3)yO2 Cathode Materials for Sodium

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On the P2-NaxCo1-y(Mn2/3Ni1/3)yO2 cathode materials for sodiumion batteries: Synthesis, electrochemical performance and redox processes occurring during the electrochemical cycling Siham Doubaji, Lu Ma, Habtom Desta Asfaw, Ilyasse Izanzar, Rui Xu, Jones Alami, Jun Lu, Tianpin Wu, Khalil Amine, Kristina Edstrom, and Ismael Saadoune ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b13472 • Publication Date (Web): 03 Nov 2017 Downloaded from http://pubs.acs.org on November 4, 2017

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ACS Applied Materials & Interfaces

On the P2-NaxCo1-y(Mn2/3Ni1/3)yO2 cathode materials for sodium-ion batteries: Synthesis, electrochemical performance and redox processes occurring during the electrochemical cycling. Siham DOUBAJIa,b, Lu MAb, , Habtom Desta ASFAWc, Ilyasse IZANZARa, Rui XUd, Jones ALAMIe, Jun LUd , Tianpin WUb, Khalil AMINEd*, Kristina EDSTRÖMc, Ismael SAADOUNEa,e a

LCME, FST Marrakesh, University Cadi Ayyad, Av AbdelkrimKhattabi, Box 511, 40000, Marrakech, Morocco

b

X-ray Science Division, Advanced Photon Sources, Argonne National Laboratory, 9700 South Cass Avenue, Lemont, IL, USA

c

d

Department of Chemistry-Ångström Laboratory, Uppsala University, Box 538, 721 21 Uppsala, Sweden

Chemical Science and Engineering Division, Argonne National Laboratory, 9700 South Cass Avenue, Lemont, IL, USA e

Materials Science and Nano-Engineering Department, Mohammed VI Polytechnic University, Lot 660-Hay My Rachid, Ben Guerir, Morocco

Abstract P2-type NaMO2 sodiated layered oxides with mixed transition metals are receiving considerable attention for use as cathodes in sodium-ion batteries. A study on the solid solution (1-y) P2-NaxCoO2-(y) P2-NaxMn2/3Ni1/3O2 (y = 0, 1/3, 1/2, 2/3, 1) reveals that changing the composition of the transition metals affects the resulting structure and the stability of pure P2 phases at various temperatures of calcination. For 0 ≤ y ≤ 1.0, the P2NaxCo(1-y)Mn2y/3Niy/3O2 solid solution compounds deliver good electrochemical performance when cycled between 2.0 and 4.2 V vs. Na+/Na with improved capacity stability in long-term cycling, especially for electrode materials with lower Co content (y = 1/2 and y = 2/3), despite lower discharge capacities being observed. The composition ½ P2-NaxCoO2-½ P2NaxMn2/3Ni1/3O2 delivers a discharge capacity of 101.04 mAh g-1 with a capacity loss of only 3% after 100 cycles and a coulombic efficiency exceeding 99.2%. Cycling this material to a higher cut-off voltage of 4.5 V vs. Na+/Na increases the specific discharge capacity to ≈ 140 mAh g-1 due to the appearance of a well-defined high voltage plateau, but after only 20 1 ACS Paragon Plus Environment


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cycles capacity retention declines to 88% and coulombic efficiency drops to around 97%. Insitu X-ray absorption near edge structure (XANES) measurements conducted on the composition NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2) in the two potential windows studied help elucidate the operating potential of each transition metal redox couple. It also reveals that at the high voltage plateau, all the transition metals are stable, raising the suspicion of possible contribution of oxygen ions in the high voltage plateau. Keywords: Na-ion batteries, P2-type

materials, energy storage, in-situ XANES

measurements, high voltage plateau. *Corresponding author; Email: [email protected]

2 ACS Paragon Plus Environment

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1. Introduction Lithium-ion batteries are the most common rechargeable batteries used in portable electronic devices, and they are also becoming an efficient energy choice used to power hybrid, plug-in, and electric vehicles. However, concerns have been raised about the availability and cost of Lithium in recent years.1 Due to the high abundance and the low cost of sodium, as well as its suitable redox potential (E(Na+/Na) = -2.71 V vs. SHE), sodium-ion batteries are regarded as possible replacements for lithium-ion ones, especially for large-scale energy-storage applications. Thus, the electrochemical energy storage community has been devoting increased attention to designing new anode and cathode materials for sodium-ion batteries, making it a robust research area within rechargeable battery systems. Sodium layered oxides with the general formula NaxMO2 (M being one or several transition metals) have been extensively studied as cathode materials for sodium-ion batteries, and their ability to electrochemically intercalate sodium ions has been known since the 1980s.2-6 Unlike LiMO2, NaxMO2 possesses different crystal structures, where the Na is located either in octahedral or in prismatic sites between the layers formed by edge-sharing MO6 octahedra. These structures, presented in Scheme 1, differ in the stacking of the oxygen layers with ABCABC for O3, ABBA for P2, and ABBCCA for P3 (according to the Delmas notation, where O and P refer, respectively, to the octahedral and prismatic sites of sodium).2 The most common structures studied are the O3 type (R-3m space group), where xNa ≈ 1 and the P2 type (P63/mmc space group), where xNa ≈ 2/3. The P2-type layered structure has an open path for Na ions, where they migrate from one prismatic site to adjacent sites through open square bottlenecks surrounded by four oxide ions.7 In addition, this structure exhibits stable electrochemical and structural behavior during cycling with only one structural transition from P2 to O2,8 or to OP4,9 or to an unknown “Z”,10 type at high voltages when the sodium



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content is less than xNa≤0.35. These characteristics offer to P2-type layered oxides better cycle life and improved energy efficiency.

--A

MO6

--B

Na

--A --B

--A --B --B --C

Nae

--A

--B --A

Naf

--B

--A

--C

--B

--C --A

--A

--B

--A

--C

--B

O3

--B

P2

P3

Scheme 1. Illustration of O3, P2, and P3 structures The material P2-NaxCoO2 was studied first in the 1980s as a cathode material for sodium-ion batteries, demonstrating that sodium ions can be reversibly extracted and inserted into the material.3,11 It was reinvestigated very recently in cell tests,12, performance

results

with

good

reversible

capacity

13

and has shown promising and

cyclability.

The

intercalation/deintercalation of sodium ions in this material is evident in well-defined steps in the voltage profile. These voltage steps are caused either by gliding of oxide-ion layers, sodium-ion ordering and oxygen defficiency.14,15 The substitution of Co with other transition metals was found to help suppress sodium ordering, where substituting Co with Mn in P2NaCo2/3Mn1/3O2,16 stabilized the structure, and the material exhibited only one voltage step at xNa ≈ 0.5 when cycled between 1.25 and 4.0 V. However, two studies on solid-solution P2NaxCo1-yMnyO2, reported later,17,18 showed that electrochemical cycling was not improved over the long term, because adding Mn actually lowered the material’s capacity retention. The substitution of Co with Ca has also been reported,19 and the presence of Ca2+ in the prismatic sites between CoO6 slabs suppresses Na ordering and stabilizes the electrochemical cycling. In fact, the substitution with inactive metals ions has showed to improve the electrochemical


