The effects of nanofiber architecture and antimony doping on the

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The effects of nanofiber architecture and antimony doping on the performance of lithium-rich layered oxides: enhancing lithium diffusivity and lattice oxygen stability Ruizhi Yu, Zhijuan Zhang, Sidra Jamil, Jiancheng Chen, Xiaohui Zhang, Xianyou Wang, Zhenhua Yang, Hongbo Shu, and Xiukang Yang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b03608 • Publication Date (Web): 26 Apr 2018 Downloaded from http://pubs.acs.org on April 29, 2018

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ACS Applied Materials & Interfaces

The effects of nanofiber architecture and antimony doping on the performance of lithium-rich layered oxides: enhancing lithium diffusivity and lattice oxygen stability Ruizhi Yu,†,§,ǁ Zhijuan Zhang,┴ Sidra Jamil,†,§,ǁ Jiancheng Chen,†,§,ǁ Xiaohui Zhang,†,§,ǁ Xianyou Wang,*,†,§,ǁ Zhenhua Yang,┴ Hongbo Shu,†,§,ǁ and Xiukang Yang*,†,§,ǁ

†

National Base for International Science & Technology Cooperation, School of Chemistry, Xiangtan

University, Xiangtan 411105, Hunan, China §

National Local Joint Engineering Laboratory for Key Materials of New Energy Storage Battery,

School of Chemistry, Xiangtan University, Xiangtan 411105, Hunan, China ǁ

Hunan Province Key Laboratory of Electrochemical Energy Storage & Conversion, School of

Chemistry, Xiangtan University, Xiangtan 411105, Hunan, China ┴

School of Materials Science and Engineering, Xiangtan University, Xiangtan 411105, Hunan, China

ABSTRACT: Li-rich layered oxides (LLOs) with high specific capacities are favorable cathode materials with high-energy density. Unfortunately, the drawbacks of LLOs such as oxygen release, low conductivity and depressed kinetics for lithium ion transport during cycling can affect the safety and rate capability. Moreover, they suffer severe capacity and voltage fading, which are major challenges for the commercializing development. To cure these issues, herein the synthesis of high performance antimony-doped LLO nanofibers by an electrospinning process is put forward. Based on the combination of theoretical analyses and experimental approaches, it can be found that the one-dimensional (1D) porous micro/nano morphology is in favor of lithium ion diffusion, and the antimony doping can expand the layered phase lattice and further improve the lithium ion diffusion coefficient. Moreover, the antimony doping can decrease the band gap and contribute extra electrons to O within Li2MnO3 phase, thereby enhancing electronic conductivity and 1

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stabilizing lattice oxygen. Benefitting from the unique architecture, reformative electronic structure and enhanced kinetics, the antimony-doped LLO nanofibers possess a high reversible capacity (272.8 mAh g-1) and initial coulombic efficiency (87.8%) at 0.1 C. Moreover, the antimony-doped LLO nanofibers show excellent cycling performance, rate capability and suppressed voltage fading. The capacity retention can reach 86.9% after 200 cycles at 1 C, and even cycling at high rate of 10 C, a capacity of 172.3 mAh g-1 can still be obtained. The favorable results can assist in developing the LLOs material with outstanding electrochemical properties.

KEYWORDS: Lithium-rich layered oxide material, Antimony doping, Nanofiber, stabilized lattice oxygen, expanded layered phase lattice

1. INTRODUCTION In the past decade, lithium-rich layered oxides (LLOs) with composite structures composed of rhombohedral LiTMO2 (space group: R 3m , TM = Mn, Ni, Co, etc.) and monoclinic Li2MnO3 (space group: C2/m) nanodomains, have received considerable attention as high-energy-density cathode materials due to the delivery of high capacity, often exceeding 250 mAh g-1, through applying a high charging voltage of 4.6 V.1-4 Recent researches have demonstrated that the high capacity resulted from joint redox processes of TMa+/TM (a+1)+ and O2-/O22-.5-7 Despite the high capacity, these materials suffer from high initial irreversible capacity, severe discharge voltage decay and insufficient cycling performance, in addition to the inferior rate capability.8-10 The high initial irreversible capacity originates from the activation of Li2MnO3 phase.11 During the initial delithiation, the large overlap of TM 3d band and O 2p band at highly charged states within Li2MnO3 phase can trigger the excessive oxidation of lattice O2- and then results in O2 gas evolution accompanied by the irreversible deintercalation of Li+ from TM layers with a continuous voltage plateau at above 4.4 V, which induces TM ions migration into the Li vacant sites, causing structure rearrangement and the subsequent failure of 2


