- Email: [email protected]
The effect of Sn substitution on the structure and oxygen activity of Na0.67Ni0.33Mn0.67O2 cathode materials for sodium ion batteries
Journal of Power Sources xxx (xxxx) xxx
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
The effect of Sn substitution on the structure and oxygen activity of Na0.67Ni0.33Mn0.67O2 cathode materials for sodium ion batteries Jinke Li a, 1, Tim Risthaus b, 1, Jun Wang b, **, Dong Zhou b, Xin He a, ***, Niloofar Ehteshami a, Vadim Murzin c, Alex Friesen b, Haidong Liu b, Xu Hou a, Marcel Diehl b, Elie Paillard a, Martin Winter a, b, Jie Li a, * a b c
Helmholtz-Institute Muenster (HI MS), IEK-12, Forschungszentrum Juelich GmbH, Corrensstr. 46, Muenster, 48149, Germany MEET Battery Research Center, Institute of Physical Chemistry, University of Muenster, Corrensstr. 46, Muenster, 48149, Germany Deutsches Elektronen-Synchrotron, Notkestr. 85, Hamburg, 22607, Germany
H I G H L I G H T S
� P2/P3 Na0.67Ni0.33Mn0.67-xSnxO2 cathode materials for SIBs are proposed. � Na0.67Ni0.33Mn0.66Sn0.01O2 exhibits high reversible capacities in 1.5–4.5 V. � Ex-situ XAS and XPS are applied to characterize the involved redox reactions. � Oxygen activity in the family of the Na0.67Ni0.33Mn0.67O2 materials. A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium ion batteries Cathode Sodium nickel manganese oxides Sn substitution Oxygen activity
A series of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) materials with mixed P2/P3 phases are synthesized with a conventional solid-state reaction method and investigated as cathode materials for sodium ion batteries. The effects of Sn substitution on the structure and electrochemical performance of the Na0.67Ni0.33Mn0.67O2 are systematically investigated. The substituted samples show smaller particle sizes compared to the pristine one and the P2:P3 phase ratio highly depends on the substitution amount. The best electrochemical performance is ob tained by Na0.67Ni0.33Mn0.66Sn0.01O2, and it delivers a discharge capacity of 245 mA h g 1 in 1.5–4.5 V (vs. Na| Naþ), which is the highest result for Na0.67Ni0.33Mn0.67O2 materials reported so far. The ex situ X-ray absorption spectroscopy and X-ray photoelectron spectroscopy measurements reveal that the oxygen ions participate in the redox reactions within the wide voltage range of 1.5–4.5 V. The increased capacity can be attributed to the smaller particle size, which results in more oxygen activity and then higher capacity.
1. Introduction Sodium ion batteries (SIBs) are attracting increasing interest as cost efficient and sustainable alternatives to the prosperous lithium ion batteries (LIBs) [1], due to the vast availability of Na resources and similar electrochemical properties between SIBs and LIBs, especially for large-scale energy storage systems [2,3]. Among reported cathode can didates, sodium layered oxides (NaMO2, M ¼ Co, Ni, Mn, Cr, Fe, V etc.) have obtained marvelous attentions in light of their high capacities and
easy synthesis route [4–6]. Generally, they are categorized into two main groups in view of structure, namely P2 and O3 type, where ‘P’ and ‘O’ refer to the occupation sites of Na ions, i.e. prismatic and octahedral, respectively, and the following digit means the number of Na layers in each repeated unit [7]. Since the standard electrochemical potential of Na|Naþ ( 2.71 V vs. SHE) is higher than that of Li|Liþ ( 3.04 V vs. SHE), the energy density of SIBs is usually lower than that of LIBs in similar cell configuration. Therefore, the development of NaxMO2 with high output voltage plateau and discharge capacity plays a key role for
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (J. Wang), [email protected] (X. He), [email protected] (J. Li). 1 These authors contributed equally to this work. https://doi.org/10.1016/j.jpowsour.2019.227554 Received 1 October 2019; Received in revised form 15 November 2019; Accepted 30 November 2019 0378-7753/© 2019 Published by Elsevier B.V.
Please cite this article as: Jinke Li, Journal of Power Sources, https://doi.org/10.1016/j.jpowsour.2019.227554
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
the practical application of SIBs [2,8]. Among various NaxMO2 materials, Na0.67Ni0.33Mn0.67O2 possesses high theoretical capacity (173 mA h g 1) and high working potential associated with the redox potential of Ni2þ|Ni4þ (>3.0 V vs. Na|Naþ), in addition to its high tolerance vs. the storage environment [9,10]. Consequently, it has attracted increasing attentions as a promising cathode material since its first exploration by Dahn’s group [5,11]. In addition to the conventional redox reactions from transition metals, the anionic redox activity in NaxMO2, namely the oxygen redox reaction, has also been explored intensively recently, and is proposed as an effective route to get cathode material with large reversible capacity [12–15]. The oxygen redox reaction of Na0.67Ni0.33Mn0.67O2 has been proven by both DFT calculation and experiments [15,16]. For example, our previous work confirmed that the excessive discharge capacity (vs. the theoretical capacity) was attributed to the oxygen activity when the Na0.67Ni0.33Mn0.67O2|Na cell was cycled in a wide voltage range of 1.5–4.5 V [15]. It is believed that the substitution of cation ions can affect the metal-oxygen covalence of the materials, thus the oxygen behaviors during charge and discharge process can be altered [17–19]. In this case, substitution of Ni or Mn in Na0.67Ni0.33Mn0.67O2 will potentially affect its oxygen activity. Sn, as non-toxic element, with a high bond dissociation energy to oxygen (Sn–O; ΔHf298 ¼ 548 (21) kJ mol 1), has been investigated as the substitution element in Li-rich cathode materials, and the oxygen activities at high voltage plateau are affected consequently [17,20]. However, it is rarely reported as the substitution element in the cathode materials with oxygen activity in SIBs. Rong et al. confirmed that, the Sn substitution in Na0.67Ni0.33Mn0.67O2 can lead to the formation of O3 phase, the increased working voltage and smooth voltage profile in the voltage range of 2.5–4.1 V [21]. However, due to the narrow voltage range applied, the influence of Sn substitution on the oxygen activity of Na0.67Ni0.33Mn0.67O2 was not discussed in that paper. Herein, a series of P2/P3 Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) materials are prepared by a combination of co-precipitation and solid-state reaction methods. Among these four active materials, the Na0.67Ni0.33Mn0.66Sn0.01O2 delivers an initial charge capacity of 166 mA h g 1 and a high discharge capacity of 245 mA h g 1 within a wide voltage range of 1.5–4.5 V. While in a narrower operating voltage range (2.0–4.2 V), the results show that Sn substitution has no influence on the improvement of capacity. Additionally, the effects of Sn substitution on the redox chemistry and safety property are also explored.
Fig. 1. XRD patterns of Na0.67Ni0.33Mn0.67O2 (black), Na0.67Ni0.33Mn0.66Sn0.01O2 (yellow), Na0.67Ni0.33Mn0.64Sn0.03O2 (green) and Na0.67Ni0.33Mn0.62Sn0.05O2 (blue). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
equipped with a 532 nm argon-ion laser. The X-ray absorption spec troscopy (XAS) measurements were carried out at the P64 beamline of the Deutsches Elektronen-Synchrotron, Hamburg, Germany, using Si (111) double-crystal monochromator (FMB Oxford). X-ray photoelec tron spectroscopy (XPS) was performed with an Axis Ultra DLD (Kratos) by using monochromatic Al-Kα radiation and the calibration of the spectra was done with respect to adventitious carbon (284.5 eV). Thermogravimetric analysis (TGA) of the precursors before annealing was performed under air atmosphere at a heating rate of 5 K min 1 from room temperature to 800 � C with a TGA Q5000 IR devices from TA Instruments. Differential scanning calorimetry (DSC) measurements were realized by using a DSC Q2000 device from TA Instruments (USA; V24.11 build 124). All the electrode samples were heated to 300 � C at 5 K min 1 in high-pressure capsules. The cycled electrode samples were disassembled and then immersed in 1 M NaPF6 in EC/DMC (1/1, vol.) with 1 wt% FEC electrolyte. 2.3. Electrochemical characterization
2. Experimental
The electrochemical behaviors were analyzed with R2032-type coin cells in two electrode configuration [22]. The electrode consists of active material, conductive carbon (carbon black, Super C65, Imerys) and poly (vinylidene difluoride) (PVdF) binder (Kynar FLEX 761A, Arkema Group) at a weight ratio of 8:1:1. 1-methyl-2-pyrrolidinone (NMP, ACROS Organics) was used as solvent during the slurry preparation and the well mixed slurry was casted on the Al foil current collector. After being dried, electrode tapes were punched into Ø13 mm discs with a mass loading of �2 mg cm 2. All cells were assembled in an Ar-filled glove box with 1 M NaPF6 dissolved in 1:1 vol ratio of ethylene car bonate (EC) and dimethyl carbonate (DMC), together with 1 wt% fluo roethylene carbonate (FEC) as electrolyte. Sodium metal (99.8%, Acros Organics) and glass fiber (GF/D, Whatman) were used as anode and seperator, respectively. The cells were cycled galvanostatically at different C rates (1 C ¼ 200 mA g 1) between 1.5 and 4.5 V or 2.0–4.2 V at 20 � C using Maccor series 4000 battery tester. Cyclic voltammetry (CV) measurements were applied on a VSP electrochemical workstation (Bio-logic) with a sweep rate of 0.1 mV s 1.
