Introduction

All electrical equipment generated by lithium-ion batteries (LIBs) require higher electrical performance to satisfy the increasing demand [1,2,3,4], which requires the applied batteries to offer better cooperation between cathodes and anodes. Although great progress has been made in anode modification [5,6,7,8], cathode materials with higher energy densities are still needed to meet this goal. Thus, new cathodes must be explored. Li-rich Mn-based cathode materials are composed of LiMO2 (where M represents Co, Ni, Mn, etc.) and Li2MnO3, whose specific capacity can exceed 250 mAh·g−1 [9,10,11] beyond LiCoO2 [12] and LiFePO4 [13, 14]. It can be considered one of the most promising materials because Li2MnO3 can activate anion redox to generate another part of electrical energy above 4.5 V. If its initial Coulomb efficiency (ICE), voltage and capacity attenuation, and poor rate performance can further increase as the next-generation cathodes based on recent studies, Li-rich Mn-based cathode materials will replace the primary cathodes [15,16,17,18].

Many methods have been proposed to solve the aforementioned disadvantages of Li-rich Mn-based cathode materials, such as surface modification and ion doping. Surface modification, such as coating with carbon [10], metal oxides [19], fluorides [20], and phosphates [21], can inhibit the transformation of the surface structure from a layered structure to a spinel structure and mitigate the side reactions. Many experiments show that coating can improve the ICE and cycle stability [10, 19,20,21,22], but an additional coating process will increase the cost, which is unfavorable for industrialization. Ion doping is a simple modification method that does not increase the extra process and cost during the synthesis process. At present, researchers have utilized many constant-value elements to modify crystal lattices, such as Na [23], Mg [24], Al [25], Fe [26], and F [27], to obtain higher ICE performance, rate performance, and cycle stability.

Al is abundant in the Earth’s crust and costs less than most metal elements. Al doping can enhance the cycle performance by stabilizing the crystal structure and inhibiting the reduction of transition metal (TM) cation [18, 28, 29]. Some researchers have reported that Al3+ doped at the Li site and TM site has achieved good rate performance and cycle stability [25, 30, 31]. For example, Nayak et al. [30] reported Al-doped Li-rich Mn-based cathode materials (Li1.2Ni0.16Mn0.56Co0.08O2). They found that doping the Al element decreased the specific capacity and significantly increased the capacity retention rate. Chen et al. [25] synthesized an Al-doped electrode material with a layered micro/nanostructure by using a porous polypropylene film as a hard template; they found that the Li1.2Ni0.18Mn0.58Al0.04O2 sample provided the highest specific capacity, reduced the charge transfer resistance, and improved the rate performance, structural stability, and lithium-ion diffusion coefficient. Although there have been reports on Al-doped Li-rich Mn-based cathode materials, the relationship between voltage attenuation and the valence state of TM cations has not been studied, and the properties of this material at high temperature and lithium-ion full cells have not been provided, which are necessary for any material before applications [32]. Therefore, the effect of Al doping must be explored.

In addition, the coating process can increase the cost for enhancing the performance of LIBs, so it is meaningful to form a coating on the material surface without additional processes. In this work, we designed a modified Pechini method to synthesize a series of Li1.2Ni0.2Mn0.6-xAlxO2 (x = 0, 0.02, 0.04, 0.06) cathode materials with an Li2CO3 coating using a one-step synthesis method, and the corresponding schematic diagram is shown in Fig. 1. In these materials, a small amount of Mn was replaced by Al. Compared with the undoped samples, the doped samples have better voltage stability, capacity stability, temperature stability, and rate performance. In addition, we have researched how Al doping affects the structural transitions and TM cation valence reductions and explored the scientific truth.