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partial Li+ ions re-intercalation during discharging.12,13 The continuous capacity fading and discharge voltage decay can lead to a decreased energy density and low accuracy in terms of measuring the state of charge (SOC), which is currently the bottleneck of commercialization of the LLOs.14,15 These degenerations have been demonstrated to be associated with the phase transformation from layered structure to spinel structure.16 Numerous studies have determined that slight octahedral TM ions in TM layers, especially the Mn4+, would migrate to tetrahedral sites in Li layers during charging process. While upon discharging, some of the TM ions may be trapped in the tetrahedral sites or migrate to octahedral sites in Li layers, which initiates the transformation from layered structure to spinel structure.17-19 Furthermore, excessive anionic redox and cationic migration during cycling can lead to large lattice distortions and rearrangements, which trigger particle cracking accompanied by the movement of grain boundary, resulting in the further structural degradation.20,21 Meng’ group presented that the visualization of Litet-V(TM)Li-Litet dumbbell (Litet is the Li+ ions in tetrahedral sites and V(TM)Li represents Li vacant sites in the TM layers) and O vacancy would facilitate the migration of TM ions.22 Cho etc. also proposed that the O2 gas loss related to activation of Li2MnO3 phase can promote the phase transformation.23 As we know, the redox potential (< 3.0 V) of spinel is much lower than layered LiMnO2 (≈3.3 V).15 Thus, the redox potential decreases with the increase of spinel phase. The poor rate capability is correlated to the low conductivity and depressed kinetics itself, which restricts their application in power battery for EVs.8 The latest study exhibited that the undesirable rate capability was also concerned with the plugged Li+ ions diffusion channel.24,25 To overcome above drawbacks, a lot of works have been conducted. Element doping, as an useful approach to boost the physical and electrochemical properties of LLOs, has been widely employed.26-28 Several studies have demonstrated that some foreign atoms (M) can transfer extra electrons to O and alleviate the excessive charge compensation from O for Li+ ions extraction in M-doped Li2MnO3, postponing the irreversible O2 gas release.29,30 Moreover, Xia’ group found that Si doping can lower 3



the overlap of TM 3d band and O 2p band, which can stabilize lattice O2- and obtain high reversible capacity.31 In addition, based on the previous study that TM ions migration is an inducement of the phase transformation and sensitive to the local atomic distribution of target site, some researchers proposed that doping alien ions into the appropriate crystallographic sites in the pristine LLOs structure to clog the kinetically favorable channel of TM ions migration can prohibit the phase transformation and the related discharge voltage decay.19,27,32,33 In our previous work,32 Mg was doped into Li (4h) site. As the segmental occupation of electrochemical inert Mg2+ ions in the Li (4h) sites, the migration path of TM ions can be blocked because of electrostatic repulsion between TM ions and Mg2+. Thus, the continuous capacity degradation and voltage fading of the LLOs material were mitigated. Simultaneously, immobile Mg2+ ions in Li layers served as fixed “pillars”, which can also enhance the rate capability. Though some foreign ions doping could increase the rate capability due to the improvement of electronic structures and lithium diffusion kinetics, numerous reports have evidenced that the morphology and surface structure are critical for Li+ ion transport and the rate capability of LLOs.22,34-38 Xu etc. reported a hetero-epitaxial nanostructural spinel/LLO composite with anisotropic Li+ ions diffusion channels, which delivered favorable discharge capacity and rate capability.22 Li etc. synthesized a 3D reticular Li1.2Ni0.2Mn0.6O2 cathode material, and it delivered a high discharge capacity of 135.7 mA h g-1 at 1000 mA g-1.36 Especially, LLOs with 1D nanorods or nanowires morphology caught much attention. For example, Yang etc. prepared a series of LLO nanorods and explored a sight to obtain pure LLOs with adjustable rate capability.37 Our group excogitated and synthesized spinel/layered heterostructured LLO Nanowires, both the rate capability and cycling performance were significantly increased.38 Motivated by the considerations above, morphology design combined with element doping may develop a feasible approach to meliorate the physical and electrochemical performance of LLOs. Thus, selecting appropriate synthetic method for controlling special morphology and substituent elements becomes crucial. Electrospinning,39-41 as 4