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performance of P2-type materials.20,21 On the other hand, P2 sodium layered oxides combining Co, Mn, and Ni have recently shown stable cyclability and good specific capacities.22-25 We have also reported, in previous work, P2-NaxCo2/3Mn2/9Ni1/9O2 as a possible candidate to use in sodium-ion batteries.26 The dependence of the crystallographic structure of these materials on the temperature was reported earlier in different studies.25,27 They tend to adopt a P3-type structure or a mixture of P3 and P2-type structures at temperatures below 800°C and P2-type structure at higher temperatures. Nevertheless, no noticeable interest has been given to the impact of the composition of the transition metals on the stable final structure. In this paper, we report the synthesis of solid-solution (1 - y) P2-NaxCoO2–(y) P2NaxMn2/3Ni1/3O2 (y = 0, 1/3, 1/2, 2/3, 1) by a simple sol gel method. The substitution of Co with Mn and Ni using these specific compositions relies on the assumption that Mn is in the +4 state, thereby avoiding the Jahn-Teller distortion associated with the presence of Mn at the trivalent state and thus stabilizing the structure. Furthermore, Ni was selected as a substituting ion to enhance the energy density since it exhibits a high redox potential, which elevates the average discharge potential. XRD results of the materials calcinated at different temperature values ranging from 700°C to 950°C are presented. The impact of changing the composition on the oxidation states of transition metals is studied using XANES measurements. All the materials are cycled between 2.0 and 4.2V vs Na+/Na at a rate of C/20 and ½ P2- NaxCoO2-½ P2-NaxMn2/3Ni1/3O2 (y = ½) along with the end members of the solid solution (y = 0 and 1) are cycled up to 4.5V vs Na+/Na. Furthermore, in-situ XANES measurements are conducted for NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2) in the two potential windows studied. This study is a step forward toward understanding the choice of cathode materials depending on the application desired, and it also raises questions about the high-voltage plateau observed in P2-type structure layered oxides. 5 ACS Paragon Plus Environment


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2. Experimental section

Synthesis of materials. NaxCo(1-y)Mn2y/3Niy/3O2 (0 ≤ y ≤ 1.0) layered oxides were synthesized via a sol-gel method using sodium, cobalt, manganese, and nickel acetates with a starting sodium molar content of 0.7 with 5 wt% excess. For each material, the stoichiometric amounts of the precursors were mixed in distilled water and stirred for 2 hours before addition of citric acid (chelating agent). The mixture was kept at 80 °C under constant stirring until a homogeneous gel was obtained. The obtained gel was then dried over night at 110 °C to obtain a powder, which was ball-milled and further heat-treated at temperatures ranging from 700°C to 950°C for 12 hours at a heat rate of 2°C/min, followed by slow cooling. The obtained materials were finally stored in an Ar-filled glove box (H2O and O2 < 1 ppm).

Structural characterization. The crystalline structure of the synthesized materials was characterized by X-ray diffraction (XRD) using a Bruker D8 Advance diffractometer equipped with Cu Kα radiation. The XRD patterns were collected in the 2θ range of 10-90° in a continuous scan mode with a step size of 0.01° and a constant counting time of 10 s/step. Lattice parameters were refined with a typical Rietveld method implemented in the FullProf program. The morphology and the size distribution of the particles were studied using the Merlin. scanning electron microscope of Zeiss make operating at an accelerating voltage of 5 kV and a working distance of 3.5 to 3.8 mm. A charge compensator was employed in all cases in order to alleviate charging effects due to the insulating nature of the oxides.

Electrochemical tests. The composite electrodes were prepared by mixing 80 wt% of the active material, 15 wt% of carbon black (Super P) conductive additive, and 5 wt% of polyvinylidene fluoride (PVDF) binder using N-methyl-2-pyrrolidone (NMP) as solvent. The slurry was then casted on an Al foil and dried at 60 °C for 3 hours in a convection oven. The electrodes were cut into 12 mm disks by a precision perforator (Hohsen) and dried overnight at 120°C in a vacuum oven within an Ar-filled glove box (M-Braun). Sodium metal was used 6 ACS Paragon Plus Environment

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as both reference and counter electrode by cutting, rolling, and pressing sodium lumps into thin plates. A glassfiber (Whatman) separator was used as separator between the sodium plate and the working electrode. The electrolyte was 0.5 M NaPF6 dissolved in polycarbonate (PC). The components were assembled in “coffee-bag” (polymer laminated aluminum pouch) cells or coin cells and vacuum-sealed in the Ar-filled glove box. All the electrochemical measurements were carried out at room temperature (25°C) via VMP2 (Bio-Logic) equipment. The charge/discharge studies were performed galvanostatically at a current rate of C/20, where a period of 20 hours is required to remove one sodium ion. In particular, two cutoff voltages were studied: the first one between 2.0 and 4.2 V and the second between 2.0 and 4.5 V. The reproducibility was checked for all the experiments, including the synthesis, by repeating all the steps at least twice.

X-ray absorption spectroscopy (XAS). XAS measurements at Co, Mn, and Ni K-edges were performed at the beamline 9-BM of the Advanced Photon Source (APS) with electron energy of 7 GeV and an average current of 100 mA. The radiation was monochromatized by a Si (111) double-crystal monochromator. XANES spectra of pristine electrodes were recorded in the fluorescence mode by PIPS detector, and in-situ XANES measurements were recorded in the transmission mode by using an operando coin cell. The cases and spacer of the coin cell used were punched with a hole (dia. =1/8 inch) in the center, and both negative and positive cases were covered with Kapton Tape as a window for the X-ray penetration (Figure S1). Data reduction and analysis were carried out with ATHENA software. 3. Results and discussion 3.1.

Synthesis and structural characterization

P2-NaxCo(1-y)Mn2y/3Niy/3O2 compounds were synthesized via a simple sol gel method. In order to achieve a pure P2-type structure, the materials were calcinated at different temperatures



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ranging from 700°C to 950°C, with a heat rate of 2°C/min, where the heat treatment was maintained for a duration of 12 hours. Figure 1 represents the XRD spectra of the solid-solution end members NaxCoO2 (y = 0) and NaxMn2/3Ni1/3O2 (y = 1.0) and the intermediate compositions NaxCo2/3Mn2/9Ni1/9O2 (y = 1/3), NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2), and NaxCo1/3Mn4/9Ni2/9O2 (y = 2/3), heat-treated at different temperatures. The XRD patterns of all samples heat treated at 700°C and of samples with (y = 0, 1/3, 1/2) at 950°C are given in Figure S2. While NaxCoO2 adopts the P2-type structure at all the chosen calcination temperatures, NaxMn2/3Ni1/3O2 adopts a P3-type structure at 700°C, and a mixture of P2 and P3 structures at 800°C and 850°C, along with the presence of an impurity peak at 2θ= 12.6°, which indicates the existence of a hydrated phase. Pure P2NaxMn2/3Ni1/3O2 was then obtained at the highest temperature of 950°C. On the other hand, XRD results of samples with a mixture of Co, Mn, and Ni were sensitive to the calcination temperature. At 700°C, the samples all adopt a P3-type structure, with the presence of small peaks of P2-type structure for the sample with high Co content (y = 1/3) (Figure S2), and at 800°C, the sample with y = 1/3 adopts mainly a P2 type structure with a small amount of P3 structure. Samples with more Mn and Ni (y = 1/2 and 2/3) adopt mainly a P3-type structure and have a small amount of P2. These results reveal that the substitution of 1-y Co with y (Mn2/3Ni1/3) increases the temperature of stability of the P2 structure. At 850°C, pure P2 phases were obtained for the three compositions, and at 900°C, some NiO impurities were identified in the compositions for y = 2/3 and ½, and their peaks slightly increased at 950°C (Figure S2). The structures obtained for all the compositions at different temperatures are summarized in Table S1. For

the

upcoming

experiments,

the

studied

materials

are

P2-type

NaxCoO2,

NaxCo2/3Mn2/9Ni1/9O2, NaxCo1/2Mn1/3Ni1/6O2, and NaxCo1/3Mn4/9Ni2/9O2 heat treated at 850°C and NaxMn2/3Ni1/3O2 heat treated at 950°C.


