Higher voltage, wider voltage plateau, longer cycle life, and faster kinetics via thermally modulated interfaces between Ramsdellite and Pyrolusite MnO2 for lithium-ion battery cathodes

Journal of Power Sources 450 (2020) 227619

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Higher voltage, wider voltage plateau, longer cycle life, and faster kinetics via thermally modulated interfaces between Ramsdellite and Pyrolusite MnO2 for lithium-ion battery cathodes Prashant Kumar Gupta a, Arihant Bhandari c, Jishnu Bhattacharya c, **, Raj Ganesh S Pala a, b, * a b c

Department of Chemical Engineering, Indian Institute of Technology, Kanpur, India Materials Science Programme, Indian Institute of Technology, Kanpur, India Department of Mechanical Engineering, Indian Institute of Technology, Kanpur, India

H I G H L I G H T S

G R A P H I C A L A B S T R A C T

� Demonstration of thermodynamics and kinetics interplay for LIB cathode. � Interfacial engineering of MnO2 poly­ morphs via thermal phase transitions. � Engineered intergrowth structure pro­ vide higher and flatter voltage profile. � “Native” structure scaffolds “NonNative” structure for better stability and Liþ mobility.

A R T I C L E I N F O

A B S T R A C T

Keywords: Lithium ion battery MnO2 Native Non-native Interface

Non-native structures have more open structure with less stable bulk formation energy compared to their native structure and hence, non-native structured cathodes are expected to generate higher discharge potential and better lithium mobility. However, non-native structures are limited in stability, which can be enhanced by scaffolding with native structure. We synthesize highly crystalline non-native Ramsdellite MnO2 (r/NN1–MnO2) and gradually thermal phase transform to native Pyrolusite MnO2 (β/N–MnO2), intimately interfaced (r/ NN1–MnO2)/(β/N–MnO2) intergrowth intermediate structures are realized. As the temperature increases, the ratio of r/NN1–MnO2 to β/N–MnO2 in intergrowth structure decreases, so does the discharge potential and lithium mobility. The wider discharge plateau for intergrowth structures is rationalized using density functional simulations followed by statistical averaging to account for interfacial intercalation sites with greater stability. Improved capacity retention for intergrowth structure is due to the higher structural stability and available free volume, where the presence of native structure acts as a stabilizing element. Electrochemical impedance spec­ troscopy correlates faster transport in the intergrowth structure with higher percentage of non-native phase with wider 2 � 1 channel compared to native phase with narrow 1 � 1 channel. The insights generated from this study have broad implications for interface engineering in multicomponent electrodes such as Li-rich-NMC oxides.

* Corresponding author. Department of Chemical Engineering, Indian Institute of Technology, Kanpur, India. ** Corresponding author. E-mail addresses: [email protected] (J. Bhattacharya), [email protected] (R. Ganesh S Pala). https://doi.org/10.1016/j.jpowsour.2019.227619 Received 28 August 2019; Received in revised form 28 November 2019; Accepted 12 December 2019 Available online 29 January 2020 0378-7753/© 2019 Elsevier B.V. All rights reserved.

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Journal of Power Sources 450 (2020) 227619

NN3-) nomenclature so that the relative stability of the phases is readily apparent. Traditionally, the utility of metastable structure in practical applications has been correlated to excess energy of crystallographic forms from its most stable form or native form [14,25]. Depending upon their stabilities of polymorphs of MnO2, we have classified them as native (N-) (i.e. the thermodynamically the lowest energy structure or the most stable) and non-native (NN-) (i.e. all other polymorphic structures other than the native structure) [14,26]. The differences in the energy of formation in different polymorphs arise due to differences in discrete translational symmetry in sub-surface (or bulk) region. Based on the energy of formation (most stable having most negative formation energy), the order of stability for different MnO2 can be written as β/N–MnO2 < γ/N–NN1–MnO2 < r/NN1–MnO2 < α/NN2–MnO2 < δ/NN3–MnO2 [25,27,28]. The γ-MnO2 is an intergrowth of native (β/N–MnO2) layers in non-native (r/NN1–MnO2) framework, which was pointed out by De Wolff in 1959, and is written alternatively as γ/N–NN1–MnO2 [29,30]. The theoretical capacity (which depends on the number of sites available for lithium binding in the crystal structure) of all the polymorphs of MnO2 is same (308 mAh/g considering stoi­ chiometric ion-transfer) as the basic unit for all polymorphic MnO2 polymorphs is the MnO6 polyhedra and one lithium ion binds to indi­ vidual MnO6 polyhedra. However, experimentally obtained capacity depends on the synthesis process as it affects the size, morphology, structural modifications, cations impurity, and cut-off potential [2, 31–33]. The non-native structure α/NN2–MnO2, with its large 2 � 2 channel, shows higher lithium storage capacity, however, the structure can easily transform during lithium insertion [34,35]. The native structure of β/N–MnO2 shows higher structural stability, but lower lithium storage capacity (0.2 Li/f.u.) [36]. Further, the excess lithium insertion results in high electrostatic interaction between lithium and oxygen in MnO2 structure which displaces half of Mn atoms into octa­ hedral sites, transforming the rutile structure into the spinel [Mn2]O4 framework of LiMn2O4 [27,28,37]. For the case of complex materials like Li(Ni,Mn,Co)O2 (NMC), it has been found that creating a composite layered-layered heterostructure with Li2MnO3 increases the structural stability and the capacity of the Li-rich NMC [38,39]. Likewise, for better overall performance of MnO2 structures, a composite intergrowth of native and non-native structure like γ/N–NN1–MnO2 is desirable as the native structure may provide the structural stability while non-native structure provides better kinetics and potential [2,30]. There can be different intergrowth structures of γ/N–NN1–MnO2 depending upon the percentage and arrangement of individual components r/NN1–MnO2 and β/N–MnO2. The exact structure of the intergrowth is hard to identify due to the presence of the De Wolff disorder and the Microtwinning in structure, as reported in detail by Chabre and Pan­ netier [30]. In our previous work, we have attributed the overall performance of the lithium-ion battery to a trade-off between the structural stability and the lithium storage capacity [11]. In order to rationalize the high per­ formance of γ/N–NN1–MnO2, we first explored the lithium intercalation mechanism in one of its component r/NN1–MnO2. Our results on r/NN1–MnO2 were combined with the findings of a study on β/N–MnO2 to provide a framework for rationalizing the performance parameters for other polymorphs of MnO2 [11]. γ/N–NN1–MnO2 has been synthesized by various techniques, which fall broadly in these three categories: electrolytic techniques (EMD), chemical method (CMD), and heat treated electrolytic techniques (HEMD), and have been explored as an electrode in lithium ion battery [2,40–42]. Each synthesis protocol shows a difference in the peak broadening and peak positioning in powder XRD pattern of γ/N–NN1–MnO2 which leads to different prop­ erties of the as-synthesized γ/N–NN1–MnO2 in the lithium-ion battery application [30,43,44]. The γ/N–NN1–MnO2 synthesized by EMD or CMD contains small amount of water at the grain boundaries as well as in the voids and hence, shows a lesser amount of lithium storage. In HEMD, heat treatment removes the water in the structure, which in­ creases the amount of lithium that can be stored. The heat treatment also