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a facile and cost-efficient technique for fabricating 1D nanofibers or nanowires, has been widely applied to prepare electrode materials. Compared with other methods (solvothermal method, co-precipitation method, etc.), the morphology of the final composite can be better controlled.40,41 Furthermore, electrospinning can homogenize the distribution of elements, which is especially beneficial to element doping.42,43 Recently, Tian etc. reported that partial substitution of antimony (Sb) helped to stabilize the structure and promote Li+ ions diffusion coefficient of Li2MoO3 based on calculation, which seemed to be superior to other elements (Mo, Nb, Pd, Ti, Sn, Tc and Ru).44 To the best of our knowledge, the Sb doped Li2MnO3-based LLOs have never been synthesized and the electrochemical mechanism and behavior were obscure. Thus, it is significant to explore the effect of Sb doping on electronic structure and lattice texture, which may give us more insight into design of high-performance LLOs. In this work, Sb doped LLO nanofibers were synthesized by electrospinning, and the relative physicochemical and electrochemical properties as well as the underlying mechanisms were studied in detail.

2. RESULTS AND DISCUSSION The morphologies of the conventional pure electrospun films (PAN-F), conventional pure Li1.2Mn0.54Ni0.13Co0.13O2+δ nanofibers (C-LMNCO), Sb-doped electrospun films (S/PAN-F) and

Sb-doped

Li1.2Mn0.513Ni0.13Co0.13Sb0.027O2+δ

nanofibers (S-LMNCO) were examined by SEM, as shown in Figure 1. The as-electrospun PAN-F (Figure 1a and b) and S/PAN-F (Figure 1c and d) precursor fibers exhibit similar smooth surfaces with diameters of ~600 nm and dozens of micrometers in lengths. Though the diameters were reduced to 150 ± 50 nm because of the decomposition of acetate salts and polymer during calcination, the morphologies of 1D nanofibers can be maintained for both C-LMNCO (Figure 1e and f) and S-LMNCO (Figure 1g and h) cathode materials. Furthermore, it can be found that each fiber of either C-LMNCO or S-LMNCO is assembled by agglomeration of nanoparticles, and their surfaces become much rougher than precursor fibers. Such 1D 5



porous micro/nano morphologies can provide more interface areas between cathode materials and electrolyte, which is expected to form a short Li+ diffusion path, facilitate electrolyte infiltrating to the nanoparticles and promote Li+ diffusion,42 as shown in the schematic diagram (Figure 1i). Additionally, the effective conducting nanofiber can also enhance the electrical contact between the nanoparticles, thus leading to an improved electrical conductivity.42 The concentrations of metal elements in C-LMNCO and S-LMNCO were performed by ICP-OES and tabulated in Table 1. The results are consistent with the expected stoichiometry within experimental errors. The crystal structures of the C-LMNCO and S-LMNCO were determined by XRD (Figure 1j). Both cathode materials deliver the characteristics of the layered α-NaFeO2 structure with the combined XRD pattern of Li2MnO3 (space group: C2/m, marked by blue color) and LiMO2 (space group: R 3m , marked by purple color).8 The intensity ratio of the (003)/(104) peaks is used for quantifying the degree of cation mixing, and a value lower than 1.2 represents cacoethic cation mixing.45,46 The I(003)/I(104) peak ratios of the C-LMNCO and S-LMNCO samples are 1.44 and 1.53, which indicates suppressed cation mixing. Moreover, the well-split (006)/(102) and (018)/(110) peaks can be clearly found, suggesting that well-defined layered phases with high crystallinity are formed in both samples.47,48 Figure 1k-n display the XPS results of C-LMNCO and S-LMNCO samples, and all binding energies (BEs) were corrected for specimen charging by referencing them to the C 1s peak (284.6 eV). The Ni 2p3/2 peaks (Figure 1k) of the C-LMNCO and S-LMNCO samples are centered at 854.5 eV and 854.9 eV, indicating the coexistence of Ni2+ (854.2-854.6 eV) and Ni3+ (855.5-855.7 eV) and the main chemical valence of Ni is +2.49,50 Moreover, the Ni 2p3/2 peak of the S-LMNCO sample shifts to a higher binding energy position compared with C-LMNCO sample, suggesting that the S-LMNCO possesses higher concentration of Ni3+. The Co 2p3/2 (Figure 1l) and Mn 2p3/2 (Figure 1m) peaks of C-LMNCO and S-LMNCO samples are located at 780.0 eV and 642.4 eV, demonstrating that the main chemical valences of Co and Mn are +3 and +4 in both 6