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NaxMn2/3Ni1/3O2 (y = 1)

950°C P2

NaxCoO2 (y = 0)

900°C P2 900°C P2 850°C P2

•

850°C P3+ ▪ P2

10

▪

• 20

30

40

P3(108)

800°C

50

800°C P2

P2(106) P2/P3(110) P2(112)/P3(113)

P2(004)/P3(006) P2(100)/P3(101) P3(102) P2(102) P2(103) P3(015) P2(104)

•

▪

•

P3(107)

• P2(002)/P3(003)

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


60

70

80

10

20

30

NaxCo1/3Mn4/9Ni2/9O2 (y = 2/3)

NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2)

* *

1.

NaxCo2/3Mn2/9Ni1/9O2(y = 1/3)

900°C P2

50

850°C P2

.

800°C

800°C

P2 +

P3+ ▪ P2 ▪ ▪ ▪ ▪

60

70

80 10

2-ThetaCu ( ° )

Figure

80

P2

▪ ▪ ▪ ▪ 40

70

850°C

800°C P3+ ▪ P2

30

60

* *

850°C P2

20

50

900°C P2+ * NiO

900°C P2+ * NiO

10

40

2-ThetaCu ( ° )

2-ThetaCu ( ° )

Powder

20

30

40

•

50

60

70

80 10

diffraction

patterns

30

40

. 50

60

70

80

2-ThetaCu ( ° )

2-ThetaCu ( ° )

X-ray

20

P3

for

NaxCo(1-y)Mn2y/3Niy/3O2

(y = 0, 1/3, 1/2, 2/3, 1) heat treated at different temperatures.“ •” hydrated phase. It should be noted that the heating rate is important because it can impact the final structure of the compound. For instance, a pure P2-NaxCo2/3Mn2/9Ni1/9O2 phase was obtained at 800°C, using faster heat rate (5°C/min).26 However, with the same rate, we couldn’t achieve a pure 9 ACS Paragon Plus Environment


P2 phase for the other compositions (y=1/2 and 2/3), for which both P2 and P3 phases were detected at 800°C while NiO impurities were identified at 850°C. XRD pattern and Rietveld-refined results of the as-prepared materials are depicted in Figure 2. The diffractograms clearly show a single phase. All Bragg diffraction lines indicate that P2NaxCo(1-y)Mn2y/3Niy/3O2 solid solution compounds crystallize in the hexagonal layered structure (P2-type structure) with the space group P63/mmc. The successful synthesis of the solid solution is confirmed also by plotting the cell’s volume as a function of the composition y, which shows a linear variation obeying Vegard’s law (Figure S3).

Measured Calculated Difference Bragg positions


NaxCoO2(y = 0)

Intensity (A.U)

NaxMn2/3Ni1/3O2(y = 1)

10

20

30

40

50

60

70

80

90

10

20

30

2-ThetaCu ( ° )

NaxCo1/3Mn4/9Ni2/9O2 (y = 2/3)


40

50

60

70

80

90

2-ThetaCu ( ° )

NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2)


20

70


NaxCo2/3Mn1/3Ni1/6O2 (y = 1/3)

Intensity (A.U)

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

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10

20

30

40

50

60

70

80

90 10

30

2-ThetaCu ( ° )

Figure

2.

Rietveld

40

50

60

80

90 10

refinement

of

20

30

40

50

60

70

80

90

2-ThetaCu ( ° )

2-ThetaCu ( ° )

as-prepared

P2-NaxCo(1-y)Mn2y/3Niy/3O2

(y = 0, 1/3, 1/2, 2/3, 1): red, observed; black, calculated; blue, difference plot; green bars, Bragg reflections. The refined parameters of the compounds are summarized in Table 1. The transition metals are located in the 2a site (0,0,0). The oxygen is in the 4f site (2/3,2/3,z) with z around 0.08 (± 0.002) for all the compositions except P2-NaxCoO2 where it is equal to 0.072 using the 10 ACS Paragon Plus Environment

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Rietveld refinement. The sodium is located in the 2b site (0,0,1/4), where the prismatic NaO6 share faces with the MO6 octahedra (Naf), and in the 2d site (2/3,1/3,1/4), where it only shares edges (Nae). The Rietveld refinement gives a Na content of approximately 0.63 (±0.02) for the following samples: NaxCoO2, NaxCo1/2Mn1/3Ni1/6O2,

NaxCo1/3Mn4/9Ni2/9O2, and

NaxCo2/3Mn2/9Ni1/9O2. For NaxMn2/3Ni1/3O2, the sodium content was higher (xNa = 0.72). The ratio of sodium content in both sites (Naf/Nae) was found to be slightly different, depending on the composition of the transition metals, where the Naf sites of samples with mixture of Co, Mn and Ni was relatively more filled (more stable) with a ratio of ≈ 0.7. Table 1. Crystallographic parameters of P2-NaxCo(1-y)Mn2y/3Niy/3O2 (y = 0, 1/3, 1/2, 2/3, 1) compounds refined by the Rietveld method. y in P2-NaxCo(1-y)Mn2y/3Niy/3O2

0

1/3

1/2

2/3

1

P63/mmc

Space group Cell parameters Profile parameters

a(Å) c(Å) U V W

2.8247(6)

2.8261(5)

2.8325(3)

2.8459(1)

10.943(5)

11.049(4)

11.082(4)

11.116(2)

0.04(1) -0.039(7)

0.09(2) -0.058(1)

0.013(1)

0.020(1)

0.11(2) -0.064(1) 0.013(2)

0.11038 0.11(3) -0.060(4) -0.06407 0.011(2) 0.01336

2.8841(7) 11.136(6) 0.2(6) -0.08(3) 0.015(4)

Refined coordinates

z (O)

0.073(1)

0.082(7)

0.079(2)

0.078(9)

0.083(7)

Occupation of Na sites

Naf Nae ratio

0.24(5) 0.39(3) 0.62

0.27(2) 0.38(1) 0.71

0.26(2) 0.36(2) 0.72

0.25(9) 0.36(9) 0.69

0.29(3) 0.43(1) 0.67

Conventional Rietveld R-factors

Rwp% RB

17.5 4.35

10.3 2.42

7.32 1.98

10.5 2.17

14.5 6.08

3.2.