1. Introduction Design of materials is central for improving the energy density of energy storage devices. This is implemented by either a new material or by a material which is re-designed by modulating the crystal structure, changing the morphology or by doping with different elements [1–8]. Transition metal oxides are being extensively explored as positive electrode material in the lithium-ion batteries [9], exist in different polymorphic forms due to the variable coordination number, stability, bond angles and the bond distances between the transition metal and oxygen atoms [2,10–12]. Various polymorphs have different lithium-ion intercalation free energies and lithium storage capacity, both of which determine the energy density of an electrochemical energy storage de­ vice. More open structures have a larger number of lithium intercalation sites and may have a more negative lithium intercalation free energy and hence, greater discharge potential, but may have inferior structural stability upon repeated lithium intercalation/deintercalation. There­ fore, the optimal performance is a trade-off between discharge potential, lithium storage capacity and structural stability [11]. A systematic variation in the energy density may be possible when properties are correlated with the systematic variation in formation energy of different polymorphs. It is expected that the host polymorphic structure with a higher (i.e. less negative) formation energy will have a more negative free energy for lithium intercalation, resulting in a greater battery voltage [13]. To emphasize the thermodynamic un­ derpinnings of potential heuristics, we adopt the nomenclature of the “native” and “non-native” structure [14]. The native structure is the bulk structure with the highest thermodynamic stability and the non-native structure differ from the native structure in their discrete translational symmetry in the sub-surface (i.e. below the outermost surface layer) regions [14–16]. Typically for large crystallites, the non-native structures have less negative free energy of formation as compared to the native structure and hence, are less stable [14]. How­ ever, as the ordering of the free energy of formation may be dependent on the size and morphology of particles, and so we adopt a nomenclature based on the differences in discrete translation symmetry with respect to the native phase [14]. Generally, the non-native structures have higher free volume and lower surface energy [14]. For a given set of thermo­ dynamic conditions there will be only one native structure, but multiple non-native structures with different discrete translation symmetry are possible [14]. The stability of different non-native structures depends on the contribution from the bulk energy and the surface energy which are related to the stabilizing conditions like temperature, pressure, ligands etc. [14,17] Typically, more open non-native structures with lower co-ordination can accommodate lithium with lesser steric hindrance compared to the denser native structure and the more open structures may exhibit facile lithium-ion transport [18]. For example, TiO2, which exists in many polymorphic structures (Rutile, Anatase, Brookite, Bronze etc.) with relatively dense Rutile structure as the native structure, has been well explored for the lithium-ion battery anode [19]. The ther­ modynamically stable native structure, Rutile TiO2, can reversibly (i.e. without structure pulverization) accommodate only 0.1 mol of lithium per mole of the formula unit (f.u.), while the other non-native structures can accommodate higher amount of lithium: Anatase (0.5 Li/f.u.), Brookite (0.1 Li/f.u.) and Bronze (0.71 Li/f.u.) [19,20]. Calculations reported in the previous studies show that native TiO2 phase yields lower open circuit voltage than the non-native structures [21,22]. Similarly, in the case of Cobaltous Oxide (CoO), the native cubic phase forms lesser number of bonds with lithium and shows only 0.8 V as the discharge potential. However, its non-native hexagonal structure, which forms a greater number of bonds with lithium, exhibits a higher po­ tential of 1.2 V [23]. MnO2 exists in more than ten polymorphic forms and is a primary choice for battery industries [2,24]. In addition to the traditional poly­ morphic nomenclature, which generally does not provide indicator in the stability, we have appended native (N-) and non-native (NN1-, NN2-, 2

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Journal of Power Sources 450 (2020) 227619

