Crystal Structures and Electrochemical Performance of Air-Stable Na2

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Article

Crystal Structures and Electrochemical Performance of Air-Stable Na Ni CuMn O in Sodium Cells 2/3

1/3-x

x

2/3

2

Lituo Zheng, Jierui Li, and M. N. Obrovac Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b04769 • Publication Date (Web): 25 Jan 2017 Downloaded from http://pubs.acs.org on February 1, 2017

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Chemistry of Materials

Crystal Structures and Electrochemical Performance of Air-Stable Na2/3Ni1/3-xCuxMn2/3O2 in Sodium Cells Lituo Zheng†, Jierui Li†, and M.N. Obrovac†* †Department of Chemistry Dalhousie University, Halifax, Nova Scotia, Canada B3H 4R2 ABSTRACT: Sodium ion batteries have garnered significant research attention in recent years due to the rising demand for large-scale energy storage solutions as well as the high abundance of sodium. P2-type layered oxide materials have been identified as promising positive electrode materials for sodium ion batteries. Previously, P2-Na2/3Ni1/3Mn2/3O2 was shown to have a high operating voltage and high capacity but suffers from a step-like voltage curve and capacity loss during cycling, potentially due to its P2-O2 transition at high voltages. One strategy to improve cycling performance has been to dope Ni2+ with other 2+ cations, such as Zn2+ or Mg2+, which improved capacity retention, but significantly decreases reversible capacity, since these ions were not electrochemically active. Since Cu2+ has been shown to be electrochemically active, we replaced Ni2+ with Cu2+, resulting in air-stable Na2/3Ni1/3-xCuxMn2/3O2 (0 ≤ x ≤ 1/3). Both Ni2+/Ni4+ and Cu2+/Cu3+ participate in the redox reaction during cycling, capacity retention was greatly improved and phase changes were suppressed during cycling without sacrificing much capacity. The material retains a P2/OP4 structure even when cycled to high voltages. The doping strategy is a promising approach for the future development of positive electrode materials for sodium ion batteries.

■ INTRODUCTION In recent years, the demand for rechargeable lithium ion batteries (LIBs) has increased significantly, as larger batteries are utilized in electric vehicles and in grid-scale energy storage systems to complement renewable energy generation.1 Ongoing increases in demand could lead to rapid increases in LIB cost due to the relatively low abundance and limited geographic availability of lithium. As a result of this concern, research interest has grown in the development of alternatives to LIBs that comprise more earth-abundant materials. Sodium ion batteries (NIBs) have been identified as a promising technology due to sodium’s low cost, virtually unlimited supply as well as its electrochemical similarities to lithium.2 Layered sodium transition metal oxides (NaxMO2, M = one or more transition metals), a subset of positive electrode materials for NIBs, have been extensively studied with good results.3 However, one of the main problem with these layered oxides is that they generally have lower operating voltages compared to their lithium counterparts. For example, the average voltage of NaCoO2 vs. Na/Na+ is ~3 V, while the average voltage of LiCoO2 vs. Li/Li+ is ~3.9 V.4 A lower operating voltage results in a lower energy density. Therefore, the development of layered oxides with higher voltages while maintaining stable cycling performance would be a significant step forward for the development of practical NIBs. In general, layered sodium transition metal oxides can be grouped, using notation proposed by Delmas et al.,5 into two main categories: O3-type (Na+ in octahedral sites, ABCABC stacking) and P2-type (Na+ in prismatic sites, ABBA stacking).

P2-type sodium layered positive electrode materials like Na2/3Fe1/2Mn1/2O2 generally have higher specific capacities than their O3-type counterparts, such as NaFe1/2Mn1/2O2.6 P2-Na2/3Ni1/3Mn1/3O2 has one of the highest average discharge voltages reported (3.7 V) and a high reported capacity (150 mAh/g).7,8 Nevertheless, it has rapidly fading reversible capacity. The fade may be induced by structural transitions during charge/discharge, as indicated by multiple plateaus in its voltage curve.7,8 Specifically, the rapid capacity loss could be caused by the P2-O2 phase transition and resulting volume change when the material is charged to above 4.2 V.8 One strategy to address this issue is to replace Mn4+ with other 4+ cations, such as Ti4+.9 Komaba et al. showed that titanium substituted Na2/3Ni1/3Mn1/2Ti1/6O2 delivers 127 mAh/g reversible capacity with an average voltage of 3.7 V.9 The voltage curves are smooth and titanium-doped samples have superior cyclability compared to the Ti-free Na2/3Ni1/3Mn1/3O2. Replacing Ni2+ with other 2+ cations (Zn2+ or Mg2+) is another strategy.10,11 Wang et al. demonstrated that substitution of Ni2+ with Mg2+ can suppress the P2-O2 phase transition and thus Mg substituted Na2/3Ni1/3Mn2/3O2 has improved capacity retention.11 Nevertheless, both of these strategies result in decreased reversible capacity. One of the reasons is that Zn2+ or Mg2+ ions are not redox active and therefore decreasing the Ni2+ content can decrease capacity. The reason that Ti4+ doping also decreases the capacity is unknown. It is possible that Ti increases the average desodiation voltage and therefore the capacity obtainable at a fixed voltage range becomes lowered as a result.