2.1. Materials synthesis All chemicals were purchased from Sigma-Aldrich and used without any further purification. Stoichiometric amounts of NiSO4⋅6H2O, MnSO4⋅H2O and SnCl4⋅5H2O were dissolved in distilled water, then the solution was slowly dropped into a 1 M NaOH aqueous solution, after wards, the slurry was continuously stirred at 60 � C for 12 h in Ar at mosphere. The obtained precipitate was filtered, washed with distilled water and dried in air at 100 � C. At the end, the precipitate was thor oughly mixed with Na2CO3 by ball milling and calcined at 500 � C for 5 h, then 800 � C for 12 h in air to obtain final products. 2.2. Materials characterization Inductively coupled plasma optical emission spectroscopy (ICP-OES) was applied on the Spectro ARCOS instrument for the element deter mination. Powder X-ray diffraction (XRD) measurements were per formed on a Bruker D8 advance diffractometer with Cu kα radiation (λ ¼ 1.5418 Å) in the 2θ range from 10� to 90� . Scanning electron micro scopy (SEM) images and energy dispersive X-ray analysis (EDX) were recorded on the EVO MA 10 microscope (Zeiss). Raman spectra were measured using a Bruker SENTERRA dispersive Raman microscope,
3. Results and discussions Initially, the composition of all samples was determined by ICP-OES analysis. The results (Table S1) reflect that the molar ratio of Ni, Mn and Sn in the four samples agrees with the intended value. The TGA-DSC 2
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 2. SEM images of (a) Na0.67Ni0.33Mn0.67O2, (b) Na0.67Ni0.33Mn0.66Sn0.01O2, (c) Na0.67Ni0.33Mn0.64Sn0.03O2 and (d) Na0.67Ni0.33Mn0.62Sn0.05O2.
curves of the pristine Na0.67Ni0.33Mn0.67O2 and 5% Sn substitution Na0.67Ni0.33Mn0.62Sn0.05O2 precursors are shown in Fig. S1. Three weight loss processes can be seen in the TGA curves, i.e., below 100 � C, 100–250 � C and 250–800 � C which can be attributed to the loss of adsorbed H2O on the surface of the precursors, the decomposition of the hydroxides and carbonate, and the formation of crystal phase, respec tively. Except the first process (<100 � C), the weight loss is nearly the same (18.5% and 18.8%) when the temperature is higher than 100 � C. The DSC curves show the similar trend and overlap each other in the high temperature range. The DSC and TGA results both indicate that, Sn substitution shows no obvious influence on the annealing process. XRD measurements were applied to examine the crystal structure of the four samples (Fig. 1). As displayed, the Bragg diffraction peaks of the four samples are well indexed with the mixture of a hexagonal P2 phase (JCPDS Nr. 54–0894 with space group P63/mmc) and a tetragonal P3 phase (JCPDS Nr. 54–0839 with space group R3m). To obtain the ratios between P2 and P3 phases, the collected XRD data of the four materials are refined with Rietveld refinements (Fig. S2 and Table S2). The Na0.67Ni0.33Mn0.67O2 structure exhibits in majority P2 phase, while the substituted materials exhibit the main characteristics of the P3 phase, reflecting that Sn substitution leads to a difference in crystal structure. It is reported that the structure of Ni, Mn based NaxMO2 depends largely on the employed pH value of the precipitation process [23] and the annealing temperature [6,24]. The high annealing temperature is favorable to P2 phase formation while at low temperature the formation of the P3 phase is preferred [25]. In this work, the applied 800 � C leads to the formation of a mix of P2 and P3 phases. And as Sn4þ is involved in the precipitation process, the pH value of the mixture solution tends to decrease owing to partial hydrolysis of Sn4þ ions, thus, the Sn substituted materials display the main feature of P3 phase. The lattice parameters (Table S3) show that with an increase in Sn substitution amount, for both the P2 and P3 phase, the a and c axes of the samples also increase gradually, which can be attributed to the differ ence of ionic radii since Sn4þ (0.71 Å) has larger ionic radius than Mn4þ (0.53 Å) and Ni2þ (0.69 Å) [26]. This good agreement between the lattice parameter and ionic radii suggests that Sn is successfully substituted in the crystal structure. Besides, to exclude the formation of SnO2 in the obtained Sn-substitution materials, the XRD measurement on the pure SnO2 was conducted (Fig. S3). When compared with the
XRD patterns of the Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0.01, 0.03, 0.05), it is concluded that no SnO2 is formed during synthesis, indicating that Sn is substituted into the crystal structure of the investigated materials. Raman analyses (Fig. S4) shows that for all four active materials, five bands are observed, namely 630, 586, 478, 380, 348 cm 1, which are associated with the A1g þ 3E2g þ E1g mode, similar to the other sodium layered oxides NaxCoO2 [27], Na0.67Fe0.5Mn0.5O2 [28] and Na0.67Alx Co1-xO2 [29]. Similar Raman spectra of Na0.67Ni0.33Mn0.67O2 were also obtained by Noguchi etc. [26] and Guo etc. [30]. Besides, no bonds corresponding to the Sn-contained impurities such as SnO2 can be seen. According to the previous reports, the A1g and E1g mode are assigned to the O vibrations and three E2g modes come from the Na and O vibrations [31,32]. With different amount of Sn, the intensities of all bands vary, and the band at 630 cm 1 representing the symmetric stretching of MnO6 octahedra gets stronger [31], indicating that the short-range structure of the materials are influenced by Sn substitution. This result also manifests the hypothesis that Sn is indeed substituted in the lattice structure instead of forming multi-phase mixtures. In materials science, ‘substitution’ is commonly used as the term describing the replacement of one element/ion by another one(s) in a given crystal structure. The amount of employed ‘impure’ element/ion often reaches level of several percent. The term ‘doping’ originates from semiconductor science, where the dopants are introduced as additives (not as substituent), in a very small amount, into pure intrinsic semi conductors to modify their electrical properties. For battery materials, ‘doping’ is, in some cases, also used to represent atomic/ionic substi tution when the amount of substituent is relatively low. Despite the nonconsistent use in battery R&D and related literature, the present work will use the term ‘substitution’ instead of ‘doping’ since the scientific approach of this work aims at substitution of a certain amount of Mn by Sn. The SEM images of the four materials are illustrated in Fig. 2. The pristine Na0.67Ni0.33Mn0.67O2 is composed of flake like primary particle. In contrast, the Sn-containing materials appear as much smaller and agglomerated spherical particles. The SEM results reveal that Sn sub stitution can significantly affect the particle shape and size of the ma terials. The EDX mapping was carried out on the 5% Sn substitution Na0.67Ni0.33Mn0.62Sn0.05O2 samples (displayed in Fig. S5) as represen tative, and the result presents a homogenous element distribution 3
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
discharged sequentially at 0.1, 0.2, 0.5, 1, 2, 5, 10 and 1 C each for five cycles. All the four materials exhibit obvious capacity fading at low current rates. Nevertheless, the highest capacities at each rate is still obtained from Na0.67Ni0.33Mn0.66Sn0.01O2. In detail, the values (average value for five cycles) are 224, 180, 145, 118, 103, 84 and 71 mA h g 1 at 0.1, 0.2, 0.5, 1, 2, 5 and 10 C, respectively. When the current density returns back to 1 C, ~96% capacity recovery is achieved. For comparison, the Na0.67Ni0.33Mn0.67-xSnxO2||Na (x ¼ 0, 0.01, 0.03 and 0.05) cells was also charged/discharged in the voltage range of 2.0–4.2 V (Fig. S6). Within the narrow voltage window of 2.0–4.2 V, the initial discharge capacities of all the electrodes decrease dramatically to less than 90 mA h g 1. However, the cycling stabilities are promising. After 50 cycles, 101%, 100%, 101% and 101% of the initial capacities are maintained with Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64Sn0.03O2 and Na0.67Ni0.33Mn0.62Sn0.05O2, respec tively. In addition, the materials also exhibit excellent rate capabilities (Fig. S6c). For example, at 10 C, Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33 Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64Sn0.03O2 and Na0.67Ni0.33Mn0.62 Sn0.05O2 maintain the discharge capacities of 78, 80, 77 and 72 mA h g 1, corresponding to 85%, 86%, 88% and 86% of their initial capacities at 0.1 C, respectively. As the current density returning to 1 C, the discharge capacities of 87, 87, 81 and 78 mA h g 1 are achieved with ~100% capacity retentions of their initial specific capacities at 1 C, suggesting good structural reversibility in the 2.0–4.2 V voltage range. Overall, in the voltage range of 1.5–4.5 V, the Na0.67Ni0.33Mn0.67xSnxO2 (x ¼ 0, 0.01, 0.03 and 0.05) materials show high a capacity of >220 mA h g 1 at 0.1 C. For the Na0.67Ni0.33Mn0.66Sn0.01O2, 245 mA h g 1 is the largest discharge capacity reported so far from the family of Na0.67Ni0.33Mn0.67O2 compounds. The value far exceeds the theoretical capacity of 172 mA h g 1 (considering only the redox reaction of Ni2þ| Ni4þ), which we attribute to the additional capacity contribution by oxygen redox [15]. Meanwhile, 0.01 Sn substitution leads to 8.4% ca pacity increase (from 226 mA h g 1 delivered by pristine sample to 245 mA h g 1), which is probably because that Sn substitution enhances the oxygen activity. However, it should be noted that an enhanced oxygen activity also leads to a relatively poor cycling performance. Though, the detrimental structural phase transition at high voltage with huge volume shrinkage or serious reactions between electrode (surface) and electro lyte would contribute to the performance fading [9]. In this case, the unsatisfactory reversibility of the oxygen redox reactions which are commonly coupled with oxygen loss [18], and the slow kinetics of the oxygen activity may also take responsibility in the fading [33,34]. Although the cycling performance and the rate capability of these ma terials can be largely enhanced when narrowing the voltage window is applied, this is achieved at the expense of ~30% of the total capacity (based on the value obtained at 50th cycle). Nevertheless, it is urgent to understand the charge compensation mechanisms of these materials upon Na extraction and insertion. Thus, CV measurements in different voltage ranges of 1.5–4.5 V and 2.0–4.2 V were performed on all four Na0.67Ni0.33Mn0.67-xSnxO2||Na cells. All electrodes for the CV measurement were selected with the same size and similar mass loading. The CV curves of the four materials in the voltage range of 1.5–4.5 V are displayed in Fig. 4. The un-substituted Na0.67Ni0.33Mn0.67O2 exhibits a series of complex current peaks. Generally speaking, the strong peaks higher than 4.2 V are due to phase transformation coupled with oxygen redox reactions, and the peaks between 2.5 and 4.0 V can be assigned to the redox reactions of Ni2þ| Ni4þ accompanied with Naþ/vacancy ordering occurring within the sodium layers. The redox peaks lower than 2.5 V refer to the Mn redox reactions [19,35]. In contrast to un-substituted Na0.67Ni0.33Mn0.67O2, the Sn substituted materials exhibit similar peak positions in the initial cycle, reflecting a similar reaction mechanism between them, regardless of the ratio of P2 and P3 phases. Na0.67Ni0.33Mn0.66Sn0.01O2 shows the strongest peaks in the voltage range of >4.2 V and <2.5 V. With higher Sn amounts (x ¼ 0.03 and 0.05), the intensity of these peaks decreases. Besides, it is notable that for all materials, the intensity of the peak
Fig. 3. (a) Initial charge and discharge curves, (b) cycling performance and (c) rate capabilities of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) in the voltage range of 1.5–4.5 V.