Impact of transition metal composition on valence state

Figure 3 presents XANES spectra of the Co, Mn, and Ni K-edge for P2-NaxCo(1y)Mn2y/3Niy/3O2

pristine electrodes. The shape of the K-edge XANES spectra of the transition 11 ACS Paragon Plus Environment


metal oxides gives information about the site symmetry and the nature of the bonding of the probed atom with its surrounding ligands, while the energy position of the absorption edge provides information about its oxidation state. Ni K-edge

Co K-edge

b)

7726

7728

7730

7732

7734

Energy (ev)

NaxCoO2 (y=0) 7706

7708

7710

7712

7714

NaxCo2/3Mn2/9Ni1/9O2 (y=1/3)

Norm. Absorbance (a.u)


a)

8344

8338

8340

7710

7720

7730

7740

7750

8342

8344

NaxCo2/3Mn2/9Ni1/9O2 (y=1/3) NaxCo1/3Mn4/9Ni2/9O2 (y=2/3) NaxMn2/3Ni1/3O2 (y=1)

7760

8330

8340

NaxCo2/3Mn2/9Ni1/9O2 (y=1/3)

NaxCo1/3Mn4/9Ni2/9O2 (y=2/3) NaxCo1/2Mn1/3Ni1/6O2 (y=1/2) NaxMn2/3Ni1/3O2 (y=1)

6540

6550

6560 Energy (ev)

6570

6580

"c" parameter (Å) "a" parameter (Å)

6558 6559 6560 6561 6562 6563 Energy (ev)

6530

8350

8360

8370

8380

Energy (eV)

c)

6544

8360

NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)

Mn K-edge

6540 6542 Energy (ev)

8356

8346

Energy (ev)

6538

8352

Energy (ev)

NaxCo1/3Mn4/9Ni2/9O2 (y=2/3)

7700

8348

Energy (ev)

NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)

Energy (ev)


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

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6590

11,15

d)

11,10 11,05 11,00 10,95 10,90 2,880 2,865 2,850 2,835 2,820

0,0

0,2

0,4

0,6

0,8

1,0

y in P2-Na2/3Co(1-y)Mn2y/3Niy/3O2

Figure 3. XANES spectra at (a) Co K-edge, (b) Ni K-edge, and (c) Mn K-edge for P2NaxCo(1-y)Mn2y/3Niy/3O2 (y = 0, 1/3, 1/2, 2/3, 1) pristine electrodes. (d) Unit cell parameters as a function of y in P2-NaxCo(1-y)Mn2y/3Niy/3O2 (y = 0, 1/3, 1/2, 2/3, 1). As shown, all the compositions present the same shape: the weak pre-edge absorption peak at low energies represents the electric dipole-forbidden transition in an ideal octahedral symmetry of a 1s electron to an unoccupied 3d orbital. The intensity of this peak is due to pure electric quadrupole coupling and/or 3d-4p orbital mixing arising from the


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noncentrosymmetric environment of slightly distorted MO6. The main absorption edge is due to the electric dipole-allowed transition of a 1s core electron to an unoccupied 4p bound state. The shoulder in the absorption spectra in the lower energy region with respect to the main absorption edge is assigned to a shakedown process due to a ligand-to-metal charge transfer involving the 1s-4p transition.28-31 A weak pre-edge peak at 7710 eV is present in the Co K-edge spectra (Figure 3a), indicating that the local environment of Co is not totally symmetric. As the Co content decreases in the material (y = 1/3, 1/2, 2/3), the features of the Co K-edge XANES spectrum remain unchanged. However, the main absorption edge shifts toward higher energy values. Cobalt ions in NaxCoO2 have a higher oxidation state than +3. The introduction of Mn and Ni in P2NaxCo(1- y)Mn2y/3Niy/3O2 reduces the Co content, leading to a decrease of the Co4+ formation rate. The main absorption edge of NaxCoO2 and NaxCo2/3Mn2/9Ni1/9O2 represents higher energy values (7729.7 eV) compared to NaxCo1/2Mn4/9Ni2/9O2 and NaxCo1/3Mn1/3Ni1/6O2 (7729.3 eV) since they contain Co4+ ions. This finding is in good agreement with our earlier investigations on the composition NaxCo2/3Mn2/9Ni1/9O2 using hard X-ray photoelectron spectroscopy (HAXPES).32 The

Ni

K-edge

XANES

spectra

of

NaxCo2/3Mn2/9Ni1/9O2,


NaxCo1/3Mn4/9Ni2/9O2, and NaxMn2/3Ni1/3O2 (Figure 3b) exhibit very weak pre-edge peaks, which means that the local environment of Ni atoms is rather symmetric. The main absorption peak represents a binding energy of 8352.8 eV for NaxCo2/3Mn2/9Ni1/9O2, which decreases gradually with the increase of Ni content to reach 8350.2 eV for NaxMn2/3Ni1/3O2. This shift is accompanied by a significant change in the peak’s shoulder that becomes sharper for NaxMn2/3Ni1/3O2 (8343.37 eV). All of these changes reveal that a lower oxidation state of Ni occurred by increasing the value of y. Since Ni ions in NaxCo2/3Mn2/9Ni1/9O2 are in the +3 valence state,32 the shift of the energy is induced by the introduction of Ni2+ ions for 13 ACS Paragon Plus Environment


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NaxCo1/2Mn1/3Ni1/6O2 and NaxCo1/3Mn4/9Ni2/9O2 and the total conversion to Ni2+ ions for NaxMn2/3Ni1/3O2. The

Mn

XANES

spectra

of



NaxCo1/3Mn4/9Ni2/9O2, and NaxMn2/3Ni1/3O2 also have the same shape (Figure 3c). The preedge region consists of two peaks, which are representative of Mn4+, and a single peak in the pre-edge region, characteristic of Mn3+.32,33 However, a small shift of the main absorption peak toward lower energies (≈0.4 eV) is evident for NaxMn2/3Ni1/3O2. As reference samples, we used pure Mn2O3 (Mn3+) and MnO2 (Mn4+) oxides. Their Mn K-edge XANES spectra are plotted in Figure S4. The difference between the main absorption peaks is ≈ 3.5 eV. Since the energy difference of the Mn K-edge main peak between y = 1/3 and y = 1 is small, and the two pre-edge peaks are preserved, the noticed shift might be explained by the presence of a very small amount of Mn3+ ions in NaxMn2/3Ni1/3O2 originating from the higher sodium content found in the material, as determined by the Rietveld refinement (xNa = 0.72). In addition, nonstoichiometric oxygen is expected in this type of material when synthesized in air. The valence state of the transition metals will be affected accordingly. Nonetheless, in the present study, we consider oxygen to be stochiometric for all the samples. The evolution of the unit cell parameter “a”, corresponding to the M-M distance, as a function of the composition (y) confirms the XAS results since this crystallographic parameter is mainly affected by the oxidation state of the transition metals. As presented in Figure 3d , only a small increase of the “a” parameter is noticed from y = 0 to y = 1/3. Indeed, the substitution of Co4+ ions (0.053 nm), present in NaxCoO2 (NaxCo3+Co4+(1-x)O2), by Mn4+ ions (0.053 nm) and low-spin Ni3+ ions (0.056 nm) having comparable ionic radii weakly influences the “a” unit cell parameter in this composition range. A gradual increase of “a” occurs between y = 1/2 and y = 2/3. This increase is caused by the presence of Ni2+ (y = 2/3) with higher ionic radii (0.64 nm). For y = 1, corresponding to the presence of only Mn and Ni


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ions, the parameter “a” reaches the highest value, 2.89 Å, due to the presence of 1/3 Ni2+ ion in the material. Figure 3d also presents the evolution of the unit cell parameter “c”, related to the interlayer distance. This hexagonal unit cell parameter shows a sharp increase from y = 0 to y = 1/3, then gradually increases to y =1. Even if both Mn and Ni ions are introduced, we believe that this increase in “c” is mainly due to the introduction of Mn wich has lower electronegativity, causing high ionicity of the M-O bonds. A negative charge is then accumulated on the oxygen ions, enhancing the coulombic repulsion between MO2 layers. This mechanism explains the increase of the “c” parameter with increasing y. The composition NaxMn2/3Ni1/3O2 (y = 1) has the largest interlayer distance, making it more vulnerable to the introduction of H2O atoms (present in air) between the sodium layers. This explains the presence of the hydrated phase in the XRD patterns of this material when calcinated at temperatures below 950°C (Figure 1). 3.3.