increases the content of β/N–MnO2 in the γ/N–NN1–MnO2 depending upon the heat treatment temperature [41,42,45]. A. P. Malloy et al. reported that the heat treatment changes the surface properties of r/NN1–MnO2 and converts it completely into β/N–MnO2 at a tempera­ ture more than 400 � C [46]. Moreover, J. B. Arnott et al. and other groups have shown that during the heat treatment of r/NN1–MnO2, the material becomes denser as the temperature increases. The volume shrinking at higher temperatures leads to a kinetic limitation in con­ verting r/NN1–MnO2 to β/N–MnO2 [47,48]. Various independent ob­ servations on the polymorphs of MnO2, as mentioned above, indicate that there exists a correlation of the synthesis process to the structure as well as the electrochemical properties of γ/N–NN1–MnO2. In order to identify the correlation between the structure and elec­ trochemical property of γ/N–NN1–MnO2, we systematically change the structure from the non-native to native via gradual thermal treatment followed by the electrochemical characterization of the different sam­ ples through a coin cell assembly. We demonstrate the tuning of voltage and capacity in the lithium-ion battery with the modulated thermal transformation of the non-native (r/NN1–MnO2) to native (β/N–MnO2) structure. Firstly, we synthesize highly crystalline r/NN1–MnO2 fol­ lowed by the modulated thermal treatment to convert it to the different intergrowth structures corresponding to different ratios of non-native r/ NN1–MnO2 and native β/N–MnO2 structure. We use these samples in fabricating lithium-ion coin cell assembly to explore the effect of phase transformation on the device-level performance. We find that as the heat treatment temperature increases, the ratio of r/NN1–MnO2 to β/N–MnO2 decreases so does the discharge potential and lithium mobility. We rationalize the trend in the discharge potential from computed formation energies obtained from the density functional theory (DFT), followed by a statistical averaging to account for the different types of available sites for lithium intercalation in the inter­ growth structures. The galvanostatic discharge profile shows a wider discharge plateau for the intergrowth structures compared to the pure native and non-native structure. We rationalize the wider discharge plateau observed in the intergrowth structures on the basis of additional interfacial sites of higher stability than native and non-native structure. The high initial discharge capacity and its retention over the cycling for the intergrowth structures compared to pure native and non-native structures is due to the trade-off between available free volume and the thermodynamic structure stability in the presence of the native structure which acts as a scaffolding substrate inhibiting the structural transformation. The kinetic performance of different samples explored via the electrochemical impedance spectroscopy (EIS) suggest that a larger percentage of the non-native structure results in better kinetics. We rationalized the lithium diffusion in different samples based on our previous study by the lower energy barrier for the lithium-ion diffusion in non-native structure with 2 � 1 channel (200 meV) as compare to native structure with 1 � 1 channel (260 meV) [11]. Overall, we demonstrate a design principle potentially generalizable to other tran­ sition metal oxides wherein the four critical aspects of performance of cathode of Lithium-ion battery, namely, discharge potential, width of voltage plateau, higher cycle life and kinetics, can be improved without changes to the chemical composition, but by a careful manipulation of the microstructure of the material.

washed with DI water multiple times to remove unreacted acid. The synthesized particles have been dried at 100 � C overnight. For thermal transition, the as-synthesized r/NN1–MnO2 particles are taken in an alumina crucible and via a ramp rate of 2 � C/min heated up to different temperatures: 300, 320, 350, 400 and 450 � C followed by retaining the same temperature for 10 h. The heat treated samples are referred to as N-NN1_300, N-NN1_320, N-NN1_350, N-NN1_400 and β/N_450 respectively throughout the manuscript, where the number at the end refers to the temperature (in oC) at which it was heated. 2.2. Physico-chemical characterization techniques X-ray diffraction (XRD, PANalytical, Germany) has been performed for a range of 10� to 80� at a scanning speed of 2� per minute to determine the crystal structure. The surface elemental composition and chemical states of components have been analyzed by X-ray photo­ electron spectroscopy (XPS). XPS spectra (PHI 5000 Versa Prob II, FEI, USA) have been measured in the binding energy range of 0–1400 eV. The binding energy scale has been calibrated using the C 1s peak at 284.6 eV. Al Kα (25 W, 15 kV) is used as an emission source. Morpho­ logical characterization has been performed by Field emission scanning electron microscopy (FESEM, ZEISS Supra 40VP, Germany) operated at 10 kV voltage. Titan G2 12 60–300 KV TEM has been used to perform high-resolution transmission electron microscopy (HRTEM). We have used carbon coated copper TEM grids to perform HRTEM. HRTEM mi­ crographs were analyzed using FFT program in Gatan Digital Microg­ raphy software. Raman data of the samples were collected on JY-Horiba (Horiba iHR550) spectrometer using a laser of 532 nm. 2.3. Electrochemical measurements Firstly, the sample is ground in a mortar and pestle for 15 min prior to the fabrication of the electrode. 70 wt% of the active material, 15 wt% of super P as a conducting additive and 15 wt% of polyvinylidene fluoride (PVDF) binder are mixed in N-Methyl 2-Pyrrolidone (NMP) solution to make a slurry. The prepared slurry is coated on a carbon coated aluminium foil current collector with the help of doctor blade and dried at 120 � C for 12 h in vacuum. The electrodes are cut in the form of a circular disc of diameter 15 mm appropriate for 2032-coin cell. Lithium foils with 750 μm thickness is used as a counter electrode. The coin cells are assembled in the argon-filled glove box. 1.0 M LiPF6 in ethylene carbonate and diethyl carbonate EC: DEC (1:1 v/v) is used as an electrolyte solution. Trilayer polypropylene membrane (Celgard 2320) is used as a separator. Approximately 2–3 mg of active material is loaded on carbon coated aluminium foil. The cyclic voltammetry (CV) is per­ formed using an Autolab Potentiostat/Galvanostat (302 N, Netherlands) at a scan rate of 0.1 mV/s. Galvanostatic charge/discharge of assembled coin cells is measured by using a battery analyzer (MTI Corporation, USA). The EIS measurements are performed in potentiostat using NOVA software for a frequency range of 100 kHz to 0.1 Hz at DC potential with AC signal amplitude of 10 mV. 2.4. Computational methodology

2. Experimental and computational details

The reaction for the initial lithium intercalation into the MnO2 cathode for dilute Li concentration, x≪1, can be written as:

2.1. Material and synthesis

MnO2 þ xLi→Lix MnO2

(1)

The initial discharge potential of the battery during the above re­ action can be calculated as [11,49,50].

Lithium manganese oxide (LiMn2O4) has been purchased from Sigma-Aldrich, India. Sulphuric acid (H2SO4) has been purchased from Fisher Scientific. We synthesize the highly crystalline r/NN1–MnO2 by slow acid etching of LiMn2O4 as reported in our previous work [11]. The spinel LiMn2O4 particles have been dispersed in 2.6 M H2SO4 and stirred at 95 � C for 48 h. Next, the solution has been centrifuged at 8000 rpm and

Vavg ¼

ELix MnO2

EMnO2 x

x⋅ELi

(2)

Here terms in the numerator are ground state energies of products and reactants of Equation (1). For a single-phase compound of MnO2, e. g. the pure native (β/N_450) and non-native structure (r/NN1), there is a 3