throughout the particle. The first charge and discharge curves of the Na0.67Ni0.33Mn0.671 xSnxO2||Na (x ¼ 0, 0.01, 0.03 and 0.05) cells at 0.1 C (1 C ¼ 200 mA g ) in the voltage range of 1.5–4.5 V are displayed in Fig. 3a. As illustrated, the initial charge and discharge capacities of the pristine sample are 161 mA h g 1 and 226 mA h g 1, respectively. In contrast, the Na0.67Ni0.33Mn0.66Sn0.01O2 (166 and 245 mA h g 1) and Na0.67Ni0.33Mn0.64Sn0.03O2 (162 and 230 mA h g 1) show higher charge and discharge capacities, while the sample with 5% Sn substitution delivers lower ones (155 and 220 mA h g 1). The cycling performance of the four materials at 1 C is compared in Fig. 3b. After 50 cycles at 1 C, the Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64 Sn0.03O2 and Na0.67Ni0.33Mn0.62Sn0.05O2 deliver discharge capacities of 122, 129, 117, 113 mA h g 1, respectively, which are associated with capacity retentions of 81%, 73%, 72% and 85%, respectively. Fig. 3c shows the rate performance of the cells which were charged and 4
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 4. Cyclic voltammograms of the Na0.67Ni0.33Mn0.67-xSnxO2||Na cell (a) x ¼ 0, (b) x ¼ 0.01, (c) x ¼ 0.03 and (d) x ¼ 0.05 at a scan rate of 0.1 mV s voltage range of 1.5–4.5 V.
higher than 4.2 V has a significantly positive correlation with that of the peak lower than 2.5 V, i.e., as the peaks >4.2 V become stronger, the peaks <2.5 V are also stronger, reflecting that the oxygen activity cor relates with the redox reaction of the Mn ions. Upon cycling, the in tensity of the peak (>4.2 V) is decreasing due to the poor reversibility of the reactions, which are responsible for the observed capacity fading in Fig. 3b. For comparison, the CV curves of the four materials in the voltage range of 2.0–4.2 V are shown in Fig. S7. For all samples, due to the initial interfacial reactions of the cells [36], the initial CV curves differ slightly with the subsequent cycles. However, their CV curves in the 2nd and 3rd cycles overlap well with each other, suggesting the highly reversible electrochemical reactions, which agrees well with their superior cycling stability in 2.0–4.2 V (Fig. S6b). The improvement in cycling perfor mance in the narrow voltage range can be attributed to the following two factors; (1) disappearance of detrimental phase transitions occur ring above 4.2 V and (2) suppression of oxygen activity and of the interfacial side reactions between electrode and electrolyte [9]. This result also suggests that for layered NaMO2 materials, the cycling per formance can be adjusted via controlling the cut-off voltage, providing useful information for cell design and operation. Furthermore, the local structure and electronic environment of nickel and manganese ions on the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 were investigated by ex situ XAS, the ac cording X-ray absorption near edge structure (XANES) spectra are pre sented in Fig. 5. Two main features are observed, i.e. the weak pre-edge peaks at lower photon energies (the electric dipole-forbidden transition of 1s electrons to unoccupied 3d orbitals) and the main edge absorption peaks at higher photon energies (the electron transition from 1s to 4p orbitals). For Na0.67Ni0.33Mn0.67O2 in the pristine state, the pre-edge and
1
in the
main edge are close to those of the NiO and MnO2 reference compounds. Therefore, the oxidation states of Ni and Mn ions are predominately 2 þ and 4þ, respectively [37]. Upon charging to 4.2 V, the Na0.67Ni0.33Mn0.67O2 displays an energy shift of ~3 eV at Ni K-edge (Fig. 5a), reflecting the increasing of Ni oxidation state. With further charging to 4.5 V, the Ni K-edge shows no clear shifts for both electrodes. Obviously, Ni ions are electrochemically active upon 4.2 V during charging and become inactive when the charging voltage goes even higher. When being discharged to 2.0 V, Ni K-edge moves back to its original position, indicating a highly reversible redox reaction of Ni ions. No further change can be distinguished between 2.0 V and 1.5 V, sug gesting that Ni ions are not responsible for the long plateau below 2.0 V in discharge. For the Mn K-edge of Na0.67Ni0.33Mn0.67O2 (Fig. 5b), no obvious shift can be observed during the charge process due to the predominate existence of Mn4þ. During discharge, the Mn K-edge shifts continuously to lower energy region, indicating the decrease of the Mn oxidation state. As for the Ni and Mn K-edge of the Na0.67Ni0.33Mn0.66Sn0.01O2 (displayed in Fig. 5c and d), similar change tendency is observed, indicating that the insertion of Sn in the crystal structure does not affect the charge compensation mechanism when compared to that of pristine Na0.67Ni0.33Mn0.67O2. In summary, the XANES results suggest for the additional voltage plateau between 4.2 and 4.5 V, that neither of Ni or Mn ions are involved in the redox reaction, but that O2 |O22 redox reactions seem to be the only possible option to balance the charge transfer. To confirm this hypothesis, ex situ O 1s XPS spectra of the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 materials at 4.2 V and 4.5 V were measured to detect possible peroxo-like species. XPS has been proven to be a powerful technique to evaluate the oxygen redox chemistry in layered oxides for LIBs and SIBs [13,38,39]. For both materials, at 4.2 V, only the 5
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 5. Ex situ XAS analysis of the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 electrodes at states of pristine, charged to 4.2 V, charged to 4.5 V, discharged to 2.0 V and discharged to 1.5 V. Normalized XANES spectra of Na0.67Ni0.33Mn0.67O2 (a) at Ni K-edge and (b) at Mn K-edge. Normalized XANES spectra of Na0.67Ni0.33Mn0.66Sn0.01O2 (c) at Ni K-edge and (d) at Mn K-edge.
lattice oxygen O2 (binding energy of ~529.5 eV) is shown in Fig. S8, while in the fully charged state of 4.5 V, a new peak at ~530.5 eV ap pears, which can be assigned to O /O22 according to previous literature [14,39,40], suggesting that the oxygen redox reactions participate in the electrochemical process and provide the capacity at the voltage plateau higher than 4.2 V. The ex situ XAS and XPS results are in good accordance with our previous conclusion about the involvement of oxygen activity, sug gesting that oxygen ions are partly extracted from the crystal lattice and are participating in the redox reactions [15]. It has been reported that the oxygen redox process has poor kinetics due to the requirement for coupled diffusion of oxygen ions, transition metals and vacancies [33, 41]. Since the substituted samples have smaller particles compared to un-substituted Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2 could exhibit enhanced oxygen activity due to the shorter diffusion way, thus, the discharge capacity is increased to 245 mA h g 1. However, in the samples with higher Sn amount (3% and 5%), the much higher bonding energy of Sn–O (△HfSn-O ¼ 548 (21) kJ mol 1 > △HfMn-O ¼ 402 (34) kJ mol 1 or △HfNi-O ¼ 391.6 (38) kJ mol 1) can suppress the release of oxygen atoms from the lattice [17,18]. Thus, smaller reversible capac and ities are obtained for Na0.67Ni0.33Mn0.64Sn0.03O2 Na0.67Ni0.33Mn0.62Sn0.05O2. In addition, in the relatively narrow voltage range of 2.0–4.2 V, the oxygen redox reactions cannot be activated, thus with increasing amount of Sn, the capacities decrease due to inert and heavy Sn4þ ions. Finally, the thermal stabilities of Na0.67Ni0.33Mn0.67O2 and Na0.67Ni0.33Mn0.66Sn0.01O2 materials are compared. For application of electrode materials, safety is one of the most decisive properties [42]. It
Fig. 6. DSC profiles of Na0.67Ni0.33Mn0.67O2 and Na0.67Ni0.33Mn0.66Sn0.01O2 at fully charged state (4.5 V) in the temperature range between 150 � C and 300 � C measured with a heating rate of 5 K min 1.
is known that substitution of materials can change the thermal stability of electrodes [17,43]. DSC measurements were carried out to assess the thermal behavior of the two electrodes after charge to 4.5 V. To simulate thermal abuse under realistic conditions, the electrode samples were immersed in the electrolytes before the DSC measurement. The DSC heat flow results in Fig. 6 illustrate, that both samples display an exothermal decomposition peak at 245 � C. However, un-substituted 6
Journal of Power Sources xxx (xxxx) xxx
J. Li et al.
Na0.67Ni0.33Mn0.67O2 shows an additional exothermal peak at 273 � C, reflecting a different thermal decomposition reaction in comparison to Na0.67Ni0.33Mn0.66Sn0.01O2. Though the onset temperature of Na0.67Ni0.33Mn0.66Sn0.01O2 and Na0.67Ni0.33Mn0.67O2 is the same, the output energy (1600 J g 1, calculated from the peak areas) of Na0.67Ni0.33Mn0.66Sn0.01O2 is higher than that of the Na0.67 Ni0.33Mn0.67O2 (760 J g 1 and 720 J g 1), revealing a worse thermal stability of the Na0.67Ni0.33Mn0.66Sn0.01O2 than that of Na0.67Ni0.33Mn0.67O2. At fully charged state, the Na0.67Ni0.33Mn0.66 Sn0.01O2 with enhanced oxygen activity results in strong side reactions between the active oxygen species and the electrolyte, and exhibits a decreased thermal stability. A similar phenomenon was also reported by other researchers [44].
[2] [3]
[4]
4. Conclusions
[5] [6]
A series of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) with mixed P2 and P3 phases have been successfully synthesized. With Sn substitution, the crystal structure evolves from P2 to P3, and the particle size decreases. The Sn substitution does not affect the charge compen sation mechanism of these materials; the long voltage plateau higher than 4.2 V is attributed to oxygen redox reactions. In addition, the introduction of 1% Sn is proven to be favorable for the enhancement of oxygen redox activity due to the smaller particle size, resulting in an impressive discharge capacity of 245 mA h g 1, which is the highest reported value for Na0.67Ni0.33Mn0.67O2 materials so far. As the Sn amount continuously increases to 3% or even to 5%, the much stronger Sn–O dissociation energy leads to the suppression of oxygen redox re actions. In the voltage range of 2.0–4.2 V, these aforesaid phenomena cannot be observed as no oxygen activity takes place. A C-rate com parison in the voltage ranges of 1.5–4.5 V vs. 2.0–4.2 V confirms that oxygen activity depends largely on the applied current density, and a high current density is detrimental to the release of oxygen activity in Na0.67Ni0.33Mn0.67Sn0.01O2 owing to its rate limitation. Besides, Na0.67Ni0.33Mn0.66Sn0.01O2 exhibits smaller thermal stability than Na0.67Ni0.33Mn0.67O2 in the fully charged state, suggesting that safety properties should be carefully considered regarding the application of oxygen activity. In summary, this work proves that with Sn substitution, the related oxygen activity of Na0.67Ni0.33Mn0.67O2 can be affected, introducing a new insight into the design of sodium layered oxides with large reversible capacities.
[7] [8] [9] [10] [11] [12]
[13]
[14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Declaration of competing interest
[25]
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
[26] [27] [28]
Acknowledgements
[29] [30]
This work is sponsored by the China Scholarship Council and the German Research Foundation (DFG, project Li 2916/2-1). The alloca tions of beamtime at P64, DESY, Hamburg, Germany is also gratefully acknowledged. We also appreciate the help from Dr. Uta Rodehorst for the discussion on XPS analysis results.
[31]
Appendix A. Supplementary data
[32] [33] [34] [35]
Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227554.