Morphological characterization

Figure S5 presents SEM images of P2-NaxCo(1-y)Mn2y/3Niy/3O2 samples. All the samples have flake-shaped particles, and regardless of the y value, they are not homogeneously distributed, with sizes ranging between 1 and 3 µm. The addition of NH3·H2O during the sol-gel process until the solution is at pH 7 was found to improve the homogeneity of the particles and, thus, the cycling stability.34 However, an attempt to increase the pH to 7, in our case, resulted in the appearance of a small impurity of nickel oxide in the XRD patterns. In addition, the materials studied delivered remarkably good electrochemical performance, despite the heterogeneity among the particles, as discussed in the next section. 3.4.

Electrochemical properties

Figure 4 a and b show the initial charge-discharge voltage profiles of NaxCoO2 (y = 0), NaxCo2/3Mn2/9Ni1/9O2

(y = 1/3),

NaxCo1/2Mn1/3Ni1/6O2

(y = 1/2),

NaxCo1/3Mn4/9Ni2/9O2 15

ACS Paragon Plus Environment


(y = 2/3), and NaxMn2/3Ni1/3O2 (y = 1) measured between 2.0 V and 4.2 V vs. Na/Na+ at a C/20 current rate.

4,0

NaxCoO2 (y=0)

4,0

3,6

+

Potential (V vs Na/Na )

3,2 NaxCoO2 (y=0)

NaxCo2/3Mn2/9Ni1/9O2 (y=1/3)

2,8

NaxCo1/2Mn1/3Ni1/6O2 (y=1/2) NaxCo1/3Mn4/9Ni2/9O2 (y=2/3)

2,4

a)

NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)

3,6

NaxCo1/3Mn4/9Ni2/9O2 (y=2/3) NaxMn2/3Ni1/3O2 (y=1)

3,2

2,8

2,4

b)

NaxMn2/3Ni1/3O2 (y=1)

2,0

2,0 0

20

40

60

80

100

0

20

c)

2,1

2,4

2,7

3,0

3,3

3,6

3,9

40

60

80

100

120

Specific discharge capacity(mAhg )

4,2

y=0

d) 200

y=1/3

100

y=2/3

y=1

95

160 140

90

-1

120

mAhg

Specific discharge capacities

180

y=1/2

100

85

80 80

60

NaxCoO2 (y=0)

NaxCo1/3Mn4/9Ni2/9O2

40

NaxCo2/3Mn2/9Ni1/9O2 (y=1/3)

20

NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)

(y=2/3)

0 2,4

2,7

3,0

3,3

3,6

3,9

10

20

30

4,2

+

40

50

75


0 2,1

140

-1

-1

Specific charge capacity(mAhg )

Coulombic efficiency %

+


NaxCo2/3Mn2/9Ni1/9O2 (y=1/3)

dQ/dV

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

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60

70

80

90

70 100

Cycle number


Figure 4. First charge (a) and discharge (b) capacities versus voltage of P2-NaxCo(1y)Mn2y/3Niy/3O2 (y

= 0,1/3, 1/2, 2/3, 1) cycled between 2 V and 4.2 V vs. Na+/Na at C/20 rate.

(c) dQ/dV plots of the second cycle and (d) evolution of the specific discharge capacity and coulombic efficiency with cycling for same materials. As shown in Figure 4b, NaxCoO2 (y=0) delivered a specific discharge capacity of 132 mAh g-1, which decreased gradually with lower Co content for the samples with a mixture of Co, 16 ACS Paragon Plus Environment

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Mn, and Ni. The capacities were 113.4 mAh g-1 for y = 1/3, 100.4 mAh g-1 for y = 1/2, and 90 mAh g-1 for y = 2/3, while the specific discharge capacity delivered by NaxMn2/3Ni1/3O2 (y = 1) was 101 mAh g-1. The NaxCoO2 exhibited 10 reversible plateaus in the studied potential window, that are clearly observed as peaks in the derivative capacity plot (dQ/dV) presented in Figure 4c. This figure also shows distinct potential jumps around 2.8 V, 3.4 V, and 4.0 V corresponding to the compositions Na2/3CoO2, Na1/2CoO2, and Na1/3CoO2, respectively. These potential jumps were found to be related to the presence of ordered phases, where additional diffraction lines in the XRD pattern for the two compositions Na2/3CoO2 and Na1/2CoO2 were reported, corresponding to the √3a×2a ×c supercell for Na1/2CoO2 and the √12a×√12a ×3c supercell for Na2/3CoO2.14,35,36 As also shown in Figure 4c, NaxMn2/3Ni1/3O2 (y = 1) presents four distinct peaks in the dQ/dV curve. The samples with a combination of Co, Mn, and Ni ions have only three peaks, and their intensity decreases with increasing x, which is manifested in lower specific capacity values. As discussed in section 3.1, the Naf site is more stable in these samples. Therefore, the substitution of Co with Mn and Ni helps stabilizing the structure by reducing Na+ ion/vacancy and thus minimizing the voltage steps occurring during cycling. The 2.8-V (i.e., xNa ≈ 0.66) potential jump noticed in the voltage profile of NaxCoO2 (y = 0) disappears for all the other samples, unlike the one present at around 3.4 V corresponding to a sodium composition of 0.5 that is noticed for all the samples. However, the sharpness of the latter dramatically decreases with decreasing Co content, declining from a voltage difference of ∆V ≈ 0.35 V for NaxCoO2 (y = 0) to ∆V ≈ 0.2 V for NaxCo2/3Mn2/9Ni1/9O2 (y = 1/3). In fact, only samples with high Co content (y = 0 and y = 1/3) present a distinct voltage jump for xNa ≈ 0.5, and according to our XANES measurements, they both have Co4+ ions in the (MO2)n layers. The same behavior was observed for the 4.0-V potential jump, and it was more noticeable in samples with higher Co content. The samples with y = 2/3 and 1 clearly indicate



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Page 18 of 37

the start of a new plateau at 4.2 V. These phenomena are caused by Na ordering driven by coulomb interactions among Na ions.37-39 However, it is clear that these ordering phases are also strongly related to the electronic composition of the (MO2)n layers, where the presence of Co4+ appears to accentuate the intensity of the potential jumps.40 The start of the ≈ 3.4 V potential jump is different for each sample: it appears after an extraction of 0.14, 0.12, and 0.10 sodium, respectively, for P2-NaxCoO2 (y = 0), P2NaxCo2/3Mn2/9Ni1/9O2 (y = 1/3), and P2-NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2). For P2NaxCo1/3Mn4/9Ni2/9O2 (y = 2/3) and NaxMn2/3Ni1/3O2 (y = 1), however, it is difficult to point out the start of the potential jump. The slight shift might be explained, simply, by a difference in the starting sodium content. In addition to the reaction of these materials with air, they react with the binder (PVDF) to form NaF during the electrode manufacturing,32 making it difficult to identify the exact sodium content. Another explanation is the different distribution of sodium in the two prismatic sites found from the Rietveld refinement for each composition. Table S2 shows the calculated values of the sodium at the 3.4-V potential jump based on the Rietveld results (for the pristine materials), assuming that the sodium extraction is mainly conducted from the Nae site at this potential.41 Despite the slightly different sodium content at each composition, the 3.4 V potential jump seems to correspond to when the two sites have similar sodium content. Figure 4d displays the evolution of the specific discharge capacity and the coulombic efficiency for the P2-NaxCo(1-y)Mn2y/3Niy/3O2 compounds in the voltage range of 2.0 and 4.2 V vs. Na/Na+ at a rate of C/20. Table 2 presents the values measured at the 1st, 25th, 50th, and 100th cycles. At 25 cycles, the capacity slightly increased for most of the compositions, and the coulombic efficiency increased more significantly for all the compositions, where it was above 99% for samples with y = 0, 1/3, and 1/2 and 98.8% for y = 2/3 and 1, and stabilized in the cycling afterwards. However, the coulombic efficiency of NaxCoO2 decreased slightly 18 ACS Paragon Plus Environment