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Recently it has been reported that Cu2+ is electrochemically active in layered sodium transition metal oxides.12–14 For instance, it was found that Cu2+ can be oxidized to Cu3+ during cycling in Na2/3Cu1/3Mn1/3O2 and Na2/3CuxMn1-xO2.12–14 Another copper-based material, P2Na7/9Cu2/9Fe1/9Mn2/3O2, also has high voltage (3.5 V) capacity and is air stable.15 Based on these results, and noting that Cu2+ and Ni2+ have similar ionic radii, copper could be a good substitution element in Na2/3Ni1/3Mn1/3O2. In this paper, we explore the structure, electrochemical performance, and air-stability of the series P2-Na2/3Ni1/3xCuxMn2/3O2 (x = 0, 1/12, 1/6, 1/4, 1/3). ■ EXPERIMENTAL SECTION Na2/3Ni1/3-xCuxMn2/3O2 compounds were prepared by solid-state reaction of Na2CO3 (99.0%, Sigma Aldrich), NiO (99 %, -325 mesh, Sigma Aldrich), CuO (99 %, Sigma Aldrich), and MnO2 (99 %, -325 mesh, Sigma Aldrich). A stoichiometric mixture of these materials was mixed by ball-milling in a Spex-8000 mill for 1 hour in air. An excess of 10% Na2CO3 was added to compensate sodium loss at high temperature. The powder mixtures were then pressed into pellets. For 1/12 ≤ x ≤ 1/3 compositions, pellets were heated at 900 °C in air for 15 hours. The Na2/3Ni1/3Mn2/3O2 sample was synthesized according to a previous report by Lu and Dahn.7 For this sample, pellets comprising the ball milled precursors were heated for 24 hours, also at 900 °C, followed by quenching in liquid nitrogen. All samples were transferred immediately to an argon filled glovebox after synthesis. Electrode preparation was conducted in glovebox. Slurries were made by mixing active materials, carbon black (from Super P, Erachem Europe) and PVDF binder (polyvinylidene fluoride, Kynar HSV 900) with a mass ratio of 80/10/10 in an appropriate amount of Nmethyl-2-pyrrolidone (Sigma Aldrich, anhydrous 99.5%). These components were mixed with two tungsten carbide balls in a Retsch PM200 rotary mill (100 rpm, 1 hour). The slurry was then coated onto aluminum foil using a coating bar with a 0.015 cm gap, and dried under vacuum at 80 °C overnight. Electrode discs 1.27 cm in diameter were punched from the as-prepared coatings. Coin cells (2325type) were assembled with sodium foils (0.04 cm in thickness, cold rolled from sodium ingot, Sigma Aldrich, ACS reagent grade) as counter / reference electrodes. Two Celgard 2300 and one polypropylene blown microfiber separator (3M Company) were used as separators. The electrolyte was 1 M NaPF6 (BASF) in a solution of ethylene carbonate, diethylcarbonate and monofluoroethylene carbonate (volume ratio 3:6:1, all from BASF). Cells were cycled at 30.0 ± 0.1°C with a Maccor Series 4000 Automated Test System at a constant current. Cycling rates were calculated by assuming a capacity of 170 mAh/g for all compositions in the Na2/3Ni1/3-xCuxMn2/3O2 series. For galvanostatic intermittent titration technique (GITT) measurements, cells were charged at a C/40 rate for 1 hour,