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Journal of Power Sources 450 (2020) 227619

single type of lithium intercalation site in the structure, and the initial intercalation voltage can be directly calculated from Equation (2). For two-phase intergrowth structures of native and non-native phases (NNN1_300, N-NN1_320, N-NN1_350 and N-NN1_400), there are multiple sites for Li intercalation into the structure. In such a case, the average intercalation voltage is computed by taking a weighted average of the voltage produced by Li intercalation into each type of site giving weight to the probability of binding into each site: X Vavg ¼ Vi Ni Pi (3) i

Where, Vi is the intercalation voltage obtained during Li intercalation into the site of type i, calculated from Equation (2), Ni is the number of sites of type i and Pi is the probability of binding into this type of site, which is calculated as: Pi ¼ P

e

Eb i kT

j Nj e

(4)

Eb j kT

Where, Ebi is the binding energy of Li in the site of type i, k is the Boltzmann’s constant and T is the room temperature (298 K). All the energies are calculated using the Density Functional Theory (DFT) as implemented in the Vienna Ab-Initio Simulation Package (VASP) [51–53]. Projector augmented wave (PAW) pseudopotentials are employed for the electronic core states [54]. All calculations are spin-polarized. GGA þ U method is employed in the PBE scheme [55]. Ionic relaxations allowing cell shape and volume to change are per­ formed for effective U-J ¼ 3.9 eV [56]. An Energy cut-off value of 400 eV and an optimized Gamma-centered k-point mesh is used based on a convergence test. The structures of the two-phase compounds are con­ structed by making an interface between the non-native (r/NN1–MnO2) and native (β/N–MnO2) phases. The previous study has shown that due to the nature of the crystallographic structure of r/NN1–MnO2 and β/N–MnO2, the formation of a coherent interface between the two structures is possible only along the (010) direction [30]. To assess the formation of the (010) interface in the two-phase intergrowth structure, we calculate the interfacial energy by taking the difference of the energy of the two-phase intergrowth structure and that of the individual phases as [57]. Einterface ¼ Eγ=N

NN1

Eβ=N

Er=NN1

Fig. 1. XRD of the as synthesized r/NN1-MnO2 and the heat-treated samples at different temperatures.

respectively in the supporting information. To obtain the percentage of r/NN1 and β/N in the intergrowth samples, we perform Rietveld refinement. We refine the samples r/NN1, N-NN1_320 and β/N_450 with the crystal structure of pure Ramsdellite and Pyrolusite by the use of MAUD-Rietveld program as shown in Fig. S3 (in the supporting information) [58]. The as-synthesized r/NN1–MnO2 structure belongs to the Pnma space group (Fig. S3(a)) and contains 100% Ramsdellite while sample β/N_450 contains 100% Pyrolusite structure with a space group of P4_2/mnm (Fig. S3(c)). The intermediate structure N-NN1_320 contains both NN1–MnO2 and N–MnO2 with the weight percentage of 76% and 24% found from the refinement, respectively (Fig. S3(b)). The calculated Rietveld parameters Rwp are 7.1%, 4.9%, and 8.6% for r/NN1, N-NN1_320, and β/N_450 respec­ tively, which confirms a precise refinement. We compute the relaxed cell parameters of pure Ramsdellite and Pyrolusite MnO2 from density functional theory as shown in Fig. 2(a) and (b), respectively. The computed lattice parameters for Ramsdellite and Pyrolusite compare closely with the experimentally obtained lattice parameters from XRD for r/NN1 and β/N_450 as shown in Tables S1 and S2, respectively, in the supporting information. To rationalize the structure of the heat-treated samples, we superpose the XRD patterns of the Ramsdellite, Pyrolusite and an intergrowth structure of RamsdellitePyrolusite (computed using the PowderPlot feature in the VESTA soft­ ware) [59] on top of the experimentally obtained XRD patterns of r/NN1, β/N_450 and N-NN1_320 as shown in Fig. 2(a), (b) and (c), respectively. The close match in the peaks suggests that the intermediate heat-treated samples are actually, the intergrowth of the non-native (r/NN1–MnO2) and native (β/N–MnO2) phases. To further investigate the structural properties of different thermally treated MnO2 samples, Raman spectroscopy was used where low wavenumber reflects the nature of local structure rather than long-range structure. The MnO2 has been considered Raman inactive by B. Stroh­ meier et al. [60] and F. Kapteijn et al. [61] while C. Julien et al. [62] and F. Buciuman et al. [63] found it to be Raman active. We have taken the Raman data which shows a poor signal for different thermally treated samples (Fig. S4 in the supporting information). The spectra for the pure non-native (r/NN1–MnO2) and native structure (β/N_450) match well with the previously reported literature [62,63]. The intermediate sam­ ples shows continuous shift in the peaks from non-native structure to the native structure as the thermal treatment temperature increases. To assess the stability of the intergrowth structure between the nonnative (r/NN1–MnO2) and native (β/N–MnO2) structure, we calculate the interfacial energy. The interfacial energy is defined as the energy

(5)

3. Results and discussions 3.1. Physical characterization and structural details The as-synthesized r/NN1–MnO2 and the heat-treated samples are analyzed through XRD (shown in Fig. 1). The sharp peaks in Fig. 1 indicate that the as-synthesized non-native structure (r/NN1–MnO2) is highly crystalline and corresponds to the orthorhombic structure of Ramsdellite MnO2. For the samples corresponding to 300 � C, 320 � C, 350 � C, 400 � C and 450 � C heat treatment temperature, r/NN1 peak at 21.89� shifts to a higher 2θ angle and finally it merges with the Pyro­ lusite peak at 28.64� (β/N_450) (Fig. 1) confirming the completion of the phase transition of the non-native structure to the native structure. The other peaks also match well and continuously shift from the ortho­ rhombic structure of r/NN1 to the tetragonal structure of β/N_450. For the samples which are heat-treated at intermediate temperatures, NNN1_300, N-NN1_320, N-NN1_350 and N-NN1_400, the ratio of Rams­ dellite (NN1-) to Pyrolusite (N-) decreases. The increase in the broadness for some of the peaks for N-NN1_300, N-NN1_320 and N-NN1_350 is due to the De Wolff disorder and Microtwinning in the intergrowth structure of NN1/N–MnO2 which depends on the relative ordering of the struc­ tures in the material [30]. The experimentally obtained lattice param­ eters from XRD for r/NN1 and β/N_450 is shown in Tables S1 and S2, 4