Fig. 1
figure 1

The process of synthesizing the Li1.2Ni0.2Mn0.6O2 sample

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Experimental section

Material synthesis

A series of materials was synthesized using a modified Pechini method. Stoichiometric amounts of Mn(CH3COO)2·4H2O, Ni(CH3COO)2·4H2O, and 5% excess LiCH3COO·H2O were mixed and dissolved in deionized water and subsequently dissolved in a mixed solution of citric acid (AC) and ethylene glycol (EG) (molar ratio = 1:4). The mixed solution was evaporated and dried to obtain a polymeric precursor. The undoped sample Li1.2Ni0.2Mn0.6O2 was obtained by 5 h of pre-sintering at 500 ℃ and 12 h of sintering at 950 ℃ in air, and it was named LRMA0. The synthesis process had the extra LiCH3COO·H2O and organic agents that could decompose to generate Li2O and CO2 to form an in situ Li2CO3 coating at high temperature. For the doped samples, the same stoichiometric aluminum acetate was used to replace the manganese acetate using the aforementioned synthesis steps. According to different doping amounts, Li1.2Ni0.2Mn0.58Al0.02O2, Li1.2Ni0.2Mn0.56Al0.04O2, and Li1.2Ni0.2Mn0.54Al0.06O2 were synthesized and named LRMA2, LRMA4, and LRMA6, respectively.

Material characterizations

X-ray diffraction (XRD) was utilized to obtain the crystallographic structure and phase state of the materials with Cu-Kα radiation in a scanning range of 10–80° at a scanning rate of 5°/min. The general structural analysis software (GSAS) was used for the Rietveld refinement. The microstructure of the materials was characterized using scanning electron microscopy (SEM). The crystallinity and microstructure of the samples were characterized using transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HR-TEM). Raman spectroscopic studies of pristine and cycled electrodes were performed by a Raman spectrometer equipped with a 532-nm laser. X-ray photoelectron spectroscopy (XPS) was used to analyze the distribution of elemental species and valence states on the surfaces of the samples, and the spectra of all elements were calibrated with the C1s peak at 284.8 eV.

Electrochemical measurements

The electrochemical properties of the synthesized materials were evaluated with CR-2032 coin-type half-cells. The active material, acetylene black, and polyvinylidene fluoride (PVDF) were uniformly mixed in N-methyl-2-pyrrolidone (NMP) at a weight ratio of 8:1:1 to prepare the electrode slurry. Then, the mixed slurry was uniformly smeared onto aluminum foil and dried at 110 ℃ overnight in a vacuum oven. The electrode materials were punched into small circles with a diameter of 10 mm. The CR-2032 coin-type half-cells were assembled in an argon glove box (H2O ≤ 0.01 ppm, O2 ≤ 0.01 ppm). Lithium metal was used as the counter electrode, Celgard 2400 was used as the diaphragm, and 1-M LiPF6 dissolved in a mixture of ethylene carbonate (EC) and methyl carbonate (DMC) (EC: DMC = 1:1) was used as the electrolyte. The galvanostatic charge‒discharge test of the coin-type half-cells was performed on the Land CT-2001A test system at different rates (1 C = 200 mAh·g−1), and the voltage range was 2–4.8 V vs. Li/Li+. An electrochemical workstation (Chenhua) was used to measure the AC impedance of the coin-type half-cells at 0.1 Hz–100 kHz with an AC amplitude of 5 mV. Cyclic voltammetry (CV) was performed on a multi-channel electrochemical workstation. Lithium-ion full cells were tested in 2032-type coin cells. The negative electrodes were made up of graphite (80% wt%), PVDF (10 wt%), and acetylene black (10 wt%). In these full cells, the mass ratio of the negative electrode to the positive electrode (N/P) was approximately 1.10.