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samples.50 In addition, a peak occurring about 538.9 eV (Figure 1n) can only be found in the S-LMNCO sample, which is comparable with the Sb 3d3/2 binding energy obtained in Sb2O3,51 evidencing the existence of Sb3+ in the S-LMNCO sample, and the local environment change originated from the replacement of Mn4+ by Sb3+ with lower valence causes the increase of concentration Ni3+ in the S-LMNCO sample. In general, the higher concentration of Ni3+ in the LLOs is beneficial to the Li+ diffusion and structural stability during cycling.50,52

Figure 1. SEM images of (a and b) PAN-F and (c and d) S/PAN-F precursor fibers, and corresponding (e and f) C-LMNCO and (g and h) S-LMNCO. (i) Schematic diagram of the Li+ diffusion pathway of the LLO nanofibers. (j) XRD patterns of C-LMNCO and S-LMNCO. XPS spectra of (k) Ni 2p, (l) Co 2p, (m) Mn 2p and (n) Sb 3d+O 1s of C-LMNCO and S-LMNCO. Table 1. Chemical Compositions of the C-LMNCO and S-LMNCO Determined by ICP-OES 7



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Normalized element content (use Ni = 0.130 for all samples)

Samples

Li1.200Mn0.540Ni0.130Co0.130O2+δ

Li

Mn

Ni

Co

Sb

1.198

0.542

0.130

0.131

\

1.196

0.514

0.130

0.129

0.026

(C-LMNCO) Li1.200Mn0.513Ni0.130Co0.130Sb0.027O2+δ (S-LMNCO)

To obtain the detail structural characterizations and elemental distribution of C-LMNCO and S-LMNCO samples, TEM images were examined with their corresponding EDS elemental mapping and fast Fourier Transform (FFT) patterns (Figure 2). Figure 2a-d (C-LMNCO) and Figure 2e-i (S-LMNCO) illustrate the low-resolution TEM images and the corresponding EDS elemental mapping, demonstrating that the element distributions in both cathode materials are homogeneous. Figure 2j and k display the high-resolution TEM (HRTEM) images of LiMO2 nanodomain with the [110] axis and Li2MO3 nanodomain with the [312] axis of the C-LMNCO sample, and the insets are their corresponding FFT patterns. The lattice spacings in Figure 2j are 0.474, 0.144 and 0.138 nm, matching the d(003), d(110) and d(113) planes.53 And three kinds of lattice spacings (0.243, 0.317 and 0.367 nm) can be observed in Figure 2k, which are indexed to (130), (021), and (111), respectively.54 Furthermore, the ideal LiMO2 and Li2MnO3 crystals along the same direction with the TEM images were investigated, as shown in Figure 2l and m. It can be concluded that the atomic configuration of C-LMNCO sample is in good accordance with the ideal atomic model. Figure 2n and o demonstrate the fine atomic arrangements LiMO2 phase along the [110] axis and Li2MO3 phase along the [001] axis of the S-LMNCO sample, and their ideal structural models were exhibited in Figure 2p and q. Obviously, the synthesized S-LMNCO sample possesses the similar crystal structure and atomic arrangement with C-LMNCO sample.

8


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Figure 2. (a) TEM image and EDS mapping of (b) Ni, (c) Co and (d) Mn of C-LMNCO sample; (e) TEM image and EDS mapping of (f) Ni, (g) Co, (h) Mn and (i) Sb of S-LMNCO sample; (j, k) HRTEM images with the corresponding FFT patterns and (l, m) simulative crystal structures of (j, k); (n, o) HRTEM images with the corresponding FFT patterns and (p, q) simulative crystal structures of (n, o). 9