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Journal of Power Sources 450 (2020) 227619

Fig. 2. (a) Experimental XRD of r/NN1–MnO2 and computed XRD of Ramsdellite MnO2. The structure and the computed lattice parameters from the DFT are shown on the right. (b) Experimental XRD of β/N_450 and computed XRD of Pyrolusite MnO2. The structure and the computed lattice parameters from the DFT are shown on the right. (c) Experimental XRD of N-NN1_320 and the computed XRD of two-phase intergrowth of Ramsdellite-Pyrolusite MnO2.

required to form the interface between two individual surfaces. We calculate the interfacial energy from Equation (5) as the difference in the energy of two-phase intergrowth structure (Fig. 3(a)) and the separated individual phases (Fig. 3(b)). The interfacial energy is found to be 293 meV/Å2. The negative value of the interfacial energy suggests a

favourable formation of the interface between the non-native (r/ NN1–MnO2) and the native (β/N–MnO2) phases. The morphology of different samples was analyzed by SEM images and shown in Fig. S5 in supporting information. All the samples have the morphology of nanorods with the length ~500 nm and width ~50 nm,

Fig. 3. The interfacial energy of the interface between Ramsdellite and Pyrolusite structures. 5

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Journal of Power Sources 450 (2020) 227619

and hence, the effect of morphology on the performance of different samples can be considered marginal. To analyze the intergrowth structures and the interface between nonnative (r/NN1–MnO2) and native (β/N–MnO2) structure, we have taken HRTEM images (Fig. 4). Fig. 4(a) shows the HRTEM image of as syn­ thesized non-native (r/NN1–MnO2) structure and the inset image shows the selected area electron diffraction (SAED) pattern. We observe well ordered fringes matching with the d-spacing of different planes calcu­ lated from the XRD spectra of non-native (r/NN1–MnO2) structure and the inset provides fast Fourier transform (FFT) image providing infor­ mation of different planes (Fig. 4(a)) (enlarged image is given in Fig. S6 in the supporting information). Similarly, the HRTEM image of pure native structure of MnO2 (β/N_450) is shown in Fig. 4(c) which suggest the presence of pure phase confirmed from the well ordered fringes matching well with the d-spacing of different planes calculated from the XRD spectra and the inset provides FFT image (enlarge image is given in Fig. S11 in the supporting information). We have also shown SAED pattern, which are well ordered bright spots for the same sample in inset image of Fig. 4(c). While the HRTEM image for the sample N-NN1_320 shows dominant non-native structure (r/NN1–MnO2) with small inter­ growth of β/N–MnO2 shown with the region R2 with yellow bordered line (enlarged image is given in Fig. S8 in the supporting information). We have characterized the different phases with the inset FFT image providing information of the planes (Fig. 4(b)). The SAED pattern (inset in Fig. 4(b)) also shows small interferences from both phases present in the intergrowth structure. We have also performed the High-angle annular dark-field (HAADF) image analysis with the elemental map­ ping for r/NN1, N-NN1_320, β/N_450-MnO2 samples and shown in Figs. S6(b) and S8(b) and S11(b) respectively in supporting information which confirms the presence on only Mn and O in the sample. Further, the detailed analysis of HRTEM for r/NN1-, N-NN1_320, β/N_450-MnO2 samples are given in Figs. S6(a) and S8(a) and S11(a) respectively in the supporting information. The HRTEM and the HAADF with elemental mapping analysis for all other intermediate samples NNN1_300, N-NN1_350, and N-NN1_400 are given in Fig. S7(a)&(b), S9 (a)&(b) and S10(a)&(b) respectively in the supporting information. The above mentioned detailed analysis of HRTEM shows that as the thermal treatment temperature increases, the region of native structure β/N–MnO2 (region R2 as mentioned in the different HRTEM images) in the sample increases, with both the parent structure native and nonnative-MnO2 retaining high crystallinity. For further characterization, we calculate the average oxidation state (AOS) of Mn from the spatial difference between the deconvoluted XPS Mn-3s peaks (ΔE3s) as shown in Fig. 5. The calculation of AOS from the XPS Mn-3s peak splitting is well established in the literature [64,65]. The value of AOS for all MnO2 samples calculated from the Mn-3s XPS spectra is given in Table 1, which shows that the oxidation state of Mn in MnO2 first increases with the increase in the heat treatment temperature for the samples N-NN1_300, N-NN1_320, and N-NN1_350 and then

Fig. 5. XPS spectra of Mn-3s orbital splitting energies of different MnO2 sam­ ples and the black line corresponds to experimentally obtained XPS spectra, the red and blue lines are split Mn-3s spectra, cyan line corresponds to the fitted spectra. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

decreases for the samples N-NN1_400 and β/N_450. The initial increase in the AOS can be attributed to the following two reasons: (a) heating r/NN1 sample up to 350 � C removes the available surface moisture as well as the water molecules present in the bulk of the sample as defects; (b) since the heating is done in the presence of air, the under-coordinated bonds get saturated during the thermal treatment, a Table 1 Calculated average oxidation state (AOS) for Mn atoms in different MnO2 structures from the XPS Mn-3s orbital splitting. Sample Name

Peak 1

Peak 2

ΔE3s

AOS ¼ 8.95–1.13 ΔE3s

r/NN1 N-NN1_300 N-NN1_320 N-NN1_350 N-NN1_400 β/N_450

84.40 84.14 84.39 84.43 84.50 84.43

88.99 88.66 88.81 88.82 88.93 89.06

4.59 4.52 4.42 4.39 4.43 4.63

3.76 3.84 3.95 3.98 3.94 3.71

Fig. 4. HRTEM image of (a) r/NN1–MnO2 (b) N-NN1_320 (c) β/N_450 with inset image of SAED pattern and FFT image providing information of different planes. 6