Results and discussions

Figure 2a shows the XRD patterns of all samples (LRMA0, LRMA2, LRMA4, and LRMA6). The main diffraction peaks correspond to the typical α-NaFeO2 layered structure with the R-3 m space group [33, 34]. A few diffraction peaks in the range of 20–23° can be considered the characteristic peak of Li2MnO3 with the C 2/m space group [35], as described in Fig. 2b. These weak diffraction peaks were caused by the orderly arrangement of Li and Mn in the TM layer. In addition, the peaks of (006)/(102) and (018)/(110) in the XRD patterns were noticeably split, which reveals that each material had an excellent layered structure [36]. Compared with the XRD patterns of the LRMA0 sample, no other diffraction peaks were found in the LRMA2, LRMA4, and LRMA6 samples, which suggest that Al successfully substituted Mn without forming a distinct impurity phase. After a careful analysis of the XRD patterns, a slight variation was observed: the 003 and 104 peaks shifted toward a large angle with the increase in doping amount of Al, as described in Figs. 2c and d. The shift of the Al-doped samples occurred because larger Ni2+ (r = 0.69 Å) was oxidized to smaller Ni3+ (r = 0.56 Å) due to the charge compensation effect of Al3+ substituting Mn4+ [30]. The corresponding variations in lattice parameters of all samples were calculated by Rietveld refinement. The fitted images and data are listed in Figs. 2e–h and Table 1. Lattice parameters “a” and “c” of the material gradually decreased with increasing doping amount, which is consistent with the variation in Figs. 2c and d. The c/a ratios of all samples were greater than 4.9, which further indicates an excellent layered structure [37]. The ratio of the diffraction peak intensity I(003)/ I(104) is usually associated with the degree of cation mixing in cathode materials. When the ratio is greater than 1.2, the material has a low degree of cation mixing and highly ordered layered structures [38]. Obviously, the ratios of all samples were greater than 1.2, which indicates that the degree of cation mixing was negligible. These results indicate that a small amount of Al doping will not change the layered crystal structure.

Fig. 2
figure 2

A XRD patterns of LRMA0, LRMA2, LRMA4, and LRMA6 samples, b–d the corresponding enlarged XRD patterns, Rietveld refinement results of e LRMA0, f LRMA2, g LRMA4, and h LRMA4 samples

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Table 1 Rietveld refinement data for as-prepared materials
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The SEM images of all samples are shown in Figs. 3a–d. The LRMA0 sample had the smallest particle size and the clearest edges. The LRMA0 sample had a more uniform particle size distribution (200–500 nm) than the others. The particle size of the doped samples increased with increasing amounts of doped Al, and there was a vaguely light-gossamer coating on the particles in Figs. 3b and c.

Fig. 3
figure 3

SEM images of a LRMA0, b LRMA2, c LRMA4, and d LRMA6

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Figures 4a and b show the structural information of the LRMA0 sample measured by TEM and HR-TEM images. In Fig. 4a, we did not observe an obvious coating, but the HR-TEM image in Fig. 4b clearly shows a coating of approximately 1.5 nm on the surface. The reason may be that the coating layer was too thin to be found in TEM images. Figs. 4c and d show the structural information of the LRMA4 sample measured by TEM and HR-TEM images. Fig. 4c shows a coating with a thickness of approximately 2.5 nm outside the grain, which can be more clearly observed in the HR-TEM images of Fig. 4d. The coating may be the pursued Li2CO3 because the generated chemical reaction formula is Li2O + CO2 = Li2CO3, where Li2O and CO2 could come from the decomposition of lithium acetate and other organic matter at high temperature. The Li2CO3 coating can hinder the direct contact between electrode and electrolyte and act as an artificial CEI film to reduce the loss of active material due to the spontaneous formation of the CEI film during the cycling process. In addition, due to the formation of a more stable CEI film on the surface, the surface side reaction in the cycling process is significantly inhibited, which effectively improves the electrochemical performance of the material [12]. Fig. 4b shows that the interplanar spacing was 0.47 nm, which corresponded to the (003) crystal planes. Fig. 4d shows that the interplanar spacings were 0.42 nm and 0.32 nm, which corresponded to the (020) and (021) crystal planes, respectively. These two crystal planes belong to the Li2MnO3 component, and the space group of this component is C 2/m. The characteristic peak of C 2/m is clearly visible in the XRD spectrum, but it was not treated in the spectrum because of its low component. To prove that the doped element was evenly distributed, EDS mapping of the LRMA4 sample is shown in Fig. 4e. Mn, Ni, O, and Al elements were present and evenly distributed in the material.