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followed by relaxing at open circuit for 5 hours between each charge. X-ray diffraction (XRD) patterns were measured with a Rigaku Ultima IV X-ray diffractometer equipped with a Cu anode X-ray tube and dual detectors. Powder samples were loaded into a gastight X-ray sample holder (DPM Solutions, Hebbville NS) in an argon-filled glove box to avoid air contamination. A D/TeX Ultra linear detector with a K-beta filter was used to measure XRD patterns of powder samples. For ex situ XRD studies, coin cells were prepared as described above, but were charged to 4.4V once and then stopped. Electrodes for ex situ XRD measurements were recovered from these cells in an argon-filled glove box, and the electrode materials were scraped off the aluminum foil and washed with dimethyl carbonate (DMC, BASF). The recovered electrode coating was then transferred onto a zero-background silicon wafer and sealed in the gastight sample holder for the ex situ XRD measurements. A scintillation detector with a diffracted beam monochromator was used to measure ex situ XRD patterns. A Phenom G2-pro scanning electron microscope (SEM, Nanoscience, Arizona) was used to determine particle size and morphology of the powder samples. Inductively coupled plasma optical emission spectrometry (ICP-OES) was used to determine the elemental composition of the products. Measurements were performed using a PerkinElmer Optima 8000. Each sample was dissolved in 2 mL aqua regia solution which was then diluted using 2% HNO3 prior to measurements. ■ RESULTS AND DISCUSSION Figure 1(a) shows the XRD patterns of Na2/3Ni1/3Samples with 0 ≤ x ≤ 1/4 were phase pure with the P2 structure and no noticeable impurities were detected. A phase pure sample could not be obtained for x = 1/3 (nominal Na2/3Cu1/3 Mn2/3O2) and an impurity phase, possibly CuO, was present in the final product. The formation of CuO impurity in Na2/3Cu1/3 Mn2/3O2 is consistent with recent reports by Kang et al. and Xu et al.13,14 The presence of CuO impurity indicates that the ratio of Cu/Mn is smaller than 1/2 in the P2-structure. We believe stoichiometric Na2/3Cu1/3 Mn2/3O2 is difficult to prepare, and that off-stoichiometric NayCu1/3-zMn2/3+zO2, where z is between 0 and 1/3, is formed instead. The presence of this off-stoichiometric phase cannot be confirmed by ICP analysis, for instance, since any Cu impurity formed as a byproduct would still be present in the sample. However, the formation of this off-stoichiometric phase can affect the trends in structure and electrochemistry, as will be discussed later. xCuxMn2/3O2.

Rietveld refinement was conducted for the XRD patterns with 0 ≤ x ≤ 1/4. The pattern calculated by Rietveld refinement of the P2 structure and a difference plot were also shown in Figure 1(a). An example of the refinement details is summarized in Table I for x = 1/6 (Na2/3Ni1/6Cu1/6Mn2/3O2). For a typical P2 structure, transition metals and sodium occupy alternating layers of


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(a) x=1/3(Na2/3Cu1/3Mn2/3O2) CuO

(b)

11.16

x=1/3(Na2/3Cu1/3Mn2/3O2)

(c)

3.294 Å 3.162 Å

11.14 11.12 x=1/4 x=1/4

3.288 Å

3.156 Å

2.92 2.9 2.88 x=1/6

x=1/6

3.288 Å 3.156 Å

2.86 0

0.1 0.2 0.3 x in Na2/3Ni1/3-xCuxMn2/3O2

0.5

Naf Nae

x=1/12

x=1/12 3.282 Å

3.156 Å

0.4

0.3 x=0 (Na2/3Ni1/3Mn2/3O2) x=0 (Na2/3Ni1/3Mn2/3O2)

(d)

0.2 0

10

20

30 40 50 60 Scattering Angle (deg.)

70

3.282 Å 3.156 Å

0.1 0.2 0.3 x in Na2/3Ni1/3-xCuxMn2/3O2

80

26 27 28 29 Scattering Angle (deg.)

Figure 1.(a) X-ray diffraction patterns and Rietveld refinements of compounds in the Na2/3Ni1/3-xCuxMn2/3O2 series. (b) Changes in lattice parameters of the Na2/3Ni1/3-xCuxMn2/3O2 series. (c) Enlarged XRD patterns of the Na2/3Ni1/3-xCuxMn2/3O2 series between 25.5° to 29.5° including d-spacings. (d) Changes in Naf and Nae site occupancies of the Na2/3Ni1/3-xCuxMn2/3O2 series.

Table I. Parameters used for Rietveld refinement of the Na2/3Cu1/6Ni1/6Mn2/3O2 using space group P63/mmc (No. 194) 2

Atom

Site

x

y

z

Occupancy

B (Å )

Naf

2b

0

0

0.25

0.245±0.007

4.2±0.4

Nae

2d

1/3

2/3

0.25

0.422±0.007

4.2±0.4

Ni

2a

0

0

0

1/6

0.7±0.1

Mn

2a

0

0

0

2/3

0.7±0.1

Cu

2a

0

0

0

1/6

0.7±0.1

O

4f

2/3

1/3

0.0892±0.0006

1.02±0.02

0.01±0.3

a = 2.899±0.0001 Å, c = 11.137±0.0008 Å, Rb = 3.67%, Rwp = 11.468%

Figure 2. SEM images of compounds in the Na2/3Ni1/3- xCuxMn2/3O2 series for x = (a) 0, (b) 1/12, (c) 1/6, (d) 1/4 and (e) 1/3