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Journal of Power Sources 450 (2020) 227619

phenomenon which is also indirectly observed through the increase in the weight percentage of the native structure with the temperature. The later decrease in the AOS for the samples N-NN1_400 and β/N_450 can be rationalized on the basis of thermal decomposition of a certain per­ centage of MnO2 to Mn2O3. It has been reported in the literature that the thermal decomposition of MnO2 starts slowly at a temperature above 400 � C and the average oxidation state start decreasing which finally converts to Mn2O3 completely at a temperature of 587 � C [66].

atom limit) using DFT as shown in Fig. S12 in the supporting informa­ tion. Then Equation (3) is used to predict the average intercalation voltage in the intergrowth structure based on the fraction of native and non-native structure found from Rietveld refinement. The average initial voltage during lithium intercalation in the intergrowth structure NNN1_320 calculated from Equation (3) is 3.32 V. Clearly, the computed values of the initial discharge voltages for the three structures (i.e. β/N, r/NN1 and N-NN1_320) show the same relative order as obtained experimentally (Fig. 6(a)). Based on the shift in the Bragg peaks in XRD of different samples, we can conclude that the percentage of the native structure increases as r/ NN1 < N-NN1_320 < N-NN1_350 < N-NN1_400 < β/N_450. The trend in the initial discharge potential is related to the thermodynamic free en­ ergy change and hence, the decrease in the initial discharge potential order is r/NN1 > N-NN1_300 > N-NN1_320 > N-NN1_350 > NNN1_400 > β/N_450 (shown in magnified form in Fig. S13 in supporting information for all samples). The overall capacity in all the samples is approximately the same (for a cutoff potential of 1.2 V) as observed in Fig. 6(a). It suggests that all the structures have the same total number of available sites for lithium intercalation. All samples show a difference in the capacity for higher cut-off potential 2.0 V as shown in Fig. 6(a) with a dotted line, for samples r/NN1-, N-NN1_320 and β/N_450 the capacity with a cut-off potential 2.0 V is 168, 238 and 200 mAh/g. However, the discharge plateau is broader for the intergrowth structures (up to x ¼ 0.67) compared to that for the pure native (up to x ¼ 0.5) or non-native structure (up to x ¼ 0.5). This suggests that the thermally treated intergrowth structures are having higher energy density compared to parent structures. From Fig. 6(a), we note that in the case of r/ NN1–MnO2, the voltage drop occurs at a lithium concentration of x ¼ 0.5 in LixMnO2. In our previous study, we rationalized the mechanism of

3.2. Characterization of discharge profile The galvanostatic discharge profiles for different samples are ob­ tained from the assembled coin cell, (the procedure is mentioned in the electrochemical characterization section 2.3). The first discharge profile is obtained at a rate of C/10 with a cutoff potential 1.2 V as shown in Fig. 6(a). The discharge profile for different MnO2 sample shows (Fig. 6 (a)) a single plateau followed by a sharp decrease in the voltage. We can clearly see that r/NN1 shows the highest initial discharge potential, β/N_450 shows the least initial discharge potential, while intergrowth structure show intermediate initial discharge potential. As the ratio of r/ NN1–MnO2 to β/N–MnO2 in the intergrowth structures increases, the initial discharge potential also increases. To rationalize the initial trend of the discharge potential among different samples, we compute the 0K energies of the non-native and the native structures of MnO2 and the Li0.05MnO2 (i.e. the dilute limit of the lithium intercalation) from DFT and find the initial output voltage from Equation (2). Initial discharge voltage is computed to be 3.43 V for Ramsdellite MnO2 (r/NN1) and 3.08 V for Pyrolusite MnO2 (β/N_450). To estimate the initial voltage in the intergrowth of r/NN1–MnO2 and β/N–MnO2, we calculate the relative stability of a lithium atom at different intercalation sites of the intergrowth structure (in the dilute Li

Fig. 6. (a) First lithium ion discharge profile for different MnO2 samples at a rate of C/10 with the cutoff potential of 1.2 V (b) The first cycle of cyclic voltammetry for different heat treated MnO2 at a scan rate of 0.1 mV/s (c) Structures of LixMnO2 beyond which voltage drop occurs for r/NN1–MnO2 and for (d) Intergrowth structure of γ/N–NN1–MnO2. Corresponding Li concentrations are (c) x ¼ 0.5 (2 Li per 4 MnO2) (d) x ¼ 0.67 (2 Li per 3 MnO2). 7

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Journal of Power Sources 450 (2020) 227619

lithium intercalation in r/NN1–MnO2 and the associated reason of voltage drop at x > 0.5,11. It was observed that lithiation occurs in the 2 � 1 channel of the r/NN1–MnO2 structure till all the channels are singly occupied, correspond to x ¼ 0.5 (shown schematically in Fig. 6(c)). Further lithium intercalation in the 2 � 1 channels (i.e. x > 0.5) results in repulsion between the immigrating lithium and the existing lithium in the same channel, leading to a shift in the lithium coordination from tetrahedral to octahedral. The shift in lithium coordination environment is accompanied by an anisotropic change in the lattice parameters and the Jahn-Teller distortion of MnO6 octahedra, accounting for a voltage drop beyond x ¼ 0.5 [11]. Similarly, a computational study by Wang et al. showed that the voltage drop in the case of Pyrolusite MnO2 occurs beyond x ¼ 0.5 [36]. The rationale provided in the aforementioned studies corroborates well with the experimental observation of voltage drop for x > 0.5 in case of r/NN1 and β/N_450 as shown in Fig. 6(a). We note that in Fig. 6(a), the discharge plateau for the intergrowth structures is wider and the voltage drop starts at x > 0.67. The extended voltage plateau in the case of intergrowth structures can be rationalized by the presence of additional interfacial sites of higher stability than NN1 and N_450 (see Fig. S12). To elaborate further, the repeating unit of the intergrowth structure is shown in Fig. 6(d). The initial lithium intercalation occurs in the 2 � 1 channel of the r/NN1–MnO2 unit which can be correlated with the higher free energy change during lithium intercalation and lower energy barrier (200 meV) for lithium diffusion in the wider 2 � 1 channel of r/NN1–MnO2 unit compared to higher diffusion barrier (260 meV) in the narrow 1 � 1 channel of β/N–MnO2 [11,36]. Further lithium intercalation into the same channel will cause high repulsion with the existing lithium. Hence, the next incoming lithium prefers to bind into the vacant 1 � 1 channel of the β/N–MnO2 unit. This occurs till all the sites in the β/N–MnO2 unit get filled, and further lithium intercalation can only occur in the 2 � 1 channel of the r/NN1–MnO2 leading to the voltage drop. In other words, voltage drop in the non-native r/NN1–MnO2 structure occurs only beyond the inter­ calation of 2 Li per 4 Mn in a primitive unit cell (Fig. 6(c)). When we extend this rationale to the intergrowth structure, the corresponding voltage drop occurs beyond the intercalation of 2 Li per 3 Mn in the primitive unit cell (Fig. 6(d)). The stoichiometric concentration of 2 Li per 3 Mn in a primitive unit cell corresponds to the 67% of Li-intercalation which is consistent with the experimental observation of voltage drop at x > 0.67. Thus the presence of additional intercalation sites in N-NN1 intergrowth structures delays the voltage drop as compared to the pure r/NN1–MnO2 structure [11], or β/N–MnO2 structure [36]. The difference in the discharge voltage plateau of different samples is