[36]
References [1] [a] M. Winter, B. Barnett, K. Xu, Chem. Rev. 118 (2018) 11433; [b] R. Wagner, N. Preschitschek, S. Passerini, J. Leker, M. Winter, J. Appl. Electrochem. 43 (2013) 481;
[37]
7
[c] J. Betz, G. Bieker, P. Meister, T. Placke, M. Winter, R. Schmuch, Adv. Energy Mater. 9 (2019) 1803170. N. Yabuuchi, K. Kubota, M. Dahbi, S. Komaba, Chem. Rev. 114 (2014) 11636. [a] J. Wang, B. Qiu, X. He, T. Risthaus, H. Liu, M.C. Stan, S. Schulze, Y. Xia, Z. Liu, M. Winter, Chem. Mater. 27 (2015) 4374; [b] X. He, J. Wang, B. Qiu, E. Paillard, C. Ma, X. Cao, H. Liu, M.C. Stan, H. Liu, T. Gallash, Y.S. Meng, J. Li, Nano Energy 27 (2016) 602. [a] N. Yabuuchi, M. Kajiyama, J. Iwatate, H. Nishikawa, S. Hitomi, R. Okuyama, R. Usui, Y. Yamada, S. Komaba, Nat. Mater. 11 (2012) 512; [b] Y. Wang, R. Xiao, Y. Hu, M. Avdeev, L. Chen, Nat. Commun. 6 (2015) 6954; [c] N. Bucher, S. Hartung, J.B. Franklin, A.M. Wise, L.Y. Lim, H. Chen, J.N. Weker, M.F. Toney, M. Srinivasan, Chem. Mater. 28 (2016) 2041; [d] S. Hartung, N. Bucher, V.S. Nair, C.Y. Ling, Y. Wang, H.E. Hoster, M. Srinivasan, ChemPhysChem 15 (2014) 2121; [e] J. Li, J. Wang, X. He, L. Zhang, A. Senyshyn, B. Yan, M. Muehlbauer, X. Cao, B. Vortmann-Westhoven, V. Kraft, H. Liu, C. Luerenbaum, G. Schumacher, E. Paillard, M. Winter, J. Li, J. Power Sources 416 (2019) 184; [f] T. Risthaus, L. Chen, J. Wang, J. Li, D. Zhou, L. Zhang, D. Ning, X. Cao, X. Zhang, G. Schumacher, M. Winter, E. Paillard, J. Li, Chem. Mater. 31 (2019) 5376. Z. Lu, J. Dahn, J. Electrochem. Soc. 148 (2001) A1225. D. Buchholz, L.G. Chagas, C. Vaalma, L. Wu, S. Passerini, J. Mater. Chem. 2 (2014) 13415. C. Delmas, C. Fouassier, P. Hagenmuller, Phys. BþC 99 (1980) 81. J. Hwang, S. Myung, Y. Sun, Chem. Soc. Rev. 46 (2017) 3529. K. Kubota, S. Kumakura, Y. Yoda, K. Kuroki, S. Komaba, Adv. Energy Mater. 8 (2018) 1703415. Z. Lu, J.R. Dahn, Chem. Mater. 13 (2001) 1252. Z. Lu, J. Dahn, J. Electrochem. Soc. 148 (2001) A710. [a] C. Ma, J. Alvarado, J. Xu, R.J. Cl�ement, M. Kodur, W. Tong, C.P. Grey, Y. S. Meng, J. Am. Chem. Soc. 139 (2017) 4835; [b] Y. Qiao, S. Guo, K. Zhu, P. Liu, X. Li, K. Jiang, C. Sun, M. Chen, H. Zhou, Energy Environ. Sci. 11 (2018) 299. U. Maitra, R.A. House, J.W. Somerville, N. Tapia-Ruiz, J.G. Lozano, N. Guerrini, R. Hao, K. Luo, L. Jin, M.A. P� erez-Osorio, F. Massel, D.M. Pickup, S. Ramos, X. Lu, D.E. McNally, A.V. Chadwick, F. Giustino, T. Schmitt, L.C. Duda, M.R. Roberts, P. G. Bruce, Nat. Chem. 10 (2018) 288. X. Rong, J. Liu, E. Hu, Y. Liu, Y. Wang, J. Wu, X. Yu, K. Page, Y.-S. Hu, W. Yang, H. Li, X.-Q. Yang, L. Chen, X. Huang, Joule 2 (2018) 125. T. Risthaus, D. Zhou, X. Cao, X. He, B. Qiu, J. Wang, L. Zhang, Z. Liu, E. Paillard, G. Schumacher, M. Winter, J. Li, J. Power Sources 395 (2018) 16. D.H. Lee, J. Xu, Y.S. Meng, Phys. Chem. Chem. Phys. 15 (2013) 3304. Q. Qiao, L. Qin, G. Li, Y. Wang, X. Gao, J. Mater. Chem. 3 (2015) 17627. Z. Deng, A. Manthiram, J. Phys. Chem. C 115 (2011) 7097. A. Konarov, J.H. Jo, J.U. Choi, Z. Bakenov, H. Yashiro, J. Kim, S.T. Myung, Nano Energy 59 (2019) 197. Y. Zhao, M. Xia, X. Hu, Z. Zhao, Y. Wang, Z. Lv, Electrochim. Acta 174 (2015) 1167. X. Rong, F. Gao, F. Ding, Y. Lu, K. Yang, H. Li, X. Huang, L. Chen, Y.-S. Hu, J. Mater. Sci. Technol. 35 (2019) 1250. R. N€ olle, K. Beltrop, F. Holtstiege, J. Kasnatscheew, T. Placke, M. Winter, Mater. Today (2019). T. Risthaus, L. Chen, J. Wang, J. Li, D. Zhou, L. Zhang, D. Ning, X. Cao, X. Zhang, G. Schumacher, M. Winter, E. Paillard, J. Li, Chem. Mater. (2019). J. Cabana, N.A. Chernova, J. Xiao, M. Roppolo, K.A. Aldi, M.S. Whittingham, C. P. Grey, Inorg. Chem. 52 (2013) 8540. [a] S.Y. Lee, J.H. Kim, Y.C. Kang, Electrochim. Acta 225 (2017) 86; [b] X. Wu, J. Guo, D. Wang, G. Zhong, M.J. McDonald, Y. Yang, J. Power Sources 281 (2015) 18; [c] L. Zheng, J. Li, M.N. Obrovac, Chem. Mater. 29 (2017) 1623. W. Zhao, H. Kirie, A. Tanaka, M. Unno, S. Yamamoto, H. Noguchi, Mater. Lett. 135 (2014) 131. Y. Shi, Y. Liu, H. Yang, C. Nie, R. Jin, J. Li, Phys. Rev. B 70 (2004) 052502. G. Singh, J.M. L� opez del Amo, M. Galceran, S. P� erez-Villar, T. Rojo, J. Mater. Chem. 3 (2015) 6954. K. Singh, B. Kirubasankar, S. Angaiah, Ionics 23 (2017) 731. P. Wang, Y. You, Y. Yin, Y. Wang, L. Wan, L. Gu, Y. Guo, Angew. Chem. Int. Ed. 55 (2016) 7445. [a] M. Sendova-Vassileva, R. Stoyanova, D. Carlier, M. Yoncheva, E. Zhecheva, C. Delmas, Raman spectroscopy study on Na2/3Mn1-xFexO2 oxides, in: Presented at Advances in Science and Technology, 2010; [b] M.A. Khan, D. Han, G. Lee, Y.-I. Kim, Y.-M. Kang, J. Alloy. Comp. 771 (2019) 987; [c] R. Stoyanova, D. Carlier, M. Sendova-Vassileva, M. Yoncheva, E. Zhecheva, D. Nihtianova, C. Delmas, J. Solid State Chem. 183 (2010) 1372. J. Qu, W. Wang, Y. Chen, G. Li, X. Li, Phys. Rev. B 73 (2006) 092518. A. van Bommel, J. Dahn, Electrochem. Solid State Lett. 13 (2010) A62. N. Yabuuchi, K. Kubota, Y. Aoki, S. Komaba, J. Phys. Chem. C 120 (2016) 875. Q. Wang, J. Meng, X. Yue, Q. Qiu, Y. Song, X. Wu, Z. Fu, Y. Xia, Z. Shadike, J. Wu, J. Am. Chem. Soc. 141 (2018) 840. [a] M.H. Han, E. Gonzalo, N. Sharma, J.M. L� opez del Amo, M. Armand, M. Avdeev, J.J. Saiz Garitaonandia, T. Rojo, Chem. Mater. 28 (2015) 106; [b] I. Hasa, D. Buchholz, S. Passerini, B. Scrosati, J. Hassoun, Adv. Energy Mater. 4 (2014) 1400083. X. Wu, G. Xu, G. Zhong, Z. Gong, M.J. McDonald, S. Zheng, R. Fu, Z. Chen, K. Amine, Y. Yang, ACS Appl. Mater. Interfaces 8 (2016) 22227.
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
[38] [a] G. Assat, D. Foix, C. Delacourt, A. Iadecola, R. Dedryv�ere, J.M. Tarascon, Nat. Commun. 8 (2017) 2219; [b] G. Assat, A. Iadecola, D. Foix, R. Dedryv� ere, J. Tarascon, ACS Energy Lett. 3 (2018) 2721. [39] [a] E. McCalla, A.M. Abakumov, M. Sauban� ere, D. Foix, E.J. Berg, G. Rousse, M. Doublet, D. Gonbeau, P. Nov� ak, G. Van Tendeloo, Science 350 (2015) 1516; [b] L. Li, H. Wang, W. Han, H. Guo, A. Hoser, Y. Chai, X. Liu, J. Electrochem. Soc. 165 (2018) A3854. [40] M. Sathiya, G. Rousse, K. Ramesha, C.P. Laisa, H. Vezin, M.T. Sougrati, M. L. Doublet, D. Foix, D. Gonbeau, W. Walker, A.S. Prakash, M. Ben Hassine, L. Dupont, J.M. Tarascon, Nat. Mater. 12 (2013) 827.
[41] N. Yabuuchi, K. Yoshii, S. Myung, I. Nakai, S. Komaba, J. Am. Chem. Soc. 133 (2011) 4404. [42] A. Friesen, F. Horsthemke, X. M€ onnighoff, G. Brunklaus, R. Krafft, M. B€ orner, T. Risthaus, M. Winter, F.M. Schappacher, J. Power Sources 334 (2016) 1. [43] M. Eilers-Rethwisch, S. Hildebrand, M. Evertz, L. Ibing, T. Dagger, M. Winter, F. M. Schappacher, J. Power Sources 397 (2018) 68. [44] X. Zhang, J. Shi, J. Liang, Y. Yin, J. Zhang, X. Yu, Y. Guo, Adv. Mater. 30 (2018) 1801751.