Fig. 4
figure 4

TEM images of LRMA0 a TEM, b HR-TEM. TEM images of LRMA4 c TEM, d HR-TEM, e the element mapping images of LRMA4

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XPS was used to directly evaluate the surface chemical composition and valence change of the LRMA0 and LRMA4 samples, as shown in Fig. 5. The spectra of the LRMA0 and LRMA4 samples are displayed in Fig. 5a. Figures 5b–f show the XPS spectra of Al 2p, Ni 2p, Mn 2p, O 1 s, and Li 1 s. The Al 2p spectra of the LRMA0 and LRMA4 samples are shown in Fig. 5b. The LRMA0 sample did not have an Al 2p peak, which suggests that the sample did not contain Al. The LRMA4 sample had an Al 2p peak, which illustrates that Al3+ was successfully doped into the lattice. In Fig. 5c, the binding energies of the Ni 2p3/2 and Ni 2p1/2 peaks were 854.8 eV and 872.5 eV [39], respectively. Two fitting peaks at 854.6 and 856 eV were obtained. The peak at 856 eV was indexed to Ni2+, and the peak at 854.6 eV was indexed to Ni3+ [40, 41]. The fitting results illustrate that both LRMA0 and LRMA4 samples contained Ni2+ and Ni3+, and LRMA4 had more Ni3+. This result corresponds to the change in the XRD pattern. Fig. 5d displays the Mn 2p spectra of samples LRMA0 and LRMA4, and the two main peaks at binding energies of 654 eV and 642 eV correspond to the peaks of Mn 2p1/2 and Mn 2p3/2 in previous literature [42], where the Mn 2p3/2 peak can be divided into two peaks: the peak with a higher binding energy is Mn4+, and the peak with a lower binding energy is Mn3+. The fitting result suggests that Mn4+ had no significant change in the LRMA0 and LRMA4 samples. Fig. 5e depicts the O 1 s XPS spectra of the LRMA0 and LRMA4 samples. The O 1 s spectrum mainly included the M–O bond (where m represents a TM cation) at 229.2 eV [43, 44], and the peak at 531.2 eV can be ascribed to the C = O bond in carbonate species (Li2CO3); Yin et al. [45] reported a similar situation in which Li2CO3 coating appeared on the material surface. The result reveals an Li2CO3 coating on the surface of LRMA4 according to the design of the in situ coating synthesis. Fig. 5f shows the Li 1 s XPS spectra of the LRMA0 and LRMA4 samples, where the Li2CO3 peak appeared at 55.3 eV for LRMA0 and LRMA4. This result is consistent with those reported in the literature [12]. Thus, the samples in this work had a thin Li2CO3 coating due to the one-step synthesis method.

Fig. 5
figure 5

XPS spectra of a survey scan, b Al2p, c Ni2p, d Mn2p, e O1s, and f Li1s for LRMA0 and LRMA4 samples