further correlated with the XPS derived change in the AOS in the following manner: (a) The initial increase in the AOS because of the loss in moisture followed by the saturation of the under-coordinated sites with oxygen, increases the number of equipotential sites in the inter­ growth structure (N-NN1_300, N-NN1_320, N-NN1_350), leading to a wider discharge plateau (as shown in Fig. 6(a)) (b) Further, heating reduces the AOS in samples N-NN1_400 and β/N_450, which decreases the charge storage capacity, leading to a narrower voltage plateau (as shown in Fig. 6(a)). For further characterization of the reaction between MnO2 host and lithium, cyclic voltammetry (CV) has been performed for a potential window of 1.2–4.2 V with a scan rate of 0.1 mV/s for all the samples (Fig. 6(b)). The CV plots for the samples show a single oxidationreduction peak corresponding to reaction Equation (1). Fig. 6(b) shows a reduction peak at 2.9 V for r/NN1 which shifts to the potential 2.5 V for β/N_450 and matches with the discharge voltage plateau in Fig. 6(a). The cycle life performance of r/NN1, N-NN1_300, N-NN1_320, NNN1_350, N-NN1_400 and β/N_450 samples are shown in Fig. 7(a) at a rate of 0.1C for a potential window of 2.2–4.0 V. From Fig. 7(a) we note, stable capacity retention in all the samples after initial few cycles. The samples r/NN1, N-NN1_320 and β/N_450 show the first discharge ca­ pacity of 160, 233, and 177 mAh/g, respectively. The capacity after 100 cycles for the samples r/NN1, N-NN1_320 and β/N_450 are 51, 149, and 75 mAh/g respectively. Fig. 7(a) shows an initial loss in capacity for all samples. The percentage loss in capacity during the first ten cycles for r/ NN1, N-NN1_300, N-NN1_320, N-NN1_350, N-NN1_400 and β/N_450 is 56.68%, 45.65%, 35.62%, 40.78%, 49.80% and 51.40% respectively. We note that intergrowth structure, N-NN1_320, shows a minimum ca­ pacity loss, while both the pure non-native structure (r/NN1) and native structure (β/N_450) show a much higher loss in capacity (56.68% and 51.40% respectively) in initial ten cycles. This trend of initial capacity loss can be rationalized by two counteracting effects: 1) The structure with higher free volume can easily undergo a structural transformation during electrochemical cycling as suggested in our previous study [11]. The order of free volume available in all the considered samples can be written as r/NN1 > N-NN1_300 > N-NN1_320 > N-NN1_350 >N-NN1_400 > β/N_450 based on the amount of native (1 � 1 channel) and non-native (2 � 1 channel) structure in different samples. 2) Due to the higher free volume available in the wider (2 � 1) channels of the non-native structure, we see that the compressive stresses developed during lithium intercalation are about 3.6 kbar less than that in the native structure. The lower stresses stabilize the non-native structures with respect to the native structures. The order of compressive stress can

Fig. 7. (a) Discharge capacity vs the number of cycles for heat treated MnO2 sample at 0.1C rate for a potential window of 2.2–4.0 V (b) Cycling performance of NNN1_320 at different discharge rate. 8

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Journal of Power Sources 450 (2020) 227619

be written as r/NN1 < N-NN1_300 < N-NN1_320 < N-NN1_350 < N-NN1_400 < β/N_450. Due to these counteracting effects, the capacity loss initially decreases and then increases, as we decrease the amount of non-native structure. Hence, an optimum structure is needed for better performance like N-NN1_320 in which a small amount of native struc­ ture provides the structural stability while the greater free volume of non-native structure provides the space for lithium intercalation. Further, we note that the capacities of all the intergrowth structures are higher than that of the pure native and non-native structures. We also observe that the sample N-NN1_320 shows the maximum capacity retention among all the considered samples even after 100 cycles. We have also shown the coulombic efficiency (CE) with the number of cycles for the sample of N-NN1_320 in the 2nd y-axis of Fig. 7(a). To assess the rate-capability, we examine the cycle life performance of the best per­ forming sample, N-NN1_320, at different discharge rate (0.1C, 0.5C, and 1C) and shown in Fig. 7(b). We note that the increase in cycling rate decreases the capacity and capacity retention due to kinetic inhibition of complete lithiation of the structure.