8
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
The effect of Sn substitution on the structure and oxygen activity of Na0.67Ni0.33Mn0.67O2 cathode materials for sodium ion batteries Jinke Li a, 1, Tim Risthaus b, 1, Jun Wang b, **, Dong Zhou b, Xin He a, ***, Niloofar Ehteshami a, Vadim Murzin c, Alex Friesen b, Haidong Liu b, Xu Hou a, Marcel Diehl b, Elie Paillard a, Martin Winter a, b, Jie Li a, * a b c
Helmholtz-Institute Muenster (HI MS), IEK-12, Forschungszentrum Juelich GmbH, Corrensstr. 46, Muenster, 48149, Germany MEET Battery Research Center, Institute of Physical Chemistry, University of Muenster, Corrensstr. 46, Muenster, 48149, Germany Deutsches Elektronen-Synchrotron, Notkestr. 85, Hamburg, 22607, Germany
H I G H L I G H T S
� P2/P3 Na0.67Ni0.33Mn0.67-xSnxO2 cathode materials for SIBs are proposed. � Na0.67Ni0.33Mn0.66Sn0.01O2 exhibits high reversible capacities in 1.5–4.5 V. � Ex-situ XAS and XPS are applied to characterize the involved redox reactions. � Oxygen activity in the family of the Na0.67Ni0.33Mn0.67O2 materials. A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium ion batteries Cathode Sodium nickel manganese oxides Sn substitution Oxygen activity
A series of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) materials with mixed P2/P3 phases are synthesized with a conventional solid-state reaction method and investigated as cathode materials for sodium ion batteries. The effects of Sn substitution on the structure and electrochemical performance of the Na0.67Ni0.33Mn0.67O2 are systematically investigated. The substituted samples show smaller particle sizes compared to the pristine one and the P2:P3 phase ratio highly depends on the substitution amount. The best electrochemical performance is ob tained by Na0.67Ni0.33Mn0.66Sn0.01O2, and it delivers a discharge capacity of 245 mA h g 1 in 1.5–4.5 V (vs. Na| Naþ), which is the highest result for Na0.67Ni0.33Mn0.67O2 materials reported so far. The ex situ X-ray absorption spectroscopy and X-ray photoelectron spectroscopy measurements reveal that the oxygen ions participate in the redox reactions within the wide voltage range of 1.5–4.5 V. The increased capacity can be attributed to the smaller particle size, which results in more oxygen activity and then higher capacity.
1. Introduction Sodium ion batteries (SIBs) are attracting increasing interest as cost efficient and sustainable alternatives to the prosperous lithium ion batteries (LIBs) [1], due to the vast availability of Na resources and similar electrochemical properties between SIBs and LIBs, especially for large-scale energy storage systems [2,3]. Among reported cathode can didates, sodium layered oxides (NaMO2, M ¼ Co, Ni, Mn, Cr, Fe, V etc.) have obtained marvelous attentions in light of their high capacities and
easy synthesis route [4–6]. Generally, they are categorized into two main groups in view of structure, namely P2 and O3 type, where ‘P’ and ‘O’ refer to the occupation sites of Na ions, i.e. prismatic and octahedral, respectively, and the following digit means the number of Na layers in each repeated unit [7]. Since the standard electrochemical potential of Na|Naþ ( 2.71 V vs. SHE) is higher than that of Li|Liþ ( 3.04 V vs. SHE), the energy density of SIBs is usually lower than that of LIBs in similar cell configuration. Therefore, the development of NaxMO2 with high output voltage plateau and discharge capacity plays a key role for
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (J. Wang), [email protected] (X. He), [email protected] (J. Li). 1 These authors contributed equally to this work. https://doi.org/10.1016/j.jpowsour.2019.227554 Received 1 October 2019; Received in revised form 15 November 2019; Accepted 30 November 2019 0378-7753/© 2019 Published by Elsevier B.V.
Please cite this article as: Jinke Li, Journal of Power Sources, https://doi.org/10.1016/j.jpowsour.2019.227554
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
the practical application of SIBs [2,8]. Among various NaxMO2 materials, Na0.67Ni0.33Mn0.67O2 possesses high theoretical capacity (173 mA h g 1) and high working potential associated with the redox potential of Ni2þ|Ni4þ (>3.0 V vs. Na|Naþ), in addition to its high tolerance vs. the storage environment [9,10]. Consequently, it has attracted increasing attentions as a promising cathode material since its first exploration by Dahn’s group [5,11]. In addition to the conventional redox reactions from transition metals, the anionic redox activity in NaxMO2, namely the oxygen redox reaction, has also been explored intensively recently, and is proposed as an effective route to get cathode material with large reversible capacity [12–15]. The oxygen redox reaction of Na0.67Ni0.33Mn0.67O2 has been proven by both DFT calculation and experiments [15,16]. For example, our previous work confirmed that the excessive discharge capacity (vs. the theoretical capacity) was attributed to the oxygen activity when the Na0.67Ni0.33Mn0.67O2|Na cell was cycled in a wide voltage range of 1.5–4.5 V [15]. It is believed that the substitution of cation ions can affect the metal-oxygen covalence of the materials, thus the oxygen behaviors during charge and discharge process can be altered [17–19]. In this case, substitution of Ni or Mn in Na0.67Ni0.33Mn0.67O2 will potentially affect its oxygen activity. Sn, as non-toxic element, with a high bond dissociation energy to oxygen (Sn–O; ΔHf298 ¼ 548 (21) kJ mol 1), has been investigated as the substitution element in Li-rich cathode materials, and the oxygen activities at high voltage plateau are affected consequently [17,20]. However, it is rarely reported as the substitution element in the cathode materials with oxygen activity in SIBs. Rong et al. confirmed that, the Sn substitution in Na0.67Ni0.33Mn0.67O2 can lead to the formation of O3 phase, the increased working voltage and smooth voltage profile in the voltage range of 2.5–4.1 V [21]. However, due to the narrow voltage range applied, the influence of Sn substitution on the oxygen activity of Na0.67Ni0.33Mn0.67O2 was not discussed in that paper. Herein, a series of P2/P3 Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) materials are prepared by a combination of co-precipitation and solid-state reaction methods. Among these four active materials, the Na0.67Ni0.33Mn0.66Sn0.01O2 delivers an initial charge capacity of 166 mA h g 1 and a high discharge capacity of 245 mA h g 1 within a wide voltage range of 1.5–4.5 V. While in a narrower operating voltage range (2.0–4.2 V), the results show that Sn substitution has no influence on the improvement of capacity. Additionally, the effects of Sn substitution on the redox chemistry and safety property are also explored.
Fig. 1. XRD patterns of Na0.67Ni0.33Mn0.67O2 (black), Na0.67Ni0.33Mn0.66Sn0.01O2 (yellow), Na0.67Ni0.33Mn0.64Sn0.03O2 (green) and Na0.67Ni0.33Mn0.62Sn0.05O2 (blue). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
equipped with a 532 nm argon-ion laser. The X-ray absorption spec troscopy (XAS) measurements were carried out at the P64 beamline of the Deutsches Elektronen-Synchrotron, Hamburg, Germany, using Si (111) double-crystal monochromator (FMB Oxford). X-ray photoelec tron spectroscopy (XPS) was performed with an Axis Ultra DLD (Kratos) by using monochromatic Al-Kα radiation and the calibration of the spectra was done with respect to adventitious carbon (284.5 eV). Thermogravimetric analysis (TGA) of the precursors before annealing was performed under air atmosphere at a heating rate of 5 K min 1 from room temperature to 800 � C with a TGA Q5000 IR devices from TA Instruments. Differential scanning calorimetry (DSC) measurements were realized by using a DSC Q2000 device from TA Instruments (USA; V24.11 build 124). All the electrode samples were heated to 300 � C at 5 K min 1 in high-pressure capsules. The cycled electrode samples were disassembled and then immersed in 1 M NaPF6 in EC/DMC (1/1, vol.) with 1 wt% FEC electrolyte. 2.3. Electrochemical characterization
2. Experimental
The electrochemical behaviors were analyzed with R2032-type coin cells in two electrode configuration [22]. The electrode consists of active material, conductive carbon (carbon black, Super C65, Imerys) and poly (vinylidene difluoride) (PVdF) binder (Kynar FLEX 761A, Arkema Group) at a weight ratio of 8:1:1. 1-methyl-2-pyrrolidinone (NMP, ACROS Organics) was used as solvent during the slurry preparation and the well mixed slurry was casted on the Al foil current collector. After being dried, electrode tapes were punched into Ø13 mm discs with a mass loading of �2 mg cm 2. All cells were assembled in an Ar-filled glove box with 1 M NaPF6 dissolved in 1:1 vol ratio of ethylene car bonate (EC) and dimethyl carbonate (DMC), together with 1 wt% fluo roethylene carbonate (FEC) as electrolyte. Sodium metal (99.8%, Acros Organics) and glass fiber (GF/D, Whatman) were used as anode and seperator, respectively. The cells were cycled galvanostatically at different C rates (1 C ¼ 200 mA g 1) between 1.5 and 4.5 V or 2.0–4.2 V at 20 � C using Maccor series 4000 battery tester. Cyclic voltammetry (CV) measurements were applied on a VSP electrochemical workstation (Bio-logic) with a sweep rate of 0.1 mV s 1.
2.1. Materials synthesis All chemicals were purchased from Sigma-Aldrich and used without any further purification. Stoichiometric amounts of NiSO4⋅6H2O, MnSO4⋅H2O and SnCl4⋅5H2O were dissolved in distilled water, then the solution was slowly dropped into a 1 M NaOH aqueous solution, after wards, the slurry was continuously stirred at 60 � C for 12 h in Ar at mosphere. The obtained precipitate was filtered, washed with distilled water and dried in air at 100 � C. At the end, the precipitate was thor oughly mixed with Na2CO3 by ball milling and calcined at 500 � C for 5 h, then 800 � C for 12 h in air to obtain final products. 2.2. Materials characterization Inductively coupled plasma optical emission spectroscopy (ICP-OES) was applied on the Spectro ARCOS instrument for the element deter mination. Powder X-ray diffraction (XRD) measurements were per formed on a Bruker D8 advance diffractometer with Cu kα radiation (λ ¼ 1.5418 Å) in the 2θ range from 10� to 90� . Scanning electron micro scopy (SEM) images and energy dispersive X-ray analysis (EDX) were recorded on the EVO MA 10 microscope (Zeiss). Raman spectra were measured using a Bruker SENTERRA dispersive Raman microscope,
3. Results and discussions Initially, the composition of all samples was determined by ICP-OES analysis. The results (Table S1) reflect that the molar ratio of Ni, Mn and Sn in the four samples agrees with the intended value. The TGA-DSC 2
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 2. SEM images of (a) Na0.67Ni0.33Mn0.67O2, (b) Na0.67Ni0.33Mn0.66Sn0.01O2, (c) Na0.67Ni0.33Mn0.64Sn0.03O2 and (d) Na0.67Ni0.33Mn0.62Sn0.05O2.