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Table II. Target compositions compared with ICP measurement results (normalized to a total transition metal population of 1.000) Composition x=0

x=1/12

x=1/6

x=1/4

x=1/3

Na

Ni

Mn

Cu

Target

0.667

0.333

0.667

0.000

ICP

0.731

0.367

0.630

0.003

Target

0.667

0.250

0.667

0.083

ICP

0.753

0.277

0.639

0.084

Target

0.667

0.167

0.667

0.167

ICP

0.775

0.190

0.636

0.174

Target

0.667

0.083

0.667

0.250

ICP

0.751

0.096

0.641

0.263

Target

0.667

0.000

0.667

0.333

ICP

0.768

0.001

0.644

0.355

octahedral and prismatic sites, respectively, with vacancies occurring in the sodium layer. There are two different prismatic sites in the sodium layer that share either faces or edges with the transition metal octahedrons MO6. These are denoted here as Naf and Nae sites, respectively. For the refinements, face-sharing and edge-sharing sodium atoms were placed in 2b and 2d sites, respectively. Transition metals were placed in 2a sites and oxygen atoms were placed in 4f sites. Na and Ni mixing was not considered in refinements due to the relatively large difference in the ionic radii of Na+ (1.02 Å) and Ni2+ (0.69 Å). The high displacement parameters refined for Na ions (~4 Å2) agree with previous reports on P2-type materials and is associated with the fast diffusion of Na ions in the Na layer. 16 Komaba et al. studied the high displacement parameter using the maximum entropy method (MEM) and found that MEM analysis also suggest that the Na ions are widely distributed.16 Figure 1(b) shows how the lattice constants a and c change with respect to x. Best fit lines according to Vegard's law are also shown for 0 ≤ x ≤ 1/4. It should be noted that the evolution of cell parameter c has a curved shape, indicating Vegard’s law is not precisely followed. It is speculated that transition metal ordering/sodium ordering also affects the lattice parameters, leading to this nonlinearity. Generally, as x is increased the lattice constants a and c both increase. However the amplitude of variation is smaller than in many other systems reported.17 The minor change in a and c with x can be attributed to the slightly smaller ionic radius of Ni2+ (0.69 Å) compared to Cu2+ (0.73 Å). One exception in the series is Na2/3Cu1/3Mn2/3O2 (shown in solid circles), which shows a decrease in lattice constants a and c, compared to Na2/3Ni1/12Cu1/4Mn2/3O2 (x=1/4). As discussed above, Na2/3Cu1/3Mn2/3O2 is actually off-stoichiometric NayCu1/3zMn2/3+zO2, which has more manganese in the P2structure than the Na2/3Ni1/3-xCuxMn2/3O2 series with 0 ≤ x

≤ 1/4, resulting in the formation of Mn3+. Since the ionic radius of Mn3+ (0.645 Å) is smaller than that of both Ni2+ (0.69 Å) and Cu2+ (0.73 Å), the lattice parameters a and c decrease from Na2/3Ni1/12Cu1/4Mn2/3O2 (x=1/4) to Na2/3+zCu1/3-zMn2/3+zO2. Figure 1(c) shows an expanded view of the Na2/3Ni1/3-xCuxMn2/3O2 XRD patterns between 25.5° and 29.5°. Two unrefined superstructure peaks at ~27.2° and ~28.3° can be observed at this region, which likely arise from the in-plane sodium/vacancy ordering. The dspacings for these peaks are also shown for comparison. The d-spacings are close to those reported by Meng et al. for Na2/3Ni1/3Mn2/3O2.8 For the Na2/3Ni1/3-xCuxMn2/3O2 series, as x increases, the low angle peak gradually shifts to the left, corresponding to an increase in d-spacings, while the higher angle peak almost does not shift. The increase in d-spacings agrees with the expansion of lattice in the a-b plane as discussed above. Meng et al. studied the sodium/vacancy ordering and ascribed these peaks to the “large zigzag” ordering of Naf ions.8 Typically for a P2 type material, there are more sodium atoms occupying Nae sites than Naf sites. This is because Naf sites have higher energy than Nae sites due to the stronger repulsion between sodium ions and adjacent transition metal ions. The relative difference leads to a higher occupancy ratio of sodium ions in Nae sites than Naf sites in order to minimize the repulsion. It can be seen from Figure 1(c) that as the copper content increases, the intensity of these superstructure peaks also increases. This could be attributed to the increasing lattice constant c, which reduces the repulsion between sodium in Naf sites and neighboring transition metals. Therefore as the copper content is increased, more sodium can be placed in Naf sites and the “large zigzag” ordering between Naf ions becomes more distinct. This model was confirmed by Rietveld refinements. The sodium occupancy ratio in each site as obtained by the refinements are shown


