinds of modification mechanism with Ni element substitution at the same time. Firstly the leading effect of Ni²⁺ is substituting the Mn²⁺ site, stabilizing the regional crystal lattice, reducing the total structural distortion extent via keeping stable ligand, and thus promoting voltage platform, alleviating capacity degradation, promoting initial discharge capacity to improve electrochemical performance. The effective Ni²⁺ is also inferred to be limited in oxidation reactions through electrochemical results. On the other hand, although pure metallic Ni will help to reduce electrical resistivity, it still leads to formation of impurity and instability of regional structure because of deficiency of [MO₄] ligand. Therefore the electrochemical performance deteriorates in high Ni element content instead. So we think there is an optimal Ni element content, under which the modified materials displays best performance. In our work, the 5.0% Ni doped sample shows optimal performance, since there is a best balance between metallic Ni and Ni²⁺ in 5.0% Ni doped sample.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (No. 51372136) and Shenzhen Basic Research Project (No. JCYJ20130402145002372).

References

1. V. Etacheri, R. Marom, R. Elazari, G. Salitra and D. Aurbach, Energy Environ. Sci., 2011, 4, 3243.
2. D. M. Kempaiah, D. Rangappa and I. Honma, Chem. Commun., 2012, 48, 2698–2700.
3. D. Rangappa, K. D. Murukanahally, T. Tomai, A. Unemoto and I. Honma, Nano Lett., 2012, 12, 1146–1151.
4. R. J. Gummow, N. Sharma, V. K. Peterson and Y. He, J. Power Sources, 2012, 197, 231–237.
5. R. Dominko, J. Power Sources, 2008, 184, 462–468.
6. M. K. Devaraju, T. Tomai, A. Unemoto and I. Honma, RSC Adv., 2013, 3, 608–615.
7. H. N. Girish and G. Q. Shao, RSC Adv., 2015, 5, 98666–98686.
8. M. K. Devaraju, Q. D. Truong and I. Honma, RSC Adv., 2013, 3, 20633.
9. W. Chen, M. Lan, D. Zhu, C. Wang, S. Xue, C. Yang, Z. Li, J. Zhang and L. Mi, RSC Adv., 2013, 3, 408–412.
10. M. E. Arroyo-de Dompablo, M. Armand, J. M. Tarascon and U. Amador, Electrochem. Commun., 2006, 8, 1292–1298.
11. R. J. Gummow and Y. He, J. Power Sources, 2014, 253, 315–331.
12. C. A. J. Fisher, N. Kuganathan and M. S. Islam, J. Mater. Chem. A, 2013, 1, 4207.
13. M. E. Arroyo-de Dompablo, R. Dominko, J. M. Gallardo-Amores, L. Dupont, G. Mali, H. Ehrenberg, J. Jamnik and E. Moran, Chem. Mater., 2008, 20, 5574–5584.
14. G. Mali, A. Meden and R. Dominko, Chem. Commun., 2010, 46, 3306–3308.
15. R. Dominko, M. Bele, A. Kokalj, M. Gaberscek and J. Jamnik, J. Power Sources, 2007, 174, 457–461.
16. V. V. Politaev, A. A. Petrenko, V. B. Nalbandyan, B. S. Medvedev and E. S. Shvetsova, J. Solid State Chem., 2007, 180, 1045–1050.
17. Y.-X. Li, Z.-L. Gong and Y. Yang, J. Power Sources, 2007, 174, 528–532.