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from 99.5% in cycle 20 to 98.8% in cycle 50. On the other hand, the capacity for NaxMn2/3Ni1/3O2 increased slightly from 100.8 mAh g-1 to 102 mAh g-1 in the 25th cycle, then dramatically decreased afterwards to only 93.92 mAh g-1 in the 50th cycle. The increase of the capacity might be due the uptake of H2O present in air, as evidenced by the fact that NaxMn2/3Ni1/3O2, unlike the other compositions, presented an impurity peak in the XRD pattern corresponding to the presence of a hydrated phase at relatively high temperatures (800°C, 850°C, and 900°C), and it only disappeared at 950°C. At 50 cycles, the tested materials offer very good capacity retention, exceeding 97% for y = 0, 1/3, 1/2, and 2/3 and 92% for y = 1. At 100 cycles, the capacity retention of the samples with y = 0 and 1/3 is still good (84% and 87%, respectively) but lower than samples with y = 1/2 and 2/3 (more than 96%). Indeed, the substitution of Co with Mn and Ni (y = 1/3 and 1/2) considerably decreased the specific discharge capacity of the materials, but it also increased remarkably their capacity retention and the stability of the coulombic efficiency during cycling. Of the samples tested, NaxMn2/3Ni1/3O2 delivered the lowest electrochemical performance, indicating that the existence of Co beneficially affected both the specific capacity and its retention during cycling. Table 2. Specific discharge capacity and coulombic efficiency values of P2-NaxCo(1y)Mn2y/3Niy/3O2 compounds

at the 1st, 25th, 50th, and 100th cycles.

y in

0

1/3

1/2

2/3

1

Specific discharge capacity (mAhg-1)

2nd cycle 25th cycle 50th cycle 100th cycle

132.07 132.53 128.58 110.6

113.42 112.66 110.53 100.48

101.04 100.52 100.5 97.93

89.59 90.17 90.13 86.97

100.83 102.01 93.92 -

Coulombic efficiency (%)

2nd cycle 25th cycle 50th cycle 100th cycle

98.89 99.50 98.64 98.82

95.93 99.12 99.02 99.02

96.79 99.31 99.45 99.27

92.77 98.85 98.83 98.80

89.52 98.79 98.73 -

P2-NaxCo(1-y)Mn2y/3Niy/3O2



½ P2-NaxCoO2-½ P2-NaxMn2/3Ni1/3O2 (y = 1/2) delivered the best electrochemical performance in long-term cycling, with a specific discharge capacity of 101 mAh g-1 and a loss of less than 3% after 100 cycles. This good performance is due to the coexistence, in the layered transition metal framework, of Co3+ that enhances the Na+ mobility, Mn4+ that stabilizes the structure, and a Ni2+/Ni3+ mixture that provides high electronic conductivity. In order to check the stability of this material at high potential, a galvanostatic test was performed at the C/20 rate from 2.0 V to 4.5 V. Figure 5 presents the voltage profiles of NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2) and the end members of the solid solution [NaxCoO2 (y = 0), NaxMn2/3Ni1/3O2 (y = 1)], as well as the evolution of their specific discharge capacity and their coulombic efficiency during the first 20 cycles. 4,5

4,5

4,0

NaxCoO2 (y=0) NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)


4,0


+

+


3,5

3,0 NaxCoO2 (y=0) NaxCo1/2Mn1/3Ni1/6O2 (y=1/2)

2,5

3,5

3,0

2,5


b)

a) 2,0

2,0

0

20

40

60

80

100

120

140

160

180

0

20

-1

40

60

80

100

120

140

160

-1

Specific discharge capacity (mAh/g )

Specific charge capacity (mAhg ) 200

c)

100

180 160

120 -1

60

100 80

40

60

Na xCoO 2 (y=0) NaxCo1/2Mn1/3Ni1/6O 2 (y=1/2) NaxMn2/3Ni1/3O 2 (y=1)

40 20

20

0

Coulombic efficiency %

80

140

mAhg

Specific discharge capacities

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

Page 20 of 37

0 0

2

4

6

8

10

12

14

16

18

20

Cycle number

Figure 5. (a,b) First charge and discharge capacities versus voltage profiles of P2NaxCo(1-y)Mn2y/3Niy/3O2 (y = 0, 1/2, 1) cycled between 2.0 V and 4.5 V vs. Na+/Na at C/20


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rate. (c) Evolution of the specific discharge capacity and coulombic efficiency with cycling for same materials. High capacity values are achieved when cycling the materials up to 4.5 V. The specific charge capacities were 145.45 mAh g-1, 152.27 mAh g-1, and 166.78 mAh g-1 for NaxCoO2 (y = 0), NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2), and NaxMn2/3Ni1/3O2 (y = 1), respectively. This performance is due to the appearance of a new plateau; its extent accounts for about 55.45 mAh g-1, 77.27 mAh g-1, and 81.77 mAh g-1 for compositions y = 0, 1/2, and 1, respectively. This plateau is more pronounced for NaxMn2/3Ni1/3O2 (y = 1) and starts also at lower voltages in comparison with samples with Co (y = 0 and 1/2). However, this plateau is not totally reversible during the first discharge. The specific discharge capacities delivered are 139.90 mAh g-1 and 144.74 mAh g-1for NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2) and NaxMn2/3Ni1/3O2 (y = 1), respectively. On the other hand, despite a specific discharge capacity of 136.57 mAh g-1, NaxCoO2 (y = 0) completely lost its voltage profile during the first discharge. After the first cycle, the discharge capacities decreased gradually. After 20 cycles a capacity retention of 94% and 88% and a coulombic efficiency of 97.6% and 98.4% were obtained for NaxMn2/3Ni1/3O2 (y = 1) and NaxCo1/2Mn1/3Ni1/6O2 (y = 1/2), respectively. While NaxCoO2 (y = 0) delivered better electrochemical performance than NaxMn2/3Ni1/3O2 (y = 1) when cycled between 2.0 V and 4.2 V vs. Na+/Na, no reversibility was possible when the sample was cycled up to 4.5 V vs. Na+/Na, leading to a very low capacity retention of only 28% after 20 cycles. Therefore, the substitution of Co with (Mn,Ni) seems to increase the capacity stability at high voltages. In conclusion, despite the higher capacity values achieved, the capacity retention was lower when we cycled the materials up to 4.5 V vs. Na+/Na with the appearance of a new plateau. The latter is well known to be associated with a phase transition, when sodium content is lower than ≈ 0.35, from P2 to O2,8,42 or to OP4,10 or to ‘Z’,11,43 types and to return to the P2-type structure during sodium insertion. Since we had cycled the three 21 ACS Paragon Plus Environment


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samples at high voltages, electrolyte decomposition is one of the main causes of the irreversibility. Although each sample might react differently with the same electrolyte, we believe that the strange behavior of NaxCoO2 (y = 0) is caused by a much more complex phenomenon and more studies need to be addressed to understand the structural evolution of this material at high voltage. 3.5.