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Figure 3. Voltage curves (left) and corresponding differential capacity curves (right) of Na2/3Ni1/3-xCuxMn2/3O2 vs. Na cells cycled between 2.5 V – 4.1 V (black lines) and 2.5 V – 4.4 V (blue lines).

in Figure 1(d). As the copper content increases, the occupancy ratio of sodium increases in Naf sites and decreases in Nae sites. The SEM images are shown in Figure 2. Apparently all of the samples have a plate-like hexagonal morphology. The particle sizes range from 1 - 2 micrometers and increases as copper content increases. Na2/3Ni1/3Mn2/3O2 has smaller particle size than the other compounds in the series, which might be due to the quenching process during its synthesis. ICP results are compared to target compositions in Table II. All calibration curves were linear, with an R2 value of 0.9999851.00000. The standard deviations are less than 3% for all the measurements. The ICP results show good agreement with the anticipated atomic ratios for all the transition metals. However, extra sodium content was observed. These values are consistent with the 10% excess Na used

Capacity (mAh/g)

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Figure 4. Discharge capacity versus cycle number of Na2/3Ni1/3-xCuxMn2/3O2 vs. Na cells cycled between (top) 2.5 V – 4.1 V and (bottom) 2.5 V – 4.4 V cycled at 1C current rate.



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120 2.5 V - 4.1 V C/20

C/5 C/2 1C

2C 5C

C/2

x=1/12 x=1/6 x=1/4 x=1/3 x=0 (Na2/3Ni1/3Mn2/3O2)

40

0

Voltage (V)

80

0

10

(a)

20 30 Cycle number

160 C/20

C/2 5C

80

x=1/12 x=1/6 x=1/4 x=1/3 x=0 (Na2/3Ni1/3Mn2/3O2)

40 0

50

2.5 V - 4.4 V C/5 C/2 1C 2C

120

40

Voltage (V)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

0

10

20 30 Cycle number

(b) 40

50

Figure 5. (a) Discharge capacity of Na2/3Ni1/3-xCuxMn2/3O2 compounds cycled at different C-rates between 2.5 V - 4.1 V and (b) between 2.5 V – 4.4 V. (c) Voltage curves of Na2/3Cu1/12Ni1/4Mn2/3O2 cycled at different C-rates between 2.5 V - 4.1 V, and (d) between 2.5 V - 4.4 V.

during synthesis. This additional Na likely exists as Na2O or Na2CO3. The initial charge/discharge (corresponding to sodium deintercalation/intercalation) curves of Na2/3Ni1/3xCuxMn2/3O2 vs Na cells cycled between 2.5 V and 4.1 V or 2.5 and 4.4 V are compared in Figure 3. Figure 3 also shows the corresponding differential capacity curves for the cells. A hysteresis of ~0.2 V between charge and discharge half-cycles was observed for all the materials. The redox-limited capacities of Na2/3Ni1/3-xCuxMn2/3O2 compounds are also shown in the figure, as calculated based on the Ni2+/Ni4+ and/or Cu2+/Cu3+ redox couples. Because of the similar atomic weight of Ni (58.69) and Cu (63.55), the formula weight varies only slightly in the composition series. Therefore the change in redox-limited capacity originates almost exclusively from the substitution of Ni2+/Ni4+ with Cu2+/Cu3+ redox couples. With increasing copper content, the amount of Ni2+ decreases and the redox-limited capacity decreases from 85 mAh/g (Na2/3Cu1/3Mn2/3O2) to 172 mAh/g (Na2/3Ni1/3Mn2/3O2). On the other hand, the sodium-limited capacity, referring to the capacity if all the sodium can be extracted, almost does not change (169 mAh/g for Na2/3Cu1/3Mn2/3O2). The voltage curve for the sample with x=0 (Na2/3Ni1/3 Mn2/3O2) agrees well with previous reports.7,8 The step-like voltage curve indicates several sodi-

um/vacancy ordering and phase transitions. The material exhibits a first charge/discharge capacity of ~80/75 mAh/g when cycled between 2.5 V and 4.1 V. Accessing voltages above 4.1 V results in a flat voltage plateau at ~4.2 V. The full capacity of this plateau is accessed by charging the material to 4.4 V, which results in a plateau capacity of about 80 mAh/g. At this point almost all the sodium is removed from the structure. A first charge/discharge capacity of ~160/140 mAh/g is achieved, with ~20 mAh/g irreversible capacity. The first cycle Coulombic efficiency is ~ 88%. Lu et al. showed that the irreversible capacity can be reduced by discharging the material to a lower cutoff voltage (2V).7 According to Meng et al.,8 the long plateau at ~4.2 V originates from the P2/O2 phase transition, which possibly causes the capacity fade. The discharge capacity rapidly fades during cycling above 4.1 V (shown in Figure 4) and after 15 cycles less than half of the initial capacity is retained. The multiple plateaus present in the Na2/3Ni1/3 Mn2/3O2 voltage curve correspond to the multiple sharp peaks in its differential capacity (Figure 3). In contrast, copper-substituted samples have smoother voltage curves and broad peaks with low intensity were observed in the differential capacity. Even when x = 1/12 in Na2/3Ni1/3xCuxMn2/3O2, the voltage curve was effectively smoothed and capacity retention improved. Na2/3Cu1/12Ni1/4Mn2/3O2