The electrochemical tests were performed on the C-LMNCO and S-LMNCO cells. Figure 3a and b exhibit the galvanostatic charge-discharge curves and the corresponding differential capacity profiles (dQ/dV vs. V) of C-LMNCO and S-LMNCO cells during the first cycle at 0.1 C between 2.0 and 4.6 V. In the initial cycle, both C-LMNCO and S-LMNCO cells deliver high charge capacities (361.7 mAh g-1 for C-LMNCO and 310.6 mAh g-1 for S-LMNCO) and discharge capacities (278.5 mAh g-1 for C-LMNCO, 272.8 mAh g-1 for S-LMNCO) with the electrochemical characteristics of LLOs. As illustrated in Figure 3a, the initial charging curves were separated into two parts: one is a short slope line, which corresponds to Li+ deintercalation from the LiMO2 structure with the Co3+/Co4+ and Ni2+/Ni4+ oxidation,20 and the capacities of C-LMNCO and S-LMNCO cells at this stage are almost the same. The other is a long plateau, representing the activation of Li2MnO3 phase and the accompanying irreversible O2 gas release, and thus the decrease of capacity in the plateau region indicates that the O2 gas release is suppressed for the S-LMNCO sample.12 Furthermore, the S-LMNCO cell exhibits the reduced polarization compared with C-LMNCO cell, as shown in Figure 3b. Figure 3c-e display the comparison of rate capability of C-LMNCO and S-LMNCO cells. Both samples deliver acceptable rate capability, but a much better performance can be obtained for the S-LMNCO. At 0.2 C, the difference of discharge capacities between C-LMNCO and S-LMNCO cells is 16 mAh g-1, while at 0.5 C, the capacity difference increases to ~32 mAh g-1, and this difference gradually increases with the increase of current rate. At 10 C, the S-LMNCO cell delivers an exceptional discharge capacity of 172.3 mAh g-1, while the C-LMNCO cell presents lower discharge capacity of 116.1 mAh g-1. Interestingly, after the high-rate measurements, the C-LMNCO cell cycled at 0.1 C exhibits a slight decrease in the discharge capacity compared with its initial value. By comparison, the S-LMNCO cell can nearly recover its initial discharge capacity, thereby suggesting its favorable reversibility. The cycling performance of the C-LMNCO and S-LMNCO cells were evaluated at 1 C between 2.0 and 4.6 V. The initial discharge capacities of the C-LMNCO and 10


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S-LMNCO cells at 1 C are 211.2 and 247.8 mA h g-1 (Figure 4a). After 200 cycles, the discharge capacity of the C-LMNCO cell declines fast and the capacity retention is only 57.4%, whereas the discharge capacity of S-LMNCO cell can retain 215.3 mA h g-1 after 200 cycles with the capacity retention of 86.9%. In order to further understand the cycling behavior, the 1st and 200th charge/discharge profiles at 1 C and their corresponding differential capacity profiles were displayed in Figure 4b and c. For the 1st charge/discharge profile, the cationic redox proceeds at the similar potentials during either charging or discharging. The dQ/dV peaks at 3.9 V are attributed to Ni2+/3+/4+ and Co3+/4+ activity and the peaks at low potentials are related to the minor Mn3+/4+ capacity.20 By comparison, anionic redox occurs asymmetrically during charging/discharging. In the process of charging, anionic redox takes place from 2.0 to 4.6 V especially above 4.2 V. During discharging, the main anionic reduction occurs below 3.6 V and the remaining reduction exists at high potential (above 4.2 V).20 After 200 cycles, the reduction peaks shift to the lower potentials than before, which are known as the voltage fading and hysteresis, suggesting the severe phase transformation and lattice O2- degeneration. Although the voltage fading can’t be eliminated for the S-LMNCO sample, the reproducibility of discharge curves and voltage plateaus during the long-term cycling is superior to the C-LMNCO sample. The larger discharge capacities and decreased overpotentials portend an improvement in specific energy for the S-LMNCO sample.

11



Figure 3. Electrochemical characteristics during initial cycling and rate capabilities of C-LMNCO and S-LMNCO cells. (a) First cycling curves of C-LMNCO and S-LMNCO cells and the related (b) dQ/dV profiles; (c) discharge capacity at various current rates and the corresponding discharge curves of the (d) C-LMNCO and (e) S-LMNCO cells.

12


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Figure 4. Cycling stability of C-LMNCO and S-LMNCO cells at 1 C. (a) Variation in discharge capacity vs. the cycle number; (b) charge/discharge curves of the 1st and 200th cycles for the C-LMNCO and S-LMNCO cells; (c) dQ/dV profiles of the C-LMNCO and S-LMNCO cells in (b).

Since the rate capability of S-LMNCO sample is superior to C-LMNCO sample, the kinetic performance of Li+ ions in both cathodes were determined. Galvanostatic intermittent titration technique (GITT) based on chronopotentiometry at nearly thermodynamic equilibrium conditions is a reliable method to estimate the Li+ diffusion coefficient ( DLi + ). Figure 5a and b illustrated the GITT profiles of C-LMNCO and S-LMNCO cells, and the insets are the t vs. V profiles of single GITT titrations. Herein, the DLi + values were calculated according to Equ. (1) as follows:50 2

DLi + =

2

4  Vm   dE dx  L2 , t

The effects of nanofiber architecture and antimony doping on the

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