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Figs. 6a and b show the first charge‒discharge curves and differential capacity vs. voltage (dQ/dV) profiles of all samples. The charge‒discharge rate was 0.1 C, and the voltage range was 2–4.8 V. The charging curves of all samples were divided into two parts: the inclined region below 4.5 V and the plateau area above 4.5 V. The inclined region was attributed to the reversible extraction of Li+ from the Li layer in LiMO2 (M = Ni, Mn) and the oxidation from Ni2+ and Ni3+ to Ni4+ corresponding to the oxidation peak before 4.5 V in the dQ/dV profiles. The plateau area is related to the activation process of the Li2MnO3 component to remove Li2O, since the sharp oxidation peak was at 4.5 V in the dQ/dV profiles [18] [10]. The dQ/dV profiles of the subsequent reduction process exhibit three reduction peaks: the reduction peak at 4.5 V corresponds to the reduction of O− to O2−; the reduction peak near 3.7 V is attributed to the reduction of Ni4+; and the peak at 3.3 V is attributed to the reduction of Mn4+ to Mn3+ [39]. With the increase in Al doping amount, the first discharge specific capacity of the material decreased. The first discharge specific capacities were approximately 268.2 mAh·g−1, 262.2 mAh·g−1, 258.6 mAh·g−1, and 253.6 mAh·g−1 for the LRMA0, LRMA2, LRMA4, and LRMA6 samples, respectively. The specific capacity decreased because Al is inert compared to Ni and Mn, which reduced the active mass. When Co is substituted by Al, similar results will be produced [46]. The ICE is found to be 78.3%, 78.9%, 79.0%, and 78.7% for LRMA0, LRMA2, LRMA4, and LRMA6, respectively. The ICE did not greatly change with the Al doping.

Fig. 6
figure 6

A First charge–discharge curves. b The corresponding dQ/dV curves at the first cycle. c Cycling stability and d voltage fading of the LRMA0, LRMA2, LRMA4, and LRMA6 samples

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To evaluate the influence of Al doping on the cyclic performance, the tests of all samples were measured at 1 C for 200 cycles after the cell was activated at 0.1 C for three cycles in a voltage range of 2–4.8 V (2–4.65 V for the 1st cycle) in Fig. 6c. The specific capacities of the LRMA0, LRMA2, LRMA4, and LRMA6 samples were approximately 188.5 mAh·g−1, 192.6 mAh·g−1, 193.1 mAh·g−1, and 190 mAh·g−1 with capacity retentions of 63.2%, 76.2%, 89.5%, and 81.1% after 200 cycles, respectively. The relationship between the discharge medium voltage (DMV) of the four cathode materials and the number of cycles is described in Fig. 6d. After 200 cycles, the DMV of LRMA0 dropped from 3.564 to 3.177 V with a DMV decay of 1.94 mV per cycle. The attenuation of the DMV for LRMA2, LRMA4, and LRMA6 decreased from 3.571 to 3.268 V, from 3.575 to 3.353 V, and from 3.567 to 3.304 V, for which the DMV decays were 1.52 mV, 1.11 mV, and 1.32 mV per cycle, respectively. Fig. 7 shows the charge‒discharge curves at 1 C. The specific capacity of LRMA4 above 3 V was 128.7 mAh·g−1, while the specific capacities of LRMA0, LRMA2, and LRMA6 were 74.4 mAh·g−1, 103.3 mAh·g−1, and 109.2 mAh·g−1, respectively. The capacity and voltage stability of Al-doped samples significantly improved. A possible reason is that the strong Al-O bond reduced the oxygen release during the cycle process [46, 47], which inhibited the reduction of the average valence state of the TM cation. The result can be affirmed by the latest research that the continuous release of oxygen reduces the average valence state of TM cations during the cycle measurements, which directly causes the voltage decay [47, 48]. In the discussion of XRD, doping Al can increase the amount of Ni3+, which can reduce the unit cell parameters. It is favorable to inhibit the free migration of TM cations to the Li layer and slow the problematic phase transition [30]. Furthermore, Li2CO3 can form a stable CEI film on the surface of the material, inhibit the side reaction on the surface, and reduce the polarization during circulation. This is beneficial to enhance the cycle stability [6]. Table 2 compares the electrochemical performances of Li-rich Mn-based cathode materials obtained in recent years [36, 46, 49,50,51,52,53,54,55]. Compared with other reported batteries, the batteries in this paper have quite excellent performance. In particular, compared with the Al-doped sample made by Fu et al., because the sample in this work had a Li2CO3 coating, our samples had better performance. In summary, doping Al can improve the cyclic stability and mitigate the DMV decay of pristine samples. According to the results, LRMA4 has the highest capacity retention rate and the lowest DMV decay.