curves of the pristine Na0.67Ni0.33Mn0.67O2 and 5% Sn substitution Na0.67Ni0.33Mn0.62Sn0.05O2 precursors are shown in Fig. S1. Three weight loss processes can be seen in the TGA curves, i.e., below 100 � C, 100–250 � C and 250–800 � C which can be attributed to the loss of adsorbed H2O on the surface of the precursors, the decomposition of the hydroxides and carbonate, and the formation of crystal phase, respec tively. Except the first process (<100 � C), the weight loss is nearly the same (18.5% and 18.8%) when the temperature is higher than 100 � C. The DSC curves show the similar trend and overlap each other in the high temperature range. The DSC and TGA results both indicate that, Sn substitution shows no obvious influence on the annealing process. XRD measurements were applied to examine the crystal structure of the four samples (Fig. 1). As displayed, the Bragg diffraction peaks of the four samples are well indexed with the mixture of a hexagonal P2 phase (JCPDS Nr. 54–0894 with space group P63/mmc) and a tetragonal P3 phase (JCPDS Nr. 54–0839 with space group R3m). To obtain the ratios between P2 and P3 phases, the collected XRD data of the four materials are refined with Rietveld refinements (Fig. S2 and Table S2). The Na0.67Ni0.33Mn0.67O2 structure exhibits in majority P2 phase, while the substituted materials exhibit the main characteristics of the P3 phase, reflecting that Sn substitution leads to a difference in crystal structure. It is reported that the structure of Ni, Mn based NaxMO2 depends largely on the employed pH value of the precipitation process [23] and the annealing temperature [6,24]. The high annealing temperature is favorable to P2 phase formation while at low temperature the formation of the P3 phase is preferred [25]. In this work, the applied 800 � C leads to the formation of a mix of P2 and P3 phases. And as Sn4þ is involved in the precipitation process, the pH value of the mixture solution tends to decrease owing to partial hydrolysis of Sn4þ ions, thus, the Sn substituted materials display the main feature of P3 phase. The lattice parameters (Table S3) show that with an increase in Sn substitution amount, for both the P2 and P3 phase, the a and c axes of the samples also increase gradually, which can be attributed to the differ ence of ionic radii since Sn4þ (0.71 Å) has larger ionic radius than Mn4þ (0.53 Å) and Ni2þ (0.69 Å) [26]. This good agreement between the lattice parameter and ionic radii suggests that Sn is successfully substituted in the crystal structure. Besides, to exclude the formation of SnO2 in the obtained Sn-substitution materials, the XRD measurement on the pure SnO2 was conducted (Fig. S3). When compared with the
XRD patterns of the Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0.01, 0.03, 0.05), it is concluded that no SnO2 is formed during synthesis, indicating that Sn is substituted into the crystal structure of the investigated materials. Raman analyses (Fig. S4) shows that for all four active materials, five bands are observed, namely 630, 586, 478, 380, 348 cm 1, which are associated with the A1g þ 3E2g þ E1g mode, similar to the other sodium layered oxides NaxCoO2 [27], Na0.67Fe0.5Mn0.5O2 [28] and Na0.67Alx Co1-xO2 [29]. Similar Raman spectra of Na0.67Ni0.33Mn0.67O2 were also obtained by Noguchi etc. [26] and Guo etc. [30]. Besides, no bonds corresponding to the Sn-contained impurities such as SnO2 can be seen. According to the previous reports, the A1g and E1g mode are assigned to the O vibrations and three E2g modes come from the Na and O vibrations [31,32]. With different amount of Sn, the intensities of all bands vary, and the band at 630 cm 1 representing the symmetric stretching of MnO6 octahedra gets stronger [31], indicating that the short-range structure of the materials are influenced by Sn substitution. This result also manifests the hypothesis that Sn is indeed substituted in the lattice structure instead of forming multi-phase mixtures. In materials science, ‘substitution’ is commonly used as the term describing the replacement of one element/ion by another one(s) in a given crystal structure. The amount of employed ‘impure’ element/ion often reaches level of several percent. The term ‘doping’ originates from semiconductor science, where the dopants are introduced as additives (not as substituent), in a very small amount, into pure intrinsic semi conductors to modify their electrical properties. For battery materials, ‘doping’ is, in some cases, also used to represent atomic/ionic substi tution when the amount of substituent is relatively low. Despite the nonconsistent use in battery R&D and related literature, the present work will use the term ‘substitution’ instead of ‘doping’ since the scientific approach of this work aims at substitution of a certain amount of Mn by Sn. The SEM images of the four materials are illustrated in Fig. 2. The pristine Na0.67Ni0.33Mn0.67O2 is composed of flake like primary particle. In contrast, the Sn-containing materials appear as much smaller and agglomerated spherical particles. The SEM results reveal that Sn sub stitution can significantly affect the particle shape and size of the ma terials. The EDX mapping was carried out on the 5% Sn substitution Na0.67Ni0.33Mn0.62Sn0.05O2 samples (displayed in Fig. S5) as represen tative, and the result presents a homogenous element distribution 3
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
discharged sequentially at 0.1, 0.2, 0.5, 1, 2, 5, 10 and 1 C each for five cycles. All the four materials exhibit obvious capacity fading at low current rates. Nevertheless, the highest capacities at each rate is still obtained from Na0.67Ni0.33Mn0.66Sn0.01O2. In detail, the values (average value for five cycles) are 224, 180, 145, 118, 103, 84 and 71 mA h g 1 at 0.1, 0.2, 0.5, 1, 2, 5 and 10 C, respectively. When the current density returns back to 1 C, ~96% capacity recovery is achieved. For comparison, the Na0.67Ni0.33Mn0.67-xSnxO2||Na (x ¼ 0, 0.01, 0.03 and 0.05) cells was also charged/discharged in the voltage range of 2.0–4.2 V (Fig. S6). Within the narrow voltage window of 2.0–4.2 V, the initial discharge capacities of all the electrodes decrease dramatically to less than 90 mA h g 1. However, the cycling stabilities are promising. After 50 cycles, 101%, 100%, 101% and 101% of the initial capacities are maintained with Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64Sn0.03O2 and Na0.67Ni0.33Mn0.62Sn0.05O2, respec tively. In addition, the materials also exhibit excellent rate capabilities (Fig. S6c). For example, at 10 C, Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33 Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64Sn0.03O2 and Na0.67Ni0.33Mn0.62 Sn0.05O2 maintain the discharge capacities of 78, 80, 77 and 72 mA h g 1, corresponding to 85%, 86%, 88% and 86% of their initial capacities at 0.1 C, respectively. As the current density returning to 1 C, the discharge capacities of 87, 87, 81 and 78 mA h g 1 are achieved with ~100% capacity retentions of their initial specific capacities at 1 C, suggesting good structural reversibility in the 2.0–4.2 V voltage range. Overall, in the voltage range of 1.5–4.5 V, the Na0.67Ni0.33Mn0.67xSnxO2 (x ¼ 0, 0.01, 0.03 and 0.05) materials show high a capacity of >220 mA h g 1 at 0.1 C. For the Na0.67Ni0.33Mn0.66Sn0.01O2, 245 mA h g 1 is the largest discharge capacity reported so far from the family of Na0.67Ni0.33Mn0.67O2 compounds. The value far exceeds the theoretical capacity of 172 mA h g 1 (considering only the redox reaction of Ni2þ| Ni4þ), which we attribute to the additional capacity contribution by oxygen redox [15]. Meanwhile, 0.01 Sn substitution leads to 8.4% ca pacity increase (from 226 mA h g 1 delivered by pristine sample to 245 mA h g 1), which is probably because that Sn substitution enhances the oxygen activity. However, it should be noted that an enhanced oxygen activity also leads to a relatively poor cycling performance. Though, the detrimental structural phase transition at high voltage with huge volume shrinkage or serious reactions between electrode (surface) and electro lyte would contribute to the performance fading [9]. In this case, the unsatisfactory reversibility of the oxygen redox reactions which are commonly coupled with oxygen loss [18], and the slow kinetics of the oxygen activity may also take responsibility in the fading [33,34]. Although the cycling performance and the rate capability of these ma terials can be largely enhanced when narrowing the voltage window is applied, this is achieved at the expense of ~30% of the total capacity (based on the value obtained at 50th cycle). Nevertheless, it is urgent to understand the charge compensation mechanisms of these materials upon Na extraction and insertion. Thus, CV measurements in different voltage ranges of 1.5–4.5 V and 2.0–4.2 V were performed on all four Na0.67Ni0.33Mn0.67-xSnxO2||Na cells. All electrodes for the CV measurement were selected with the same size and similar mass loading. The CV curves of the four materials in the voltage range of 1.5–4.5 V are displayed in Fig. 4. The un-substituted Na0.67Ni0.33Mn0.67O2 exhibits a series of complex current peaks. Generally speaking, the strong peaks higher than 4.2 V are due to phase transformation coupled with oxygen redox reactions, and the peaks between 2.5 and 4.0 V can be assigned to the redox reactions of Ni2þ| Ni4þ accompanied with Naþ/vacancy ordering occurring within the sodium layers. The redox peaks lower than 2.5 V refer to the Mn redox reactions [19,35]. In contrast to un-substituted Na0.67Ni0.33Mn0.67O2, the Sn substituted materials exhibit similar peak positions in the initial cycle, reflecting a similar reaction mechanism between them, regardless of the ratio of P2 and P3 phases. Na0.67Ni0.33Mn0.66Sn0.01O2 shows the strongest peaks in the voltage range of >4.2 V and <2.5 V. With higher Sn amounts (x ¼ 0.03 and 0.05), the intensity of these peaks decreases. Besides, it is notable that for all materials, the intensity of the peak
Fig. 3. (a) Initial charge and discharge curves, (b) cycling performance and (c) rate capabilities of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) in the voltage range of 1.5–4.5 V.
throughout the particle. The first charge and discharge curves of the Na0.67Ni0.33Mn0.671 xSnxO2||Na (x ¼ 0, 0.01, 0.03 and 0.05) cells at 0.1 C (1 C ¼ 200 mA g ) in the voltage range of 1.5–4.5 V are displayed in Fig. 3a. As illustrated, the initial charge and discharge capacities of the pristine sample are 161 mA h g 1 and 226 mA h g 1, respectively. In contrast, the Na0.67Ni0.33Mn0.66Sn0.01O2 (166 and 245 mA h g 1) and Na0.67Ni0.33Mn0.64Sn0.03O2 (162 and 230 mA h g 1) show higher charge and discharge capacities, while the sample with 5% Sn substitution delivers lower ones (155 and 220 mA h g 1). The cycling performance of the four materials at 1 C is compared in Fig. 3b. After 50 cycles at 1 C, the Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2, Na0.67Ni0.33Mn0.64 Sn0.03O2 and Na0.67Ni0.33Mn0.62Sn0.05O2 deliver discharge capacities of 122, 129, 117, 113 mA h g 1, respectively, which are associated with capacity retentions of 81%, 73%, 72% and 85%, respectively. Fig. 3c shows the rate performance of the cells which were charged and 4
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 4. Cyclic voltammograms of the Na0.67Ni0.33Mn0.67-xSnxO2||Na cell (a) x ¼ 0, (b) x ¼ 0.01, (c) x ¼ 0.03 and (d) x ¼ 0.05 at a scan rate of 0.1 mV s voltage range of 1.5–4.5 V.