In-situ XANES measurements of ½ P2-NaxCoO2-½ P2-NaxMn2/3Ni1/3O2

To understand the evolution of transition metals during cycling in the two potential windows studied,

we

performed

in-situ

XANES measurements for ½

P2-NaxCoO2-½ P2-

NaxMn2/3Ni1/3O2, with the first cycle between 2 V and 4.2 V and the second between 2.0 V and 4.5 V.

3.5.1. First cycle: 2.0-4.2V window Figure 6 presents Co, Mn, and Ni K edge XANES spectra for the first cycle during charge and discharge, and the insets present their main peak edge at the most significant voltages. During the first cycle, no significant change is noticed in the shape of Co, Mn, and Ni K-edge spectra features, only a shift of the main peak edge corresponding to oxidation (sodium extraction) or reduction (sodium insertion) of the transition metal ions. During charge, the Co K-edge main peak shows a continuous shift toward higher energy values. The energy difference between OCV and 4.2 V is approximately 0.8 eV, meaning that Co3+ oxidized to Co4+. Likewise, the Ni K-edge XANES main peak shows an energy difference of around 0.84 eV between OCV and 4.2 V; however, this shift is observed mainly between OCV and 3.7 V, after which the main Ni K-edge peak stays constant (see inset for Ni K-edge in Figure 6). Since the pristine material is mainly Ni3+ (Section 2), we believe that total oxidation of Ni ions to Ni4+ occurs during charge. On the other hand, Mn4+ ions are stable all along the charge process.


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During discharge, the main peak of the Co K-edge XANES spectra undergoes a continuous shift toward lower energy values, which stops around 2.5 V (see inset to Co K-edge in Figure 7). Meanwhile, the Ni K-edge main peak only started shifting toward lower energy values after 3.5 V (inset to Ni K-edge in Figure 6). According to our electrochemical results, more sodium ions are inserted into the P2-NaxCo1/2Mn1/3Ni1/6O2 during the first discharge (0.4 Na+) then are extracted during the first charge (0.3 Na+), since we cycled it beyond its OCV value. This allows the material to deliver higher discharge capacity during the first cycle, and the difference needs to be compensated by reduction of transition metals beyond their first state. Based on our in-situ XANES results, the Co ions regained their initial valence state, but Ni ions were reduced to a lower valence state, since the Ni K-edge main peak slightly shifted to a lower energy value (by 0.3 eV) than the one recorded for OCV, suggesting that Ni3+ reduced to Ni2+ at lower voltages. The Mn tetravalent ions were also reduced to a trivalent state after 3.0 V (i.e., OCV value), where the main peak in the Mn K-edge showed an energy difference of 0.5 eV between 3.0 V and 2.0 V (inset to Mn K-edge in Figure 6). Therefore, the extra capacity recorded during the first discharge is mainly due to simultaneous reduction of tetravalent Mn ions along with reduction of Ni ions to the +2 valence state.



7726 7728 7730 7732 7734

Energy (ev)

7710

7720

7730

OCV 4,2V

6558 6559 6560 6561 6562 6563

Energy (ev)

7750 6530

7740

6540

6550

6560

6570

6580

4,2V 2,5V 2,0V

4,2V 2,0V

7726 7728 7730 7732 7734

8330

8340

7740

7750 6530

6540

8350

8360

8370

4,2V 3,5V 2,0V

4,2V 2,0V

Energy (ev)

Energy (ev)

7730

Energy (ev)

8348 8350 8352 8354 8356

6558 6559 6560 6561 6562 6563

Energy (eV)

7720

8348 8350 8352 8354 8356

Energy (ev)

4,2V 3,0V 2,0V

4.2V 2.0V

7710

OCV 4,2V

Energy (ev)

Energy (ev)


OCV 3,7V 4,2V

OCV 4,2V

OCV Ch4.2V


OCV 3,3V 4,2V

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Mn K-edge

Norm. A bsorbance (a.u)

N orm . A bso rban ce (a.u)

Charge

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Discharge

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Figure 6. In-situ XANES spectra at Co, Mn, and Ni K-edge for P2-NaxCo1/2Mn1/3Ni1/6O2 cycled between 2.0 V and 4.2 V.

3.5.2. Second cycle: 2.0-4.5V window The Co, Mn, and Ni K edge XANES spectra recorded for the second cycle during both charge and discharge are presented in Figure 7, and the inset plots show their main peak edge at the most significant voltage values. During charge, all transition metal ions participate in the oxidation process. The Mn K-edge main peak shifts toward higher energies up to 3.5 V, indicating that Mn ions regain their tetravalent state at low voltages. On the other hand, both Ni and Co K-edge main peaks start shifting toward higher energy values after 3.0 V, corresponding to the oxidation of Ni ions to Ni4+ and Co ions to Co4+. Cells charged to 4.5 V exhibited a very high capacity due to the new 24 ACS Paragon Plus Environment

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voltage plateau, where between 4.2 V and 4.5 V the material delivers a specific capacity of ≈ 80 mAh g-1, which corresponds to the extraction of 0.3 Na+. We expected that this sodium extraction would be compensated by the oxidation of Co3+ since it is the main transition metal in our material, with a composition of ½, and both Mn and Ni ions are already mainly oxidized at this potential (4.2 V). However, only a very small shift occurs after 4.2 V in the Co K-edge spectrum, and from 4.4 V to 4.5 V, which corresponds to a specific capacity of 55 mAh g-1 (i.e., 0.21 Na+), no shift is noticed. In fact, all the transition metal K-edge main peaks during full charge to 4.5 V have the same energy positions as during full charge to 4.2 V (Figure S6). In addition, as explained before, this high voltage plateau is due to a phase transition related to the sodium content, and a recent study on P2-Na2/3− z[Mn1/2Fe1/2]O2 showed that at high voltage, the transition metals migrate into a tetrahedral coordination environment.42 Our XANES measurements showed no significant change in the shape of the transition metal K-edge spectra features (pre-edge peak, shoulder, and main peak), indicating that the local structure is preserved in our case. Therefore, the origin of the capacity at this high voltage is probably the contribution of oxygen. In fact, the valence state of cobalt increases less for NaCoO2 than LiCoO2 during alkali metal extraction, and it stops increasing after a sodium composition of ≈ 0.4 for NaCoO2.44 During discharge, the main peaks of the Co, Mn, and Ni K-edge XANES spectra are stable between 4.5 V and 4.2 V (high voltage plateau). The Ni and Mn ions show the same behavior as during the first discharge between 4.2 V and 2.0 V, where the Ni ions are active below 3.6 V and Mn ions below 3.0 V, as evident in the insets to Figure 7. However, the main peak energy difference is lower than the one recorded during the first discharge, with an energy difference of only 0.12 eV for Ni ions and 0.16 eV for Mn ions. On the other hand, the Co Kedge main peak shows a very small shift from 4.2 V to 2.5 V, in contrast to what we found during the first discharge from 4.2 V to 2.0 V. Despite a high specific discharge capacity of 25 ACS Paragon Plus Environment


140 mAh g-1, which corresponds to the insertion of 0.55 Na+, all the transition metals show very little change in their oxidation degree and are stable in the high voltage plateau. This behavior of transition metals during discharge strengthens the hypothesiss of contribution of oxygen to the electrochemical process. This hypothesis is also supported by the decomposition of sodium carbonates found at the surface of the pristine electrodes during charge and their dissolution after 4.2 V to form again during discharge, as observed in our recent study with P2-NaxCo2/3Mn2/9Ni1/9O2 using the HAXPES technique.