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tution

Figure

6:

Sodium

xCuxMn2/3O2 calculated

112

012 013

011

110

002

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diffusion coefficient in Na2/3Ni1/3from GITT as a function of voltage.

has a reversible capacity of ~130 mAh/g when cycled between 2.5 V and 4.4 V. Higher copper contents can smooth the voltage curve further, at the expense, however, of lower reversible capacities. On the other hand, when x was increased to 1/3 (Na2/3Cu1/3Mn2/3O2), the steplike voltage curve reappears and sharp peaks were found in the differential capacity curve. The voltage curve agrees with previous reports of Na2/3Cu1/3Mn2/3O2. However, the observed capacity for Na2/3Cu1/3Mn2/3O2 is larger than its theoretical capacity based solely on transition metal oxidation when it is cycled between 2.5 V – 4.4 V. As discussed earlier, this could be due to the formation of off stoichiometric NayCu1/3-zMn2/3+zO2. The presence of Mn3+ in this structure would contribute to the total capacity. On the other hand, when Na2/3Cu1/3Mn2/3O2 is cycled between 2.5 V – 4.4 V, an irreversible plateau appears at ~4.3 V, corresponding to an irreversible sharp peak at ~4.3 V in differential capacity curve. Combined with the anomalously large voltage hysteresis and the extra capacity, this suggests an irreversible structural change, possibly due to oxygen redox reaction. Further studies are needed to understand the source of extra capacity in Na2/3Cu1/3Mn2/3O2 when charged to 4.4 V. The reversible capacities of all samples exceed their theoretical capacity based solely on either the Ni2+/Ni4+ (red dotted line) or Cu2+/Cu3+ (green dotted line) redox couples. Therefore both Ni2+ and Cu2+ participate in the redox reaction during cell cycling. Interestingly, when these materials were cycled between 2.5 V and 4.1 V, the reversible capacity actually increases from ~75 mAh/g for Na2/3Ni1/3Mn2/3O2 (x=0) to ~90 mAh/g for Na2/3Cu1/12Ni1/4Mn2/3O2 and Na2/3Cu1/6Ni1/6Mn2/3O2 (x=1/12 and 1/6), despite the decreased nickel content. This is caused by the lower redox potential of Cu2+/Cu3+ than Ni2+/Ni4+, resulting in a higher capacity utilization ratio in the low voltage range. As can be seen in Figure 3, while only half of the capacity is obtained for Na2/3Ni1/3Mn2/3O2 when cycled between 2.5 V and 4.1 V, ~80% of the capacity is obtained for Na2/3Cu1/3Mn2/3O2 at this voltage range. Compared with Mg or Zn, 10,11 which are redox inactive, Cu substitution results in a higher capacity at similar substi-

No sign of O2 phase Figure 7. Ex situ X-ray diffraction patterns of compounds in the Na2/3Ni1/3-xCuxMn2/3O2 series after being fully charged to 4.4 V. level. The high capacity originates from the inclusion of the active Cu2+/Cu3+ redox couple, rather than the inactive Zn2+ and Mg2+. The

cycling performance of Na2/3Ni1/3vs. Na cells are shown in Figure 4. When cycled between 2.5 V and 4.1 V at 1C, all the materials show good capacity retention after 30 cycles. When cycled between 2.5 V and 4.4 V at 1C, severe capacity fade was observed for Na2/3Ni1/3Mn2/3O2 and less than half of its first cycle capacity was retained after 15 cycles. Copper substitution effectively improves the capacity retention of electrodes cycled at this voltage range. Na2/3Cu1/12Ni1/4Mn2/3O2 retained 75% of its initial capacity after 30 cycles. Higher copper content is beneficial for better capacity retention. For example, Na2/3Cu1/6Ni1/6Mn2/3O2 has a capacity retention of 97% after 30 cycles. xCuxMn2/3O2