higher than 4.2 V has a significantly positive correlation with that of the peak lower than 2.5 V, i.e., as the peaks >4.2 V become stronger, the peaks <2.5 V are also stronger, reflecting that the oxygen activity cor relates with the redox reaction of the Mn ions. Upon cycling, the in tensity of the peak (>4.2 V) is decreasing due to the poor reversibility of the reactions, which are responsible for the observed capacity fading in Fig. 3b. For comparison, the CV curves of the four materials in the voltage range of 2.0–4.2 V are shown in Fig. S7. For all samples, due to the initial interfacial reactions of the cells [36], the initial CV curves differ slightly with the subsequent cycles. However, their CV curves in the 2nd and 3rd cycles overlap well with each other, suggesting the highly reversible electrochemical reactions, which agrees well with their superior cycling stability in 2.0–4.2 V (Fig. S6b). The improvement in cycling perfor mance in the narrow voltage range can be attributed to the following two factors; (1) disappearance of detrimental phase transitions occur ring above 4.2 V and (2) suppression of oxygen activity and of the interfacial side reactions between electrode and electrolyte [9]. This result also suggests that for layered NaMO2 materials, the cycling per formance can be adjusted via controlling the cut-off voltage, providing useful information for cell design and operation. Furthermore, the local structure and electronic environment of nickel and manganese ions on the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 were investigated by ex situ XAS, the ac cording X-ray absorption near edge structure (XANES) spectra are pre sented in Fig. 5. Two main features are observed, i.e. the weak pre-edge peaks at lower photon energies (the electric dipole-forbidden transition of 1s electrons to unoccupied 3d orbitals) and the main edge absorption peaks at higher photon energies (the electron transition from 1s to 4p orbitals). For Na0.67Ni0.33Mn0.67O2 in the pristine state, the pre-edge and
1
in the
main edge are close to those of the NiO and MnO2 reference compounds. Therefore, the oxidation states of Ni and Mn ions are predominately 2 þ and 4þ, respectively [37]. Upon charging to 4.2 V, the Na0.67Ni0.33Mn0.67O2 displays an energy shift of ~3 eV at Ni K-edge (Fig. 5a), reflecting the increasing of Ni oxidation state. With further charging to 4.5 V, the Ni K-edge shows no clear shifts for both electrodes. Obviously, Ni ions are electrochemically active upon 4.2 V during charging and become inactive when the charging voltage goes even higher. When being discharged to 2.0 V, Ni K-edge moves back to its original position, indicating a highly reversible redox reaction of Ni ions. No further change can be distinguished between 2.0 V and 1.5 V, sug gesting that Ni ions are not responsible for the long plateau below 2.0 V in discharge. For the Mn K-edge of Na0.67Ni0.33Mn0.67O2 (Fig. 5b), no obvious shift can be observed during the charge process due to the predominate existence of Mn4þ. During discharge, the Mn K-edge shifts continuously to lower energy region, indicating the decrease of the Mn oxidation state. As for the Ni and Mn K-edge of the Na0.67Ni0.33Mn0.66Sn0.01O2 (displayed in Fig. 5c and d), similar change tendency is observed, indicating that the insertion of Sn in the crystal structure does not affect the charge compensation mechanism when compared to that of pristine Na0.67Ni0.33Mn0.67O2. In summary, the XANES results suggest for the additional voltage plateau between 4.2 and 4.5 V, that neither of Ni or Mn ions are involved in the redox reaction, but that O2 |O22 redox reactions seem to be the only possible option to balance the charge transfer. To confirm this hypothesis, ex situ O 1s XPS spectra of the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 materials at 4.2 V and 4.5 V were measured to detect possible peroxo-like species. XPS has been proven to be a powerful technique to evaluate the oxygen redox chemistry in layered oxides for LIBs and SIBs [13,38,39]. For both materials, at 4.2 V, only the 5
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
Fig. 5. Ex situ XAS analysis of the Na0.67Ni0.33Mn0.67O2 and the Na0.67Ni0.33Mn0.66Sn0.01O2 electrodes at states of pristine, charged to 4.2 V, charged to 4.5 V, discharged to 2.0 V and discharged to 1.5 V. Normalized XANES spectra of Na0.67Ni0.33Mn0.67O2 (a) at Ni K-edge and (b) at Mn K-edge. Normalized XANES spectra of Na0.67Ni0.33Mn0.66Sn0.01O2 (c) at Ni K-edge and (d) at Mn K-edge.
lattice oxygen O2 (binding energy of ~529.5 eV) is shown in Fig. S8, while in the fully charged state of 4.5 V, a new peak at ~530.5 eV ap pears, which can be assigned to O /O22 according to previous literature [14,39,40], suggesting that the oxygen redox reactions participate in the electrochemical process and provide the capacity at the voltage plateau higher than 4.2 V. The ex situ XAS and XPS results are in good accordance with our previous conclusion about the involvement of oxygen activity, sug gesting that oxygen ions are partly extracted from the crystal lattice and are participating in the redox reactions [15]. It has been reported that the oxygen redox process has poor kinetics due to the requirement for coupled diffusion of oxygen ions, transition metals and vacancies [33, 41]. Since the substituted samples have smaller particles compared to un-substituted Na0.67Ni0.33Mn0.67O2, Na0.67Ni0.33Mn0.66Sn0.01O2 could exhibit enhanced oxygen activity due to the shorter diffusion way, thus, the discharge capacity is increased to 245 mA h g 1. However, in the samples with higher Sn amount (3% and 5%), the much higher bonding energy of Sn–O (△HfSn-O ¼ 548 (21) kJ mol 1 > △HfMn-O ¼ 402 (34) kJ mol 1 or △HfNi-O ¼ 391.6 (38) kJ mol 1) can suppress the release of oxygen atoms from the lattice [17,18]. Thus, smaller reversible capac and ities are obtained for Na0.67Ni0.33Mn0.64Sn0.03O2 Na0.67Ni0.33Mn0.62Sn0.05O2. In addition, in the relatively narrow voltage range of 2.0–4.2 V, the oxygen redox reactions cannot be activated, thus with increasing amount of Sn, the capacities decrease due to inert and heavy Sn4þ ions. Finally, the thermal stabilities of Na0.67Ni0.33Mn0.67O2 and Na0.67Ni0.33Mn0.66Sn0.01O2 materials are compared. For application of electrode materials, safety is one of the most decisive properties [42]. It
Fig. 6. DSC profiles of Na0.67Ni0.33Mn0.67O2 and Na0.67Ni0.33Mn0.66Sn0.01O2 at fully charged state (4.5 V) in the temperature range between 150 � C and 300 � C measured with a heating rate of 5 K min 1.
is known that substitution of materials can change the thermal stability of electrodes [17,43]. DSC measurements were carried out to assess the thermal behavior of the two electrodes after charge to 4.5 V. To simulate thermal abuse under realistic conditions, the electrode samples were immersed in the electrolytes before the DSC measurement. The DSC heat flow results in Fig. 6 illustrate, that both samples display an exothermal decomposition peak at 245 � C. However, un-substituted 6
Journal of Power Sources xxx (xxxx) xxx
J. Li et al.
Na0.67Ni0.33Mn0.67O2 shows an additional exothermal peak at 273 � C, reflecting a different thermal decomposition reaction in comparison to Na0.67Ni0.33Mn0.66Sn0.01O2. Though the onset temperature of Na0.67Ni0.33Mn0.66Sn0.01O2 and Na0.67Ni0.33Mn0.67O2 is the same, the output energy (1600 J g 1, calculated from the peak areas) of Na0.67Ni0.33Mn0.66Sn0.01O2 is higher than that of the Na0.67 Ni0.33Mn0.67O2 (760 J g 1 and 720 J g 1), revealing a worse thermal stability of the Na0.67Ni0.33Mn0.66Sn0.01O2 than that of Na0.67Ni0.33Mn0.67O2. At fully charged state, the Na0.67Ni0.33Mn0.66 Sn0.01O2 with enhanced oxygen activity results in strong side reactions between the active oxygen species and the electrolyte, and exhibits a decreased thermal stability. A similar phenomenon was also reported by other researchers [44].
[2] [3]
[4]
4. Conclusions
[5] [6]
A series of Na0.67Ni0.33Mn0.67-xSnxO2 (x ¼ 0, 0.01, 0.03, 0.05) with mixed P2 and P3 phases have been successfully synthesized. With Sn substitution, the crystal structure evolves from P2 to P3, and the particle size decreases. The Sn substitution does not affect the charge compen sation mechanism of these materials; the long voltage plateau higher than 4.2 V is attributed to oxygen redox reactions. In addition, the introduction of 1% Sn is proven to be favorable for the enhancement of oxygen redox activity due to the smaller particle size, resulting in an impressive discharge capacity of 245 mA h g 1, which is the highest reported value for Na0.67Ni0.33Mn0.67O2 materials so far. As the Sn amount continuously increases to 3% or even to 5%, the much stronger Sn–O dissociation energy leads to the suppression of oxygen redox re actions. In the voltage range of 2.0–4.2 V, these aforesaid phenomena cannot be observed as no oxygen activity takes place. A C-rate com parison in the voltage ranges of 1.5–4.5 V vs. 2.0–4.2 V confirms that oxygen activity depends largely on the applied current density, and a high current density is detrimental to the release of oxygen activity in Na0.67Ni0.33Mn0.67Sn0.01O2 owing to its rate limitation. Besides, Na0.67Ni0.33Mn0.66Sn0.01O2 exhibits smaller thermal stability than Na0.67Ni0.33Mn0.67O2 in the fully charged state, suggesting that safety properties should be carefully considered regarding the application of oxygen activity. In summary, this work proves that with Sn substitution, the related oxygen activity of Na0.67Ni0.33Mn0.67O2 can be affected, introducing a new insight into the design of sodium layered oxides with large reversible capacities.
[7] [8] [9] [10] [11] [12]
[13]
[14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Declaration of competing interest
[25]
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
[26] [27] [28]
Acknowledgements
[29] [30]
This work is sponsored by the China Scholarship Council and the German Research Foundation (DFG, project Li 2916/2-1). The alloca tions of beamtime at P64, DESY, Hamburg, Germany is also gratefully acknowledged. We also appreciate the help from Dr. Uta Rodehorst for the discussion on XPS analysis results.
[31]
Appendix A. Supplementary data
[32] [33] [34] [35]
Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227554.