Their

counterparts

(lithium

carbonates)

decompose

due

to

electrochemical oxidation, but they are not reversibly formed by electrochemical reduction, suggesting that they can be formed as a byproduct of the oxygen reduction reaction.45 In fact, this hypothesis has been confirmed for Na0.78Ni0.23Mn0.69O2 by Ma et al.46, where they have found oxygen vacancies at the surface of the material.

Mn K-edge

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Ni K-edge

2,0V 3,5V 4,2V 4,5V

2,0V 4,5V

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Norm . A bsorbance (a.u)

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Figure 7. In-situ XANES spectra at Co, Mn, and Ni K-edge for P2-NaxCo1/2Mn1/3Ni1/6O2 cycled between 2.0 V and 4.5 V. 4. Conclusions (1-y) P2-NaxCoO–y P2-NaxMn2/3Ni1/3O2 (y = 0, 1/3, 1/2, 2/3, 1) compounds were successively synthesized via a simple sol gel method. We found that changing the composition of the transition metals in samples synthesized by the same method, having the same starting sodium content and treated under the same conditions, impacts the resulting structure. By changing the temperature of calcination from 700°C to 950°C, the substitution of 1-y Co with y (Mn2/3, Ni1/3) increased the temperature of stability of the P2-type structure. XANES measurements conducted on pristine electrodes revealed a higher oxidation state for the Co ions in samples with higher Co content (y = 0 and y = 1/3), confirming the presence of Co4+ ions. Manganese ions were found at the tetravalent state for the samples. On the other hand, the oxidation state of Ni ions decreased with increasing y (Mn2/3Ni1/3), confirming the introduction of Ni2+ ions. All the materials gave good electrochemical performance, when cycled between 2.0 and 4.2 V vs. Na+/Na at a rate of C/20. After cell testing for 50 cycles, the end members of the solid solution delivered good capacity retention: 97% for P2-NaxCoO2 (y = 0), with an initial discharge capacity of 132 mAh g-1, and 92% for P2-NaxMn2/3Ni1/3O2 (y = 1), with an initial discharge capacity of 100 mAh g-1. After 100 cycles, the samples with a mixture of Co, Mn, and Ni, especially the one with lower Co content (y = 1/2, 2/3), showed a remarkably higher stability, with a loss of only 3% of the initial capacity. Also, only one potential jump was observed in their voltage profiles, and its intensity decreased with decreasing Co content. Even though increasing y (Mn2/3Ni1/3) lowered the discharge capacity, it remarkably increased the capacity retention.



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Since the ½ P2-NaxCoO2-½ P2-NaxMn2/3Ni1/3O2 composition gave the best electrochemical performance, we cycled it up to 4.5 V vs. Na+/Na along with the end members of the solid solution for comparison. Higher capacity values were obtained due to the presence of the high voltage plateau, which is usually observed in P2-type structure layered materials. However, the three materials (y= 0, ½, 1) showed lower stability with cycling, where NaxCoO2 (y = 0) completely lost its voltage profile during the first discharge, and the capacity retention dropped to less than 94% for the other samples after only 20 cycles. The evolution of the valence state of the transition metals of NaxCo1/2Mn4/9Ni2/9O2 (y = 1/2) in these two potential windows was then studied by in-situ XANES measurements, where we found that between 2 and 4.2 V vs Na+/Na, Mn is activated when we discharge the material below its OCV value, both Mn and Ni ions are active at low voltages and Co ions are active along the potential window. At the high voltage plateau where the electrochemical reactions are expected to be mainly based on Co3+/Co4+ redox couple, all the transition metals are found stable and only small shifts in TM K-edge main peaks are noticed during discharge despite a discharge capacity of ≈ 140 mAhg-1. This new finding raises questions about the origin of the high voltage plateau and the phase transition that accompanies it and further studies are needed in order to understand and improve the stability of P2-type materials at high voltages. Acknowledgments We would like to acknowledge V. A. Oltean for assistance and fruitful discussion with Dr. K. Kubota. Argonne National Laboratory is operated by the U. S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. Use of the Advanced Photon Source (9-BM) was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under contract No. DE-AC02-06CH11357. We also would like to acknowledge StandUp for Energy – a Swedish Strategic Research area and the Swedish Research Council, under Contract No. 2015-05106. 28 ACS Paragon Plus Environment

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Supporting information. XRD patterns for NaxCo(1-y)Mn2y/3Niy/3O2 heat treated at 700°C and 950°C, Volume cell of P2-NaxCo(1-y)Mn2y/3Niy/3O2 as a function of the composition (y), SEM images of NaxCo(1-y)Mn2y/3Niy/3O2 , XANES spectra at Mn K-edge for Mn2O3 and MnO2, XANES spectra.



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Figures and tables captions Scheme 1. Illustration of O3, P2 and P3 structures. Figure 1. Powder X-ray diffraction patterns for NaxCo(1-y)Mn2y/3Niy/3O2 (y=0, 1/3, 1/2, 2/3, 1) heat treated at different temperatures.“ •” hydrated phase. Figure 2. Rietveld refinement of as-prepared P2-NaxCo(1-y)Mn2y/3Niy/3O2 (y=0, 1/3, 1/2, 2/3, 1): red, observed; black, calculated; blue, difference plot; green bars, Bragg reflections. Table1. Crystallographic parameters of P2-NaxCo(1-y)Mn2y/3Niy/3O2 (y=0, 1/3, 1/2, 2/3, 1)compounds refined by the Rietveld method. Figure 3. XANES spectra at (a) Co K-edge, (b) Ni K-edge, and (c) Mn K-edge for P2NaxCo(1-y)Mn2y/3Niy/3O2 (y=0, 1/3, 1/2, 2/3, 1) pristine electrodes. (d) Unit cell parameters as a function of y in P2-NaxCo(1-y)Mn2y/3Niy/3O2 (y=0, 1/3, 1/2, 2/3, 1). Figure 4. First charge (a) and discharge (b) capacities versus voltage of y)Mn2y/3Niy/3O2

P2-NaxCo(1-

(y=0,1/3, 1/2, 2/3, 1) cycled between 2 V and 4.2 V vs. Na+/Na at C/20 rate.

(c) dQ/dV plots of the second cycle and d) evolution of specific discharge capacity and coulombic efficiency with cycling for same materials. Table 2. Specific discharge capacity and coulombic efficiency values of P2-NaxCo(1x)Mn2x/3Nix/3O2 compounds

at the 1st, 25th, 50th, and 100th cycles.

Figure 5. (a,b) First charge and discharge capacities versus voltage profiles of P2-NaxCo(1y)Mn2y/3Niy/3O2

(y= 0, 1/2, 1) cycled between 2 V and 4.5 V vs. Na+/Na at C/20 rate. (c)

Evolution of the specific discharge capacity and coulombic efficiency with cycling for same materials. Figure 6. In-situ XANES spectra at Co, Mn, and Ni K-edge for P2-NaxCo1/2Mn1/3Ni1/6O2 cycled between 2.0 V and 4.2 V. 30 ACS Paragon Plus Environment

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[44] Valkeapää M., Katsumata Y., Asako I., Motohashi T., Chan T. S., Liu R. S., Chenc J.M., Yamauchia H., Karppinen M., Charge compensation and oxidation in NaxCoO2−δ and LixCoO2−δ studied by XANES, J. Solid State Chem., 2007, 180(5), 1608-1615. [45] Yabuuchi N., Yoshii K., Myung S. T., Nakai I., Komaba S., Detailed studies of a highcapacity electrode material for rechargeable batteries, Li2MnO3−LiCo1/3Ni1/3Mn1/3O2, J.

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Table of Contents Graphic


3)yO2 Cathode Materials for Sodium

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