within the electrode. We show the variation in the fitted parameters (R2 and C2) for different heat-treated samples in Fig. 8(b). We observe that the drop in the R2 is marginal for samples N-NN1_300, N-NN1_320, and N-NN1_350 compared to r/NN1 and it increases sharply for samples NNN1_400 and β/N_450. We rationalize this change in the resistance R2 for different samples with the channel size available for the lithium intercalation. As the channel size decreases, the ion transport resistance increases. The small decrease in the R2 for the intergrowth samples NNN1_300, N-NN1_320, and N-NN1_350 is due to the presence of the stable native structure which inhibits any structural change and de­ creases the resistance for the lithium transport. While, for sample NNN1_400 and β/N_450 the percentage of the native structure is large, and native structure has smaller 1 � 1 channel size 1 � 1, smaller than the non-native with bigger channel size 2 � 1, hence higher resistance for the lithium diffusion. Moreover, we note that the pseudocapacitance C2 can be correlated with the amount of charge stored in different MnO2 samples (see the right-hand side vertical axis of Fig. 8(b)) and it shows high values for the intergrowth samples than the pure native or the nonnative structures. Based on the impedance analysis, we can conclude that there exists an optimum ratio of the native and the non-native phases which provides faster transport property and this optimal point is achieved for the heat treatment temperature of 320 � C. As we have seen for the MnO2, the intergrowth structures (N-NN1MnO2) have higher energy density and better cycle life, we believe that the insights obtained from the present study can be extended to known high-capacity cathode materials. The rationale for the improved per­ formance of the intergrowth structures is that the native component provides the structural stability, while the non-native component pro­ vides the higher voltage. Additionally, the intergrowth structures pro­ vide a wider discharge plateau due to the presence of additional thermodynamically stable interfacial sites. Similarly, for the case of Li (Ni,Mn,Co)O2 (NMC), it has been found that creating a composite het­ erostructure with Li2MnO3 increases the structural stability and the capacity of the Li-rich NMC [38,39]. Further, controlled modulation of the ratio of layered-NMC and Li2MnO3 can help in designing the cathode for optimum performance. The observations for these structures can be extended to design improved cathode materials for the next-generation batteries.

3.3. Measurement of lithium transport properties The power density of an electrochemical energy storage device de­ pends on the kinetics of lithium ion transport in the host crystal struc­ ture. The electrochemical impedance spectroscopy (EIS) has been used to obtain greater insight into lithium ion diffusion in the host crystal structure [67]. EIS analysis was performed after first discharge and NOVA 1.11 software has been used for obtaining the equivalent circuit. After first discharge, the cell has been kept for 6 h in an open circuit condition to reach the equilibrium. The measured EIS data at 1.2 V along with the equivalent circuit is shown in Fig. 8(a). For all the samples, we notice one small semicircle at the high-frequency region followed by a second semicircle for the medium frequency with a small tail at the low-frequency region. From experiments, we see that the first semicircle in the high-frequency regime corresponds to the electrode-electrolyte interface formation. We note that the diameter of this semicircle does not change much with the potential and it is almost constant for all heat treated MnO2 samples. In the equivalent circuit, the semicircle in the high-frequency zone is fitted with an R1C1 circuit, and Rs stands for the solution resistance (all EIS fitted parameters are given in Table S14 in supporting information). In the low-frequency regime, we observe major changes in the EIS spectra for different heat-treated samples. In the equivalent circuit, R2 represents the charge transfer resistance and the C2 can be correlated with the pseudocapacitance (which is charge stored in the electrode) while the Warburg element (W) represents the mass transfer effect

4. Conclusion In general, there are two common approaches towards material design: either explore a material with a new chemical composition or modify the material within a particular chemical composition in order to obtain better performance. Here, we demonstrate a specific design

Fig. 8. (a) EIS spectra of different MnO2 samples with their fitted equivalent circuit at 1.2 V (b) the variation in the value of the fitted parameter for different samples. 9

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Journal of Power Sources 450 (2020) 227619

principle that fits in the second category via structural alterations be­ tween two MnO2 polymorphs (r/NN1- and β/N–MnO2) by modulated thermal phase transition. We first synthesize the pure phase of nonnative structure (r/NN1–MnO2) and thermally transform it gradually to the native structure (Pyrolusite (β/N–MnO2)) to produce intimately interfaced (r/NN1–MnO2)/(β/N–MnO2) intergrowth structures. Such a procedure to generate intergrowth microstructures help in optimizing the structure for four important material properties: higher voltages, wider voltage plateau, better cycle life and kinetics. We observe that as the temperature for phase transition increases, the ratio of r/NN1–MnO2 to β/N–MnO2 decreases so does the discharge potential and lithium mobility. This is because the more “open” non-native structures will bind to Lithium with more negative free energy change and also provide facile lithium-ion diffusion pathways. We rationalize the trend in the discharge potential and the wider voltage plateau for intergrowth structures from density functional theory based computed formation energies followed by a statistical averaging to account for the different types of available sites for lithium intercalation in the intergrowth structures. Specifically, in the intergrowth structures, the grain bound­ aries between the native and non-native structures offer a certain per­ centage of binding sites to improve the voltage plateau observed during discharge. The optimized intergrowth structure corresponds to the sample heated at 320 � C, which has a composition of 76:24 of r/NN1:β/ N–MnO2 (wt:wt) respectively. In the intergrowth structure, the native structure acts as scaffolding support due to the higher thermodynamic stability and inhibits the structural transformation and hence shows improved cycle life. The cycle life of the best performing sample, NNN1_320, is analyzed at different rates and it shows capacity retention of ~64%, 50% and 38% at 0.1C, 0.5C, and 1C rate, respectively, after 100 cycles. The lithium transport properties are also rationalized with impedance spectroscopy and we find that the larger percentage of native phase inhibits the lithium transport in the intergrowth structure due to the smaller channel of size 1 � 1, compared to the non-native structure with 2 � 1 sized channel. In this work, we demonstrate that systematic variation and improved electrochemical performance for the cathode material with greater voltage, a broader voltage plateau, better cycle life, and faster kinetics, can be achieved through the modulation of the polymorphic phase-composition without changing or introducing a new chemical composition.

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Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments Authors are grateful to the Indian Space Research Organization and the Science and Engineering Research Board, Department of Science and Technology for supporting this work via grants STC/CHE/2014098 and SERB/F/11147/2017–2018 respectively. We also thank the Center for Nanoscience at IIT Kanpur for providing the experimental facilities. We thank Prof. Rajeev Gupta from the Department of Physics, IIT Kanpur for extending the Raman facility for material characterization. RGP thanks Prof. K. S. Gandhi, Indian Institute of Science for critically reviewing the manuscript and offering valuable suggestions. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227619. References [1] J.B. Goodenough, Design considerations, Solid State Ion. 69 (3–4) (1994) 184–198.

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