Fig. 7
figure 7

Selected charge/discharge potential profiles of the a LRMA0, b LRMA2, c LRMA4, and d LRMA6 samples at 1 C

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Table 2 Comparison of the electrochemical performance of Li-rich Mn-based cathode materials
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Figure 8a shows the rate capability of all samples. The rate performance was tested at current densities of 0.2 C, 0.5 C, 1 C, 2 C, 5 C, and 0.2 C. Obviously, the doped samples had slightly lower specific capacities than the undoped sample at low current densities, but with increasing current densities, the specific capacity of the doped samples gradually exceeded that of the undoped sample. In particular, compared with 246.8 mAh·g−1, 211.6 mAh·g−1, 190.7 mAh·g−1, 167.6 mAh·g−1, and 136.4 mAh·g−1 for the LRMA0 sample, the LRMA4 sample showed capacities of 241.6 mAh·g−1, 217.5 mAh·g−1, 203.5 mAh·g−1, 182 mAh·g−1, and 151 mAh·g−1 at 0.2 C, 0.5 C, 1 C, 2 C, and 5 C, respectively. After the high current density test, the capacity of all samples was well recovered to 0.2 C, and the specific capacity of the LRMA4 sample was recovered to 235.4 mAh·g−1, which is higher than that of the LRMA0 sample (232.1 mAh·g−1), which illustrates that the LRMA4 sample had better structural stability and capacity retention rate after the rapid Li+ insertion/extraction process. Al doping slightly improves the rate performance of the samples because the lithium-ion diffusion coefficient increases with increasing Al amount, as shown in Figs. 8b–d.

Fig. 8
figure 8

a Rate capability of the LRMA0, LRMA2, LRMA4, and LRMA6 samples at different rates. CV profiles of b LRMA0 and c LRMA4 samples with various sweeping rates from 0.1 to 1 mV·s−1, d Ip − v1/2 relationships for both samples. Nyquist diagrams of the LRMA0 and LRMA4 samples e before cycling and f after 200 cycles at 1C

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Fig. 8b–d show the CV experiments of LRMA0 and LRMA4 electrodes in the range of 0.1–1 mV to further investigate the influence of Al doping on the lithium-ion diffusion kinetics. The lithium-ion diffusion coefficient was calculated according to the following Randles − Sevcik equation [56]:

$${I}_{p}=\left(2.69\times {10}^{5}\right){n}^{3/2}{C}_{Li}A{D}_{Li}^{1/2}{v}^{1/2}$$

Ip is the peak current at specific scanning rates, n is the number of transferred electrons, CLi is the molar concentration of Li+ in the cathode material, A is the electrode area, DLi is the lithium-ion diffusion coefficient, and v is the scanning rate. According to Randles − Sevcik equation, we fit a straight line, which reflects the value of DLi, as shown in Fig. 8d. The lithium-ion diffusion coefficients of the LRMA0 and LRMA4 electrodes were 1.07 × 10−12 cm2·s−1 and 2.22 × 10−12 cm2·s−1, respectively. The results show that the lithium-ion diffusion coefficient improved after the Al doping. The electrochemical behavior of doped Al was also investigated, and electrochemical impedance spectroscopy (EIS) tests of the LRMA0 and LRMA4 samples were performed. Fig. 8e shows the Nyquist plots of the LRMA0 and LRMA4 samples, which consisted of a semicircle and a line before cycling. The high-frequency region semicircle reflects the charge transfer impedance (Rct), and the low-frequency region diagonal line represents Weber impedance. Fig. 8f displays the Nyquist plots for LRMA0 and LRMA4 samples, which consisted of two semicircles and a line after 200 cycles at 1 C. The high-frequency region semicircle corresponds to the surface film impedance (Rsf), the intermediate-frequency region semicircle corresponds to the charge transfer impedance (Rct), and the low-frequency region diagonal line is associated with the diffusion of lithium ions in the crystal lattice [23]. Table 3 shows the impedance fitting values of the LRMA0 and LRMA4 samples. The Rct values of the LRMA0 and LRMA4 samples before cycling were 252.5 Ω and 130.2 Ω, respectively, which suggests that the LRMA4 sample had better electrochemical kinetics. After 200 cycles, the Rsf values of the LRMA0 and LRMA4 samples were 48.5 Ω and 27.4 Ω, and the Rct values of the LRMA0 and LRMA4 samples were 627.5 Ω and 270.8 Ω, respectively. Although the impedance of the LRMA0 and LRMA4 samples increased after cycling, the difference was obvious. For comparison, the impedance increase of the LRMA4 sample was lower during cycling.