Figure 5a and Figure 5b show the discharge capacity for Na2/3Ni1/3-xCuxMn2/3O2 vs. Na cells cycled at different C-rates within 2.5 V - 4.1 V and 2.5 V – 4.4 V, respectively. Figure 5c and Figure 5d show the voltage curves of Na2/3Cu1/12Ni1/4Mn2/3O2 cycled at different C-rates between 2.5 V - 4.1 V and 2.5 V – 4.4 V, respectively. All of the materials with Cu substitution exhibit improved rate capability, especially in the voltage range of 2.5 V – 4.4 V. When cycled between 2.5 V and 4.1 V,



Figure 7 shows ex situ XRD patterns measured for each compound in the Na2/3Ni1/3-xCuxMn2/3O2 series after charging to 4.4 V. After charging to 4.4 V, Cu-free Na2/3Ni1/3 Mn2/3O2 (x=0) has an O2-type structure with random stacking faults. This is consistent with a P2-O2 phase transition during desodiation, as reported by Lu et al.18 This P2-O2 phase transformation has been associated with the rapid capacity fade.8 In contrast, for copper substituted samples, ex situ XRD patterns taken after charging to 4.4 V (Figure 7) show no sign of a P2-O2 phase transition. At ~20°, where the (002) peak of O2 phase usually occurs, no peak was observed for the copper substituted samples. The major reflections corresponding to the P2 phase are maintained after being charged to 4.4 V. The result indicates that the addition of Cu, similar to Mg,11 can also suppress the P2-O2 phase transition. A closer examination shows that excepting x=1/4, a splitting of (002) peak for all other compositions can be observed. Such peak splitting has been assigned to the O2-OP4 phase transition.20 The absence of such splitting for the material with x=1/4 is consistent with the results in Figure 3. In Figure 3, all the while other materials have a plateau at ~4.2 V, except for the one with x=1/4. This result is similar to a previous report by Billaud et al. who studied the Na2/3MgxMn1-xO2 series and observed in the ex-situ XRD patterns that only the materials with a plateau at ~4.2 V have such OP4 peak splitting.20 Therefore, instead of the O2-P2 transition, the material with x=1/4 retains its P2 phase, while a P2-OP4 transition occurs for other Cudoped materials. The suppression of O2-P2 phase transi-

Intensity (a.u.)

tions in the Cu substituted materials might be responsible for their improved cycling performance.

Intensity (a.u.)

Na2/3Cu1/12Ni1/4Mn2/3O2 has a discharge capacity of 90, 88, 87, 86 and 84 mAh/g at current rates of C/20, C/5, C/2, 1C and 2C with an average voltage of ~3.4 V. Even at a high rate of 5C (850mA/g), a discharge capacity of 80 mAh/g was retained with an average voltage of ~3.3 V, corresponding to 89% of its capacity at C/20. When cycled between 2.5 V and 4.4 V, Na2/3Cu1/12Ni1/4Mn2/3O2 has a discharge capacity of 130, 122, 116, 110 and 102 mAh/g at current rates of C/20, C/5, C/2, 1C and 2C, all exceeding 75% of its capacity at C/20. Figure 6 shows the sodium diffusion coefficients (DNa) in Na2/3Ni1/3-xCuxMn2/3O2 as calculated from GITT measurements. All samples have a similar DNa at voltages below 4.0 V. However, a rapid DNa drop was observed for Na2/3Ni1/3Mn2/3O2 at voltages above 4.0 V. This phenomenon agrees with the results obtained by Meng et al., who observed that the minimum value of DNa occurred at low sodium content.8 The P2-O2 phase transition is responsible for the low DNa, as Meng et al. show that sodium has a much lower diffusion barrier in P2 structure than in O2 structure.8 For comparison, the DNa for Na2/3Ni1/3xCuxMn2/3O2 with x=1/12 and x=1/6 is much higher at voltages larger than 4.0 V, which indicates a faster sodium diffusivity, leading to a better rate capability at high voltage ranges. Materials with high copper content also show smoother DNa vs. voltage curve, suggesting a more stable structure at different sodium content. The result also implies that copper substitution could suppress the P2-O2 phase transition at high voltage.

Intensity (a.u.)

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Figure 8. X-ray diffraction patterns of Na2/3CuxNi1/3after air exposure, compared to pristine materials.

xMn2/3O2

Unlike many other Na-M-O materials that readily react with water, Na2/3Ni1/3Mn2/3O2 has been reported to be air-stable and does not uptake water into its structure.19 Lu et al. speculated that the superlattice ordering of transition metals induces a strong coupling between adjacent MO2 sheets, which prevents the intercalation of water.19 However, it has also been reported that as Co3+ is added to form P2-Na2/3[CoxNi1/3-xMn2/3]O2, the superlattice ordering is suppressed and the materials become hygroscopic. P2-Na2/3[CoxNi1/3-xMn2/3]O2 readily reacts with water in the air and forms a hydrated phase, which can be easily spotted as the diffraction pattern changes significantly from pristine phase even after only a few hours of air exposure.19 The stability of materials against moisture is important as hygroscopic materials require a dry handling system, which would increase the cost of manufacturing.21 For this reason we studied the air-sensitivity of copper doped P2-Na2/3[CuxNi1/3-xMn2/3]O2 Figure 8 shows the XRD patterns of P2Na2/3[CuxNi1/3-xMn2/3]O2 powders after being exposed to air