[36]
References [1] [a] M. Winter, B. Barnett, K. Xu, Chem. Rev. 118 (2018) 11433; [b] R. Wagner, N. Preschitschek, S. Passerini, J. Leker, M. Winter, J. Appl. Electrochem. 43 (2013) 481;
[37]
7
[c] J. Betz, G. Bieker, P. Meister, T. Placke, M. Winter, R. Schmuch, Adv. Energy Mater. 9 (2019) 1803170. N. Yabuuchi, K. Kubota, M. Dahbi, S. Komaba, Chem. Rev. 114 (2014) 11636. [a] J. Wang, B. Qiu, X. He, T. Risthaus, H. Liu, M.C. Stan, S. Schulze, Y. Xia, Z. Liu, M. Winter, Chem. Mater. 27 (2015) 4374; [b] X. He, J. Wang, B. Qiu, E. Paillard, C. Ma, X. Cao, H. Liu, M.C. Stan, H. Liu, T. Gallash, Y.S. Meng, J. Li, Nano Energy 27 (2016) 602. [a] N. Yabuuchi, M. Kajiyama, J. Iwatate, H. Nishikawa, S. Hitomi, R. Okuyama, R. Usui, Y. Yamada, S. Komaba, Nat. Mater. 11 (2012) 512; [b] Y. Wang, R. Xiao, Y. Hu, M. Avdeev, L. Chen, Nat. Commun. 6 (2015) 6954; [c] N. Bucher, S. Hartung, J.B. Franklin, A.M. Wise, L.Y. Lim, H. Chen, J.N. Weker, M.F. Toney, M. Srinivasan, Chem. Mater. 28 (2016) 2041; [d] S. Hartung, N. Bucher, V.S. Nair, C.Y. Ling, Y. Wang, H.E. Hoster, M. Srinivasan, ChemPhysChem 15 (2014) 2121; [e] J. Li, J. Wang, X. He, L. Zhang, A. Senyshyn, B. Yan, M. Muehlbauer, X. Cao, B. Vortmann-Westhoven, V. Kraft, H. Liu, C. Luerenbaum, G. Schumacher, E. Paillard, M. Winter, J. Li, J. Power Sources 416 (2019) 184; [f] T. Risthaus, L. Chen, J. Wang, J. Li, D. Zhou, L. Zhang, D. Ning, X. Cao, X. Zhang, G. Schumacher, M. Winter, E. Paillard, J. Li, Chem. Mater. 31 (2019) 5376. Z. Lu, J. Dahn, J. Electrochem. Soc. 148 (2001) A1225. D. Buchholz, L.G. Chagas, C. Vaalma, L. Wu, S. Passerini, J. Mater. Chem. 2 (2014) 13415. C. Delmas, C. Fouassier, P. Hagenmuller, Phys. BþC 99 (1980) 81. J. Hwang, S. Myung, Y. Sun, Chem. Soc. Rev. 46 (2017) 3529. K. Kubota, S. Kumakura, Y. Yoda, K. Kuroki, S. Komaba, Adv. Energy Mater. 8 (2018) 1703415. Z. Lu, J.R. Dahn, Chem. Mater. 13 (2001) 1252. Z. Lu, J. Dahn, J. Electrochem. Soc. 148 (2001) A710. [a] C. Ma, J. Alvarado, J. Xu, R.J. Cl�ement, M. Kodur, W. Tong, C.P. Grey, Y. S. Meng, J. Am. Chem. Soc. 139 (2017) 4835; [b] Y. Qiao, S. Guo, K. Zhu, P. Liu, X. Li, K. Jiang, C. Sun, M. Chen, H. Zhou, Energy Environ. Sci. 11 (2018) 299. U. Maitra, R.A. House, J.W. Somerville, N. Tapia-Ruiz, J.G. Lozano, N. Guerrini, R. Hao, K. Luo, L. Jin, M.A. P� erez-Osorio, F. Massel, D.M. Pickup, S. Ramos, X. Lu, D.E. McNally, A.V. Chadwick, F. Giustino, T. Schmitt, L.C. Duda, M.R. Roberts, P. G. Bruce, Nat. Chem. 10 (2018) 288. X. Rong, J. Liu, E. Hu, Y. Liu, Y. Wang, J. Wu, X. Yu, K. Page, Y.-S. Hu, W. Yang, H. Li, X.-Q. Yang, L. Chen, X. Huang, Joule 2 (2018) 125. T. Risthaus, D. Zhou, X. Cao, X. He, B. Qiu, J. Wang, L. Zhang, Z. Liu, E. Paillard, G. Schumacher, M. Winter, J. Li, J. Power Sources 395 (2018) 16. D.H. Lee, J. Xu, Y.S. Meng, Phys. Chem. Chem. Phys. 15 (2013) 3304. Q. Qiao, L. Qin, G. Li, Y. Wang, X. Gao, J. Mater. Chem. 3 (2015) 17627. Z. Deng, A. Manthiram, J. Phys. Chem. C 115 (2011) 7097. A. Konarov, J.H. Jo, J.U. Choi, Z. Bakenov, H. Yashiro, J. Kim, S.T. Myung, Nano Energy 59 (2019) 197. Y. Zhao, M. Xia, X. Hu, Z. Zhao, Y. Wang, Z. Lv, Electrochim. Acta 174 (2015) 1167. X. Rong, F. Gao, F. Ding, Y. Lu, K. Yang, H. Li, X. Huang, L. Chen, Y.-S. Hu, J. Mater. Sci. Technol. 35 (2019) 1250. R. N€ olle, K. Beltrop, F. Holtstiege, J. Kasnatscheew, T. Placke, M. Winter, Mater. Today (2019). T. Risthaus, L. Chen, J. Wang, J. Li, D. Zhou, L. Zhang, D. Ning, X. Cao, X. Zhang, G. Schumacher, M. Winter, E. Paillard, J. Li, Chem. Mater. (2019). J. Cabana, N.A. Chernova, J. Xiao, M. Roppolo, K.A. Aldi, M.S. Whittingham, C. P. Grey, Inorg. Chem. 52 (2013) 8540. [a] S.Y. Lee, J.H. Kim, Y.C. Kang, Electrochim. Acta 225 (2017) 86; [b] X. Wu, J. Guo, D. Wang, G. Zhong, M.J. McDonald, Y. Yang, J. Power Sources 281 (2015) 18; [c] L. Zheng, J. Li, M.N. Obrovac, Chem. Mater. 29 (2017) 1623. W. Zhao, H. Kirie, A. Tanaka, M. Unno, S. Yamamoto, H. Noguchi, Mater. Lett. 135 (2014) 131. Y. Shi, Y. Liu, H. Yang, C. Nie, R. Jin, J. Li, Phys. Rev. B 70 (2004) 052502. G. Singh, J.M. L� opez del Amo, M. Galceran, S. P� erez-Villar, T. Rojo, J. Mater. Chem. 3 (2015) 6954. K. Singh, B. Kirubasankar, S. Angaiah, Ionics 23 (2017) 731. P. Wang, Y. You, Y. Yin, Y. Wang, L. Wan, L. Gu, Y. Guo, Angew. Chem. Int. Ed. 55 (2016) 7445. [a] M. Sendova-Vassileva, R. Stoyanova, D. Carlier, M. Yoncheva, E. Zhecheva, C. Delmas, Raman spectroscopy study on Na2/3Mn1-xFexO2 oxides, in: Presented at Advances in Science and Technology, 2010; [b] M.A. Khan, D. Han, G. Lee, Y.-I. Kim, Y.-M. Kang, J. Alloy. Comp. 771 (2019) 987; [c] R. Stoyanova, D. Carlier, M. Sendova-Vassileva, M. Yoncheva, E. Zhecheva, D. Nihtianova, C. Delmas, J. Solid State Chem. 183 (2010) 1372. J. Qu, W. Wang, Y. Chen, G. Li, X. Li, Phys. Rev. B 73 (2006) 092518. A. van Bommel, J. Dahn, Electrochem. Solid State Lett. 13 (2010) A62. N. Yabuuchi, K. Kubota, Y. Aoki, S. Komaba, J. Phys. Chem. C 120 (2016) 875. Q. Wang, J. Meng, X. Yue, Q. Qiu, Y. Song, X. Wu, Z. Fu, Y. Xia, Z. Shadike, J. Wu, J. Am. Chem. Soc. 141 (2018) 840. [a] M.H. Han, E. Gonzalo, N. Sharma, J.M. L� opez del Amo, M. Armand, M. Avdeev, J.J. Saiz Garitaonandia, T. Rojo, Chem. Mater. 28 (2015) 106; [b] I. Hasa, D. Buchholz, S. Passerini, B. Scrosati, J. Hassoun, Adv. Energy Mater. 4 (2014) 1400083. X. Wu, G. Xu, G. Zhong, Z. Gong, M.J. McDonald, S. Zheng, R. Fu, Z. Chen, K. Amine, Y. Yang, ACS Appl. Mater. Interfaces 8 (2016) 22227.
J. Li et al.
Journal of Power Sources xxx (xxxx) xxx
[38] [a] G. Assat, D. Foix, C. Delacourt, A. Iadecola, R. Dedryv�ere, J.M. Tarascon, Nat. Commun. 8 (2017) 2219; [b] G. Assat, A. Iadecola, D. Foix, R. Dedryv� ere, J. Tarascon, ACS Energy Lett. 3 (2018) 2721. [39] [a] E. McCalla, A.M. Abakumov, M. Sauban� ere, D. Foix, E.J. Berg, G. Rousse, M. Doublet, D. Gonbeau, P. Nov� ak, G. Van Tendeloo, Science 350 (2015) 1516; [b] L. Li, H. Wang, W. Han, H. Guo, A. Hoser, Y. Chai, X. Liu, J. Electrochem. Soc. 165 (2018) A3854. [40] M. Sathiya, G. Rousse, K. Ramesha, C.P. Laisa, H. Vezin, M.T. Sougrati, M. L. Doublet, D. Foix, D. Gonbeau, W. Walker, A.S. Prakash, M. Ben Hassine, L. Dupont, J.M. Tarascon, Nat. Mater. 12 (2013) 827.
[41] N. Yabuuchi, K. Yoshii, S. Myung, I. Nakai, S. Komaba, J. Am. Chem. Soc. 133 (2011) 4404. [42] A. Friesen, F. Horsthemke, X. M€ onnighoff, G. Brunklaus, R. Krafft, M. B€ orner, T. Risthaus, M. Winter, F.M. Schappacher, J. Power Sources 334 (2016) 1. [43] M. Eilers-Rethwisch, S. Hildebrand, M. Evertz, L. Ibing, T. Dagger, M. Winter, F. M. Schappacher, J. Power Sources 397 (2018) 68. [44] X. Zhang, J. Shi, J. Liang, Y. Yin, J. Zhang, X. Yu, Y. Guo, Adv. Mater. 30 (2018) 1801751.
8