Table 3 The fitted data from EIS spectra and the lithium-ion diffusion coefficient is calculated by the variable scanning speed CV
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To further understand the fundamental reasons for the difference in electrochemical properties of materials, ex situ XRD analysis was conducted for the LRMA0 and LEMA4 samples after 200 cycles, as shown in Fig. 9a. If the diffraction peaks become broader and less intense, it is suggested that the crystal structure of the sample is distorted. For comparison, the XRD diffraction peaks of the LRMA4 sample had a smaller change, which indicates that the LRMA4 sample had less structural change. Furthermore, the ratio of I003/I104 is usually used to indicate the degree of cation mixing [56]. In contrast, the ratio of LRMA0 was lower than 1.2, which implies its higher cation mixing degree. Fig. 9b shows the Raman spectra of LRMA0 and LRMA4 before cycling. The LRMA0 and LRMA4 samples showed two distinct Raman peaks near 482 and 600 cm−1, which were indexed to the Eg bending and A1g stretching of the layered hexagonal R-3 m symmetry. The weak peaks at 425 cm−1 were attributed to the fingerprint vibration of Li2MnO3 with the C2/m space group. The Raman spectra of the LRMA0 and LRMA4 samples after 200 cycles are shown in Fig. 9c. Compared with the Raman spectrum before cycling, the peak at 425 cm−1 disappeared, since the Li2MnO3 phase disappeared during cycling. Moreover, the LRMA0 and LRMA4 samples before cycling near 600 cm−1 shifted to 622 cm−1 and 604 cm−1 in the spectra of the LRMA0 and LRMA4 samples after cycling measurements. This is the blueshift in Raman spectra after the cycling measurement of Li-rich manganese-based materials [57]. The slight blueshift in the Raman spectra of these Li-rich manganese-based materials is associated with the formation of the spinel phase during cycling. More blueshift corresponds to more spinel phase. In this connection, Al-doped Li-rich manganese-based materials can obviously inhibit the transition from the layer to the spinel phase; hence, they can improve the electrochemical performance because Al doping reduces the lattice parameters and inhibits the migration of TM ions [46, 58].

Fig. 9
figure 9

A The ex situ X-ray diffraction patterns of LRMA0 and LRMA4 samples after 200 cycles. Raman spectra of b before cycling and c after 200 cycles for LRMA0 and LRMA4 samples. XPS spectra of d survey scan, e Ni 2p and f Mn 2p for LRMA0 and LRMA4 samples after 200 cycles