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Voltage (V)

for 5 days and 21 days. Compared to XRD patterns of un-

Voltage (V)

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Chemistry of Materials exposed samples, no noticeable change was observed. ferred back into the glovebox without any further treatment and incorporated into coin cells as described above. Figure 9 shows the first two cycles of the voltage curves for electrodes exposed to air for 5 days, compared to the voltage curve of electrodes without air exposure. It was found that all the cells with air exposed electrodes have a higher open circuit voltage (~0.4 V) compared to cells made with air-protected electrodes. This suggests that the sodium content in the air exposed samples is slightly less than the as-synthesized samples. However, except for the Na2/3Cu1/3Mn2/3O2 (x=1/3), the average voltage and reversible capacity did not change and there is almost no decay in electrochemical performance, indicating the airstability of these materials. For Na2/3Cu1/3Mn2/3O2 (x=1/3), a smaller reversible capacity was obtained after air exposure. The electrochemical test results agree with the XRD measurements that the Na2/3CuxNi1/3-xMn2/3O2 series (0 ≤ x ≤ 1/4) are air-stable, while Na2/3Cu1/3Mn2/3O2 is airsensitive. The air-stability of these materials demonstrates the feasibility of using Na2/3CuxNi1/3-xMn2/3O2 for practical sodium ion batteries. ■ CONCLUSIONS

Figure 9. Voltage curves of Na2/3CuxNi1/3-xMn2/3O2 after 5 days of air exposure (right), compared to pristine materials (left). The positions of the Bragg peaks do not change and no new peaks are formed. Even after 3 weeks of air exposure, XRD patterns of Na2/3[CuxNi1/3-xMn2/3]O2 materials show no significant change, except for Na2/3Cu1/3Mn2/3O2. It can be observed that after air exposure of Na2/3Cu1/3Mn2/3O2, there appears a weak peak formed at ~14° and a broad peak at ~28°, which are indicated by arrows in Figure 7. This is indicative of water intercalation as reported by Lu et al., who observed a similar XRD pattern for air-exposed Na2/3Co1/6Ni1/6Mn2/3O2.19 Lu and Dahn proposed that these two peaks are due to a shift of the (002) and (004) peaks of the pristine phase upon the intercalation of water. We speculate that copper does not disturb transition metal ordering due to the similar ionic radii of Cu2+ (0.73)and Ni2+(0.69), while the relatively large ionic radii difference of to Co3+(0.545) and Ni2+(0.69) would suppress transition metal superlattice ordering. Lu and Dahn also suggested that the presence of Mn3+, which forms as Co3+ is added, also possibly aids the intercalation of water. However, when Cu2+ is added to replace Ni2+, the oxidation state of Mn does not change and presumably no Mn3+ is formed. This may prevent water intercalation and be the reason why Cu imparts air stability to these materials. It can be concluded that water cannot be intercalated into Na2/3CuxNi1/3-xMn2/3O2 (0 ≤ x ≤ 1/4) under the air exposure conditions employed here, but can be intercalated into Na2/3Cu1/3Mn2/3O2. The air stability of the materials was further studied in electrochemical tests. Electrode discs were intentionally taken from the glovebox and exposed to air on a lab bench. After 5 days, these electrodes were trans-

P2-type Na2/3Ni1/3Mn2/3O2 could be a promising cathode candidate in sodium ion batteries for large scale applications such as grid storage due to its high capacity, high operating voltage and air stability. However, this material suffers from its step-like voltage curve and rapid capacity fade during cycling. A doping strategy that substitutes part of the nickel with copper was proven to be an effective way to smooth the voltage curve and improve capacity retention and rate capability. Both Ni2+/Ni4+ and Cu2+/Cu3+ participate in the redox reaction during cycling, therefore less capacity reduction is observed compared to when Ni is substituted with Mg or Zn. Ex situ XRD studies reveal that these materials do not convert to O2 phase when charged to a high voltage, which might be the reason for their improved cycling performance. The doping strategy is a promising approach for the future development of positive electrode materials for sodium ion batteries.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT The authors would like to recognize financial support from 3M Canada and NSERC under the auspices of the Industrial Research Grants and Discovery Grants programs.

REFERENCES 1. Dunn, B.; Kamath, H.; Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices. Science. 2011, 334, 928–935.



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Crystal Structures and Electrochemical Performance of Air-Stable Na2

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