18. V. Aravindan, K. Karthikeyan, J. W. Lee, S. Madhavi and Y. S. Lee, J. Phys. D: Appl. Phys., 2011, 44, 152001.
19. I. Belharouak, A. Abouimrane and K. Amine, J. Phys. Chem. C, 2009, 113, 20733–20737.
20. Q. Zhang, Q. Zhuang, S. Xu, X. Qiu, Y. Cui, Y. Shi and Y. Qiang, Ionics, 2011, 18, 487–494.
21. C. Lai, Z. Wang, Q. Xu and H. Li, Mater. Lett., 2015, 139, 134–137.
22. N. Kuganathan and M. S. Islam, Chem. Mater., 2009, 21, 5196–5202.
23. S. Choi, S. J. Kim, Y. J. Yun, S. S. Lee, S.-Y. Choi and H.-K. Jung, Mater. Lett., 2013, 105, 113–116.
24. M. Wang, M. Yang, L. Ma and X. Shen, RSC Adv., 2015, 5, 1612–1618.
25. S. Zhang, Z. Lin, L. Ji, Y. Li, G. Xu, L. Xue, S. Li, Y. Lu, O. Toprakci and X. Zhang, J. Mater. Chem., 2012, 22, 14661.
26. C. Hwang, T. Kim, Y. Noh, W. Cha, J. Shim, K. Kwak, K. M. Ok and K.-K. Lee, Mater. Lett., 2016, 164, 270–273.
27. M. Zhang, S. Zhao, Q. Chen and G. Yan, RSC Adv., 2014, 4, 30876.
28. R. Chen, R. Heinzmann, S. Mangold, V. S. K. Chakravadhanula, H. Hahn and S. Indris, J. Phys. Chem. C, 2013, 117, 884–893.
29. S. Zhang, Y. Li, G. Xu, S. Li, Y. Lu, O. Toprakci and X. Zhang, J. Power Sources, 2012, 213, 10–15.
30. M. Bini, S. Ferrari, D. Capsoni, C. Spreafico, C. Tealdi and P. Mustarelli, J. Solid State Chem., 2013, 200, 70–75.
31. R. Dominko, I. Arčon, A. Kodre, D. Hanžel and M. Gaberšček, J. Power Sources, 2009, 189, 51–58.
32. A. Saracibar, Z. Wang, K. J. Carroll, Y. S. Meng and M. E. A.-d. Dompablo, J. Mater. Chem. A, 2015, 3, 6004–6011.
33. Y.-C. Wang, S.-X. Zhao, P.-Y. Zhai, F. Li and C.-W. Nan, J. Alloys Compd., 2014, 614, 271–276.
34. R. Dominko, M. Bele, M. Gaberšček, A. Meden, M. Remškar and J. Jamnik, Electrochem. Commun., 2006, 8, 217–222.
35. Y. Zhao, C. Wu, J. Li and L. Guan, J. Mater. Chem. A, 2013, 1, 3856.
36. H. Gong, Y. Zhu, L. Wang, D. Wei, J. Liang and Y. Qian, J. Power Sources, 2014, 246, 192–197.
37. Q. Liu, W. Liu, D. Li and Z. Wang, Mater. Lett., 2016, 162, 87–90.
38. T. Muraliganth, K. R. Stroukoff and A. Manthiram, Chem. Mater., 2010, 22, 5754–5761.
39. M. Kim, X. M. Chen, X. Wang, C. S. Nelson, R. Budakian, P. Abbamonte and S. L. Cooper, Phys. Rev. B: Condens. Matter Mater. Phys., 2011, 84, 174424.
40. P.-Y. Zhai, S.-X. Zhao, H.-M. Cheng, J.-W. Zhao and C.-W. Nan, Electrochim. Acta, 2015, 153, 217–224.
41. C. M. Julien and M. Massot, Mater. Sci. Eng., B, 2003, 97, 217–230.
42. S. H. Oh, K. Y. Chung, S. H. Jeon, C. S. Kim, W. I. Cho and B. W. Cho, J. Alloys Compd., 2009, 469, 244–250.
43. G. B. Zhong, Y. Y. Wang, X. J. Zhao, Q. S. Wang, Y. Yu and C. H. Chen, J. Power Sources, 2012, 216, 368–375.
44. Y. Xue, Z. Wang, F. Yu, Y. Zhang and G. Yin, J. Mater. Chem. A, 2014, 2, 4185.
45. J. Zhu, Z. Tang, H. Tang, Q. Xu and X. Zhang, J. Electroanal. Chem., 2016, 761, 37–45.
46. J. Zhang, J. Shen, T. Wang, C. Wei, Y. Ma, C. Zhu and Y. Yue, Electrochim. Acta, 2013, 111, 691–697.

---