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Fig. 9d–f show the XPS spectra of LRMA0 and LRMA4 after 200 cycles. In the XPS spectrum of Fig. 9d, F, N, and P elements were additionally detected compared with the spectrum of Fig. 5a before the cycle measurement. Elements F and P came from the electrolyte LiPF6, and element N came from NMP. In Fig. 9e, Ni 2p3/2 fit three peaks at approximately 854.8 eV, 856.6 eV, and 858.3 eV, which corresponded to Ni2+, Ni3+, and Ni4+, respectively [59]. For comparison, the peak area of Ni4+ in the LRMA0 sample was the largest, which indicates that LRMA0 had more Ni4+. Ni4+ can cause oxidation reactions of materials, which results in electrolyte decomposition and consumption of cathode active materials to decrease the performance [60] [47]. Fig. 9f displays the Mn 2p spectra of the LRMA0 and LRMA4 samples, and Mn 2p3/2 was fitted to the two peaks of Mn4+ and Mn3+. The fitting results show that LRMA0 contained more Mn3+. Mn3+ underwent a disproportionation reaction and induced the Jahn–Teller effect, which resulted in the loss of cathode active materials and structural collapse. Additionally, most significantly, the increase in Mn3+ decreased the DMV [47]; therefore, the LRMA4 sample showed less attenuation of DMV than LRMA0 in Fig. 6d. The reason is that the Al-O bond (512 kJ·mol−1), which is stronger than the Mn–O bond (402 kJ·mol−1), reduces oxygen release to reduce the production of Mn3+.

Finally, the high-temperature and full-cell performance of the samples were investigated, which are vital tests for any electrode material before commercialization. The cycling property of LRMA0 and LRMA4 samples at 1 C at 55 ℃ is described in Fig. 10a. The specific capacities of the LRMA0 and LRMA4 samples after 200 cycles were approximately 218.5 mAh·g−1 and 230.4 mAh·g−1, and the capacity retention rates were 31.9% and 59.1%, respectively. The cycle performance at high temperature improved because Al doping stabilized the sample structure. The electrochemical performance of the LRMA4 sample in lithium-ion full cells is shown in Fig. 10b. Commercial graphite was selected as the anode material, and the activation process was identical to that in Fig. 6c. Then, the full cell was performed at 1 C for the next 200 cycles. The specific capacity at 1 C was 170.53 mAh·g−1, and the capacity retention rate was 95.6% after 200 cycles at 1 C, which is higher than that of the half cell (89.5%) due to the reduced electrode cross-talk [26]. The specific energy density, which was calculated based on potential profiles, displayed maximum values of approximately 443.7 Wh·kg−1 at 0.1 C and approximately 328.1 Wh·kg−1 at 1 C. After 200 cycles, the specific energy density was 286.9 Wh·kg−1, and the retention rate was 87.4%. It is quite promising that the energy density of 200 Wh·kg−1 could be maintained after 200 cycles at 1 C in commercial batteries, which is considered to satisfy the benchmark requirements of application in electric vehicles.

Fig. 10
figure 10

A Cycling performance of LRMA0 and LRMA4 samples at 1 C and 55 °C. b Energy density and specific capacity of the LRMA4/graphite lithium-ion full cell. The specific energy density was calculated based on the mass of both active materials (cathode and anode)

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Conclusion

A series of Li1.2Ni0.2Mn0.6-xAlxO2 (x = 0, 0.02, 0.04, 0.06) materials was synthesized using a modified Pechini method, and the results prove the effectiveness and feasibility of the Al doping strategy. The results show that doping Al reduces the transition from the layered structure to the spinel structure by lowering the unit cell parameters, reduces oxygen release by introducing stronger Al-O bonds, and suppresses the generation of Mn3+ during the cycle. The capacity retention rates of the LRMA2, LRMA4, and LRMA6 samples were 76.2%, 89.5%, and 81.1%, respectively, after 200 cycles at 1 C, while the capacity retention rate of LRMA0 was only 63.2%. Al doping can significantly inhibit the decay of the DMV during cycling, and the voltage decay of the LRMA4 sample was only 1.11 mV per cycle. The specific capacity of the Al-doped LRMA4 sample at 5 C reached 151 mAh·g−1. Finally, the modified Pechini method can form a Li2CO3 coating on the surface of the material in situ using a one-step method, which decreases the cost. Thus, LRMA4 has the characteristics of fast charge‒discharge performance, good cycle stability, and low cost, and it is an ideal choice for the next-generation lithium-ion battery.