All-solid-state sodium-ion batteries operating at room temperature based on NASICON-type NaTi2(PO4)3 cathode and ceramic NASICON solid electrolyte: A complete in situ synchrotron X-ray study

Chemical Engineering Journal 472 (2023) 144509

Available online 29 June 2023

1385-8947/© 2023 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by- nc-nd/4.0/).

All-solid-state sodium-ion batteries operating at room temperature based on NASICON-type NaTi ₂ (PO ₄ ) ₃ cathode and ceramic NASICON solid electrolyte: A complete in situ synchrotron X-ray study

Bidhan Pandit

^a,*

, Morten Johansen

^b

, Bettina P. Andersen

^b

, Cynthia S. Martínez-Cisneros

^a

, Belen Levenfeld

^a

, Dorthe B. Ravnsbæk

^b

, Alejandro Varez

^a,*

aDepartment of Materials Science and Engineering and Chemical Engineering, Universidad Carlos III de Madrid, Avenida de la Universidad 30, 28911 Legan´es, Madrid, Spain

bCenter for Integrated Materials Research, Department of Chemistry, Aarhus University, Langelandsgade 140, DK-8000 Aarhus, Denmark

A R T I C L E I N F O Keywords:

NaTi2(PO4)3

NASICON Cathode material Sodium-ion battery All-solid-state battery

A B S T R A C T

All-solid-state sodium-ion batteries that work at ambient temperature are a potential approach for large-scale energy storage systems. Nowadays, ceramic solid electrolytes are gaining attention because of their good ionic conductivity and excellent mechanical and chemical stabilities. Furthermore, a good interface between electrode and solid electrolyte is also required to achieve successful cell performances. In this work, sintered ceramic layer electrolyte Na3.16Zr1.84Y0.16Si2PO12, with high ionic conductivity (0.202 mS/cm at room temperature), are prepared by using uniaxial pressing followed by a sintering process. The conductive carbon coated NASICON material (NaTi2(PO4)3/C) exhibits, as cathode material, enhanced rate capability and stability for sodium ion batteries for high carbon (18.95 %) coated sample. At C/10, the optimized cathode (with higher carbon content) achieves a remarkable initial discharge capacity of 107.3 mAh/g (reversible capacity of 101.4 mAh/g), a suf- ficient rate capability up to a rate of 10C, and a long cycle life (capacity retention of 58% after 950 cycles). The one-stage reversible biphasic reaction mechanism and potential-dependent structure–property of NaTi2(PO4)3 can be explained by employing in situ X-ray synchrotron method. Sequential Rietveld refinements of the in situ data show the evolution of the Na-poor NaTi2(PO4)3 and Na-rich Na3Ti2(PO4)3 phase fractions (wt%), unit cell characteristics, and unit cell volume. The design of an all-solid-state sodium ion half-cell with a NaTi2(PO4)3/C cathode and a Na3.16Zr1.84Y0.16Si2PO12 solid-state electrolyte interface results in stable capacity of 83.6 mAh/g at C/10 and excellent reversible capacity at high C-rate. The results show that sintered NASICON-based electrolytes can significantly contribute for the fabrication of all-solid-state sodium-ion battery due to the superior con- ductivity and stability.

1. Introduction

Since the start of the 21st century, modern technologies reliant on fossil fuels have corresponded a host of serious environmental chal- lenges. Green energy generation, such as solar and tide energy, has arisen as a hot research topic for long-term and sustainable development of energy [1,2]. These renewable energy sources, on the other hand, are distinguished by their location, performance and stability [3]. Second- ary battery system for electrochemical energy storage is one of the most effective methods of converting and storing energy [4]. Lithium-ion batteries are widely used in our everyday life due to their high output

voltage and energy density, such as in portable devices and electric automobiles [5,6]. Li-ion batteries, on the other hand, associate with flammable, hazardous chemical solvents in electrolytes and face a growing scarcity of lithium supplies [7]. As a result, the safety and economic issues arise in the way of large-scale energy storage applica- tions [8,9]. Recently, rechargeable sodium-ion batteries (SIBs) have successfully proved their merits and attracted the attention of re- searchers [10,11]. Because of its abundance in the Earth’s crust, sodium is substantially cheaper than lithium [12–15].

Traditional organic electrolytes are widely used for sodium ion batteries due to their higher ionic conductivity [16]. Sodium-ion

* Corresponding authors.

E-mail addresses: [email protected] (B. Pandit), [email protected] (A. Varez).

Contents lists available at ScienceDirect

Chemical Engineering Journal

journal homepage: www.elsevier.com/locate/cej

https://doi.org/10.1016/j.cej.2023.144509

Received 31 March 2023; Received in revised form 15 June 2023; Accepted 28 June 2023

batteries with liquid electrolytes, on the other hand, may pose safety risks and have a low energy density. Unfortunately, when using a standard liquid electrolyte, Na-ion batteries, like other alkaline metal batteries, have side reactions, dendrite development, and safety issues [17]. Although liquid electrolytes are widely used owing to their low cost, high ionic conductivity and ease of use, solid electrolytes are often preferred since they are nontoxic, non-flammable, non-corrosive, non- volatile and thermally stable across a wide temperature range [18,19].

All-solid-state sodium batteries (ASSBs) based on solid electrolytes outperform orthodox SIBs in terms of safety, packing efficiency and energy density [20,21]. Because of air stability, good ionic conductivity, high ionic transference number and large voltage window, NASICON solid electrolytes are promising alternative to many other solid-state electrolytes [22,23].

NASICON-type materials, of general formula A_xM₂(XO₄)₃, exhibit very interesting properties, such as high thermal and chemical stability, high ionic conductivity and low thermal expansion [24]. Depending on the cation A and transition metals M, NASICON-type materials can serve as cathode, anode and electrolyte in Li-ion or Na-ion batteries. Thus, if M =V or Ti, with ability to oxidize or reduce (e.g. V³⁺/V⁴⁺or Ti³⁺/Ti⁴⁺) can be used as cathode or anode material in Li-ion or Na-ion batteries [7,25]. Meanwhile, Na1+xZr2SixP3−xO12(0 ≤x ≤3), NZSP where M =Zr which is difficult to reduce from its stable oxidation state +4, the ma- terial is a good electrolyte for solid-state Na-ion batteries.

NASICON-type NaTi2(PO4)3 (NTP) has gained significant interest as potential cathode for SIBs due to its outstanding low cost, safety, good conductivity, and open 3D structure with a wide interstitial site [26,27].

The NASICON-type electrode material is associated with an open 3D framework with two TiO6 octahedra isolated by three PO4 tetrahedra sharing all corner oxygens without sharing the edges or faces [28,29]. In addition, when compared to Ti⁴⁺/Ti³⁺redox couple, NTP has a great theoretical capacity of 133 mAh/g at about 2.1 V (vs. Na/Na⁺) [30].

Despite these advantages, pure NTP has a low electrical conductivity, which limits its usage as cathodic material [31]. It has been investigated to improve the electrochemical performance of NASICON materials using two widely used methods: altering the particle size of electrode materials to reduce both electronic and ionic migration channels, and enhancing the electronic conductivity [32]. Several efforts have been attempted to advance the electrical conductivity of NTP by carbon coating [33], metal doping [3], and nanoarchitecture fabrication [34].

In recent years, many processes, including sol–gel [35], hydrothermal [36], spray-drying [37], and electrospinning [38] have been used to prepare the NTP. Because of the ease of scale-up and the simplicity, the one-step solid-state process is a popular method for producing graphene or carbon decorated NTP.

On the other hand, NZSP structured solid-state electrolyte is also 3D framework made of connected polyhedral, forming a skeleton with many interstitial sites suited for the migration of monovalent cations [39]. The structure exhibits great chemical and thermal stability, that is important in ensuring the battery’s dependable performance. NZSP electrolyte has an ionic conductivity in the order of 10⁻² S/cm at working temperatures ranging from 50 to 100 ^◦C [40]. As a result, combining these approaches to build a solid-state battery with NASICON-based cathode and electrolyte will be an attractive feature.

To improve the cyclability and rate capability, NTP with controlled particle size, optimum porous structure and carbon coating are required.

We established a one-step solid-state approach to synthesis NTP with carbon coating layers (NTP/ C) using sugar as the carbon source. X-ray diffraction (XRD), Raman spectroscopy, field emission-scanning elec- tron microscopy (FESEM), high resolution transmission electron mi- croscopy (HRTEM), and X-ray photoelectron spectroscopy (XPS) were used to examine the materials. The highly crystalline NTP nanoparticles coated with carbon exfoliated layer are expected to advance the elec- trical conductivity of the material and electrolyte/electrode interface contact. The NTP/C electrode material combines the benefits of the carbon network and the NASICON structure to advance the

electrochemical performance of the NTP cathode. Furthermore, the carbon network can function as a continuous electron transmission route [41,42]. The associated porous structure of the NTP/C allows the elec- trolyte to penetrate into the electroactive sites [43]. Two different, high and low, coated NTP/C composite has been prepared. The high coated NTP/C electrode outperforms the lower carbon sample in terms of rate performance. The NTP/C sample with high wt% of carbon coating has a first discharge capacity of 107.3 mAh/g at C/10 and a reversible charge capacity of 101.4 mAh/g with excellent stability up to 950 cycles. The low total (grain boundary + bulk) conductivity at RT of sintered Na3.16Zr1.84Y0.16Si2PO12 is probably key in order to be used as electro- lyte in ASSB operating at RT. We are not aware of any studies that focus on solid-state batteries that use NASICON-based carbon layer coated NTP with sintered NASICON-based Na3.16Zr1.84Y0.16Si2PO12 electrolyte.

The developed solid-state battery has a strong reversible capacity (about 83.6 mAh/g at C/10) and high-rate capabilities (up to C/5), with cycle stability equivalent to liquid electrolyte-based sodium half cells. This all- solid-state architecture will provide more potential electrodes for high- performance rechargeable room temperature Na-ion batteries.

2. Experimental

2.1. Synthesis and preparation of electrodes and electrolytes

5 mmol Na₂CO₃ (Sigma-Aldrich, 99.9%), 20 mmol anatase TiO₂ (Sigma-Aldrich, 99.8%), 10 mmol NH4H2PO4 (Sigma-Aldrich, 99.9%), and equivalent quantities of sugar (C12H22O11, Merck, >99%) such as 1 g and 2 g were added altogether as precursors in a stainless steel vessel.

The precursors were ball-milled for 8 h in a planetary ball mill (PM 200) at 400 rpm. Following ball milling, the material was compressed into pellets and placed in oven at 800 ^◦C in an Ar atmosphere for 12 h. The associated materials were labeled NTPC1 and NTPC2, as per the low and high carbon content, individually.

The solid-state electrolyte Na3.16Zr1.84Y0.16Si2PO12 was synthesized using a typical solid-state reaction technique [44]. This composition is selected as it is demonstrated that the doping of rare-earth element in NASICON increases its density and reduce the grain-boundary resistance [45]. The following reagents were used in stoichiometric quantities as precursors: Na2CO3 (Sigma-Aldrich, 99.9%), NH4H2PO4 (Sigma-Aldrich, 99.9%), SiO2 (Merck), and Yttria-stabilized zirconia powder (8 mol % YSZ from Tosoh). In summary, powders were ball-milled in wet (ethanol) in zirconia jars for 24 h at 350 rpm with zirconia balls (diameter 5 mm). Ethanol is removed from the mixture by heating at 80

◦C overnight. Afterwards, the material is subjected to two consecutive thermal treatments in air atmosphere: i) preheating at 500 ^◦C for 4 h, followed by 800 ^◦C for 4 h; ii) calcination at 1100 ^◦C for 4 h.

NASICON pellets (13 mm in diameter) are obtained by uniaxial pressing applying 220 MPa during 5 min. Pellets are subsequently sin- tered in atmospheric air for 10 h at 1200 ^◦C. To prevent sodium and phosphorus evaporation caused by the high sintering temperature, pellets are coated with the NASICON powders during sintering.

2.2. Material characterization

Cu-Kα radiation (λ =l.5406 nm) was used in the Bruker AXS D8 Advance for the XRD investigation with a diffraction angle (2θ) range of 10–80^◦. XPS (PHI 5000 VersaProbe II ULVAC INC) was used to measure the composition, structure, and crystallographic information of elec- trode materials. In a JEOL Model JSM − 6390LV, the surface morphology and related composition were examined using FESEM and energy dispersive X-ray (EDX) spectroscopy. A LaB6 source and model JEOL 2100 were used for HRTEM analysis.

2.3. Electrochemical characterization

The electrochemical performance of NaTi2(PO4)3/C as the cathode

material was examined using CR2032 coin-type half batteries. The active material, conductive additive (9 % C65 +9 % vapor grown car- bon fibers), and polyvinylidene fluoride (PVdF) binder were combined in a mass ratio of 70:18:12 using N-methyl-2-pyrrolidone (NMP) as the solvent to produce a viscous slurry (300 μL − 1 ml). The resulting slurry was evenly coated as a current collector on clean aluminum foil and dried for 12 h in vacuum at 80 ^◦C. The active material mass loading was 2–5 mg. Sodium was used for both the reference and counter electrodes.

The separator and electrolyte were used of glass fiber GF/D (Whatman) and a 1.0 M NaClO4 solution in propylene carbonate (PC), respectively.

The CR2032 coin-type half-cells were constructed in a Jacomex glove- box filled with pure argon, with O2 and H2O levels regulated to less than 1 ppm. The liquid electrolyte cells were prepared at pressure of 800 psi.

To ensure complete electrolyte penetration, the cells were kept at rest for 12 h at 25 ^◦C before conducting electrochemical tests. In case of solid- state battery, the Na3.16Zr1.84Y0.16Si2PO12 solid-state electrolyte was sandwiched in between NaTi2(PO4)3/C cathode and sodium metal.

The Galvanostatic charge/discharge profiles, rate capabilities and cycle performances were investigated at various current densities in the potential range of 1.5–2.8 V. The specific capacities were estimated using the active material mass loading on the working electrode. All electrochemical experiments were performed at a temperature of 25 ^◦C.

The ionic conductivity of the solid-state electrolyte is evaluated by applying electrochemical impedance spectroscopy (EIS) as a function of temperature (− 30 ^◦C to 90 ^◦C), while heating, using an Impedance/

Gain-Phase Analyzer SI1260 (Solartron) coupled to the Interface 1287.

For this purpose, both electrolyte faces are previously covered with gold ink, acting as blocking electrodes, and a sinusoidal signal (100 mV) is applied in the frequency range from 0.1 Hz to 1 MHz.

2.4. Operando synchrotron powder X-ray diffraction and sequential Rietveld refinement

For the operando synchrotron powder X-ray diffraction (SR-PXRD), cathode pellets were prepared by mixing active material with carbon black (conductive additive) and polytetrafluoroethylene (PTFE) binder with the mass ratio of 80:15:5. An AMPIX battery cell developed for transmission X-ray scattering studies was mounted with a cathode pellet sample [46]. The microporous glass fiber (Whatman GF/B) separator, the Na-metal disc anode and 1 M NaClO4 in PC liquid electrolyte were used to fabricate the cell which was mounted on the diffractometer at beamline DanMAX, MAX IV, Lund, Sweden [47]. A BioLogic VMP3 potentiostat was used to analyze the electrochemical performance of the AMPIX battery cell, which was cycled between 1.5 and 2.8 V at a C/5 rate. It should be noted that the calculated capacity is only based on the weight percentage (wt %) of NaTi2(PO4)3in the pellet, i.e.,xin Nax-

Ti₂(PO₄)₃is associated only with insertion/extraction in NaTi₂(PO₄)₃. PXRD data were collected during charge–discharge on a DECTRIS PILATUS3 X 2 M CdTe area detector. Every two minutes, PXRD patterns were recorded using an X-ray exposure duration of 2s with the wave- length of λ =0.354130 Å. For easy comparison with the angular Bragg reflections, the operando diffraction data is shown in Q-space (Q =4πsin (θ)/λ). Each diffractogram was normalized and a background was sub- tracted by collecting scattering data from an AMPIX battery test cell engaging just electrolyte and separator.

Rietveld refinements of all the collectedoperandoPXRD data were carried out in a sequential manner using the Fullprof software [48]. In all refinements, the diffraction peaks were defined using a Pseudo-Voigt profile function and the background was defined by linear interpolation of manually selected refined points. The primary PXRD pattern was refined by using the structural models of NaTi2(PO4)3(space group: R- 3c) [49] and Na3V2(PO4)3 (space group: R-3c, exchanging V for Ti) [50]

in order to get an initial well-described reference point. For both phases, these refinements were applied sequentially from scan 10 to scan 222.

For the sequential refinement scale factors, lattice parameters, atomic positions, overall vibration factors and profile parameters were refined.

3. Results and discussion

3.1. Structural and morphology characterizations of NaTi2(PO4)3

Fig. 1a depicts the XRD patterns of the as-prepared NaTi₂(PO₄)₃/C.

All of the diffraction peaks are unambiguously ascribed to the well- crystallized NASICON-type NaTi2(PO4)3 with the rhombohedral space group R-3c (JCPDS number 85–2265), confirming previous findings [51,52]. Furthermore, no peaks associated with the carbon material are visible in the XRD profile, indicating the amorphous character of the coated carbon layer. The findings imply that the solid-state approach might be used to create a contaminant-free NaTi2(PO4)3/C material.

Fig. 1b depicts the associated crystal structure of prepared NaTi2(PO4)3. TiO6 octahedra and PO4 tetrahedra are combined in an open 3D framework to develop NaTi₂(PO₄)₃. All the PO₄ tetrahedron is linked to four TiO6 octahedra, which in turn are linked to six PO4

tetrahedra. By exchanging apical oxygen, they were able to generate a robust three-dimensional structure [53]. XPS and TG analysis is briefly discussed in supporting information (Fig. S1).

Fig. 2a, b show SEM images of the NTPC1 and NTPC2 powders. The particle sizes in both samples ranged from hundreds of nanometers to several micrometers. The particle size of NTPC2 is somewhat larger than that of pure NTPC1, possibly due to the thick carbon coating generated on the NaTi2(PO4)3 particle during the high-temperature calcination process. The elemental composition of NTPC2 was evaluated using EDS spectrum (Fig. S2 and Table T1), which revealed that the elemental content has an approximate element ratio of Na:Ti:P:O = 1:2:3:12, compatible with the development of NaTi2(PO4)3 composition. The EDS elemental mapping images indicate uniform distributions of Na, Ti, P, and O from the NTP, as well as C from the carbon coating layer (Fig. 2c, d). The composite’s homogenous distribution of C may help in electron transport and Na⁺diffusion kinetics.

HRTEM were utilized to further analyze the in-depth structure of the samples. As demonstrated in Fig. 2e, a thin carbon coating shell (15–30 nm) is efficiently coated on the surface of the NaTi2(PO4)3 particle core.

This carbon layer has the potential to significantly increase active ma- terial conductivity [54]. Furthermore, the HRTEM image clearly reveals the lattice fringes of NaTi2(PO4)3, indicating that the particles are well- crystallized.

3.2. Battery performance of the NTPC electrodes in liquid electrolyte The Na-storage capacities of NaTi2(PO4)3 were examined using gal- vanostatic tests throughout a voltage range of 1.5 to 2.8 V in half-cell vs.

Na in 1 M NaClO4/PC electrolyte at C/10 (1C =63 mA/g, depending on the active material), and the results are shown in Fig. 3. The NaTi2(PO4)3

cathode has an open-circuit voltage (OCV) of 2.6 V (against Na/Na⁺).

The first discharge and charge capabilities indicate that 1.4 Na⁺may be introduced in first cycle and subsequently 1.3 Na⁺can be removed in next cycle (Fig. 3a). The NTPC2 electrode, on the other hand, can able to let insert 1.7 Na⁺and extract 1.6 Na⁺on subsequent cycle as depicted in Fig. 3b. Each sample’s Ti⁴⁺/Ti³⁺redox couple peaks have been identi- fied. It corresponds to the observed phase transition throughout Na⁺ insertion/extraction. NTPC2 has more well-defined redox peaks and a less potential difference among the cathodic and anodic peaks than NTPC1, indicating less electrochemical polarization and greater reversibility of the electrode. The NTPC1 has an initial discharge ca- pacity of 88.7 mAh/g, approaching towards theoretical capacity. The NTPC2 exhibits a larger discharge capacity of 107.3 mAh/g and a better initial Coulombic efficiency (95 %) than the NTPC1, indicating efficient use of NTP particle with a highly reversible sodiation/desodiation process.

The first three differential capacity (dQ/dV) curves derived from the obtained galvanostatic profiles to understand better the char- ge–discharge processes. The NTP1 and NTPC2 cells’ dQ/dV curves in Fig. 3c, d exhibit a sharp discharge peak at 2.11 V and a sharp charge

Fig. 1. (a) XRD of NTPC1 and NTPC2 powder samples. (b) Crystal structure of NaTi2(PO4)3.

Fig. 2.(a, b) FESEM images of NTPC1 and NTPC2 samples. (c, d) EDS elemental mapping of white square associated with NTPC2 sample. (e) Associated TEM image.

peak at 2.15 V, corresponding to flat charge/discharge plateaus. Over the whole 1.5–2.8 V voltage range, no additional peaks of side reactions are found, and the sharp charge/discharge peaks are maintained without any changes in position with the cycling, suggesting a very stable and reversible redox activities [55]. However, after the first cycle, the capacity decreases, which might be due to the dissolution of

electrode material as no voltage change is seen when cycling. After the first two cycles, the coulombic efficiency reaches 100% in both situa- tions. When compared to the first cycle, the discharge capabilities of both samples fall a little during the second cycle. The formation of solid electrolyte interphase layer (SEI) might cause mostly the permanent capacity loss [56,57].

Fig. 3. Electrochemical charge/discharge curves per Na⁺ion and derivative of Galvanostatic charge/discharge of first three cycles for (a and c) NTPC1, (b and d) NTPC2 samples in half sodium cells in 1 M NaClO4/PC electrolyte at C/10.

Fig. 4.The rate capability (between C/10 and 10C) and Galvanostatic charge/discharge profiles for (a and b) NTPC1, (c and d) NTPC2 samples in half-cell vs. Na in 1 M NaClO4/PC electrolyte.

The discharge–charge patterns of the electrodes at various C-rates are shown in Fig. 4. When the current density increases from C/10 to 10C, the discharge capacities of the electrodes show moderate capabil- ities even at high rates, and reversibly recover to the starting capacity values when the current rate returns to low. The capacity drops of the NTPC2 electrode (Fig. 4c, d) towards the high current rate 10C is less than that of the NTPC1 electrode (Fig. 4a, b), implying that the NTPC2 electrode has superior rate performance. Notably, the capacity fading of the electrodes are less at high C-rates, which may be attributed to the high carbon content. Carbon may be capable of enhancing electron transfer and improving accessibility between NTP particles and elec- trolyte due to its high number of fast transport channels [58]. Na⁺can occupy two separate sites in the NASICON NTP structure and rapidly transport among one another due to local heating produced by higher

charge/discharge rates. This phenomenon has been seen in other NASICON-type materials, and it is considered to be produced by struc- tural rearrangement caused by temperature change, resulting in elec- trochemical performance instability. When the current rate changes for NTPC electrode, the first cycle becomes intermediate cycle, that’s why it provides incomplete sodiation. But after 1^st cycle, it gets stable to the next cycle and gives stable Columbic efficiency for the remaining cycles ate the same current rate. This phenomenon is normal during rate capability studies reported in literatures [59–61].

The cycling performances of both samples at relatively high C/5 rate are depicted in Fig. S3. The NTPC1 electrode has a reversible charge capacity of 46.5 mAh/g and low capacity retention (17.3%) after 950 cycles, which is due to the thin carbon shell’s less intrinsic electric conductivity (Fig. S3a). Except for the primary cycle, the reversible

Fig. 5.(a) Contour plot of the collected scattering data from in situ X-ray synchrotron studies with a Scan# from 10 to 222. (b) Collected galvanostatic charge/

discharge data. (c) Extracted wt %. (d) Lattice parameters a and c. (e) Cell volumes (ų) extracted from the sequential Rietveld refinement.

capacity of the NTPC2 electrode (68.6 mAh/g) is clearly larger than that of the NTPC1 (Fig. S3b). Moreover, it shows high capacity retention 58%

after 950 cycles. Furthermore, after a few cycles, the discharge capacity of the NTPC1 electrode rapidly falls when compared to the NTPC2 electrode, despite the fact that the Coulombic efficiency remains close to 100% in both situations. Fig. S4 depicts charge–discharge curves for NTPC1 and NTPC2 after various cycles. After the same cycles, the discharge plateaus of NTPC2 are better maintained than those of NTPC1.

This indicates how a carbon coating may increase NTPC2

′

s cycle stability significantly. The high conductivity of coated carbon, which speeds up electron transmission and prevents electrons from partially accumula- tion on the electrode’s surface, may be responsible for NTPC2

′

s improved cycle stability. This can help to avoid electrode degradation and maintain structural integrity, leading in higher cycle performance.

3.3. Dynamic charge–discharge phase evolution

During the discharge and charge processes of a sodium battery, the sodiation and desodiation of the cathode, NTPC, occurs with the following phase transformation:

NaTi2(PO4)₃+2Na⁺+2 e⁻←Sodiation/expansion

Desodiation/shrinkage→Na3Ti2(PO4)₃ For deeper understanding of structural phase transformation and the volume change of the cathode during discharge–charge processes, PXRD patterns (Fig. 5a) were collected from the positive electrode surface during the cycling of a battery, Galvanostatic charged/discharged be- tween 1.5 V and 2.8 V at C/5 rate (Fig. 5b). In Fig. 5, the collected operando PXRD is directly associated to concurrently recorded galva- nostatic charge/discharge profile.

Rietveld refinement was performed using NASICON-type NaTi2(PO4)3 with R-3c space group as the structural model [49]. The refined lattice parameters of the NASICON-type phase were a =8.486(2) Å, c =2.1.800(8) Å, and V =1359.3(7) ų, which corresponded rather well with previously published references [62]. As it has already been shown above, sodium insertion and extraction in NaTi2(PO4)3 proceeds via two clear flat redox plateaus at approximately 2.10 and 2.15 V (Fig. 5a). After the very beginning of the Na⁺intercalation, PXPD pat- terns reflect as expected a first order phase transition, i.e. the peaks from the initial NaTi2(PO4)3 phase gradually disappear and new peaks grad- ually appear at lower 2θvalues during sodium insertion. The new peaks originate from the isostructural sodiated Na3Ti2(PO4)3. This behavior is reversed upon sodium de-insertion, i.e. NaTi2(PO4)3 reforms. No other phases are observed in the SR-PXRD data. Nor are significant changes in the in the angular peak positions of, which would signify Na solid so- lution behavior. So, it is conclude that a typical bi-phasic transition occurred in the NaTi₂(PO₄)₃ electrode during cycling in accordance with previous electrochemical and computational studies [63].

TheoperandoPXRD data were further analyzed by sequential Riet- veld refinement using structural models for the two rhombohedral phases, i.e. Na-poor NaTi2(PO4)3 and Na-rich Na3Ti2(PO4)3. Fig. 5c–e shows the evolution phase fractions (wt%), unit cell parameters and unit cell volume for NaTi2(PO4)3 and Na3Ti2(PO4)3. During discharge, the wt

% of the two phases change linearly and symmetrically around 50%

state-of-charge as expected. During the charge process, the reformation of the Na-poor NaTi2(PO4)3 is lacking somewhat, i.e. the changes in wt%

of the two phases are not linear and symmetrical around 50 state-of- charge. As the sum of the wt% for the two phases by definition is unity, the observation may originate from the Na-poor NaTi2(PO4)3

initially forming in a less crystalline state, e.g. small nanocrystals, which may form in high numbers but are not large enough to diffract the X- rays. Hence, they need to grow in size to be observed by the SR-PXRD.

The unit cell parameters show that the phase transition is indeed bi- phasic with a relatively large volume misfit between the Na-poor and Na-rich phase. During the majority of the charge–discharge process the unit cell parameters of NaTi2(PO4)3 area =8.48 Å andc =21.84 Å, they

for Na3Ti2(PO4)3 area =8.86 Å andc =21.66 Å, respectively. It should be noted that at the initial stage of discharge, some solid solution behavior (continuously changing unit cell parameters) are observed for the Na-poor NaTi2(PO4)3, i.e. a small amount of Na-ions can be inter- calated before the phase transition is induced. Similarly, some solid solution bahavior is observed near the end of discharge for the Na-rich Na3Ti2(PO4)3, i.e. the phase is stable with a small amount of Na- vacancies. Refinement profiles at different change-discharge points are shown in Fig. 6.

Based on the operando synchrotron XRD analysis, It can be confirmed that the strong reversible crystal structure and Ti’s consistent valence change were both the contributors to the battery performance of the material. Although the overall patterns remain same, the fact that the indexed peaks change reversibly throughout the charge/discharge pro- cess implies that the NASICON crystal structure is highly maintained during electrochemical activities. The sequentially refined PXRD pat- terns demonstrate the reversibility of the two phases with stable elec- trochemical activities, which is particularly helpful for the material’s mechanical stability.

3.4. Ionic conductivity of ceramic electrolyte (Na_3.16Zr_1.84Y_0.16Si₂PO₁₂) Fig. 7 presents the typical impedance dataset obtained at 30 ^◦C.

Fig. 7a, corresponding to the impedance complex plane plots (Z’ vs Z’’), which depicts two semicircles, a big one at low frequencies and a smaller one at the high frequency range (close to the origin). Plots of log conductance, Y’, against log frequency (Fig. 7b), show a frequency in- dependent plateau over a relatively wide frequency range. At low fre- quency, a dispersion occurs and, as described below, this is associated with the charge blocking at the sample-ion blocking electrode.

The plot of the real part of capacity, C’, against frequency (Fig. S5), shows a high frequency plateau with a value in the range of ≅10⁻¹⁰F, which is too high to be assigned to the bulk; hence, it was assigned to grain boundary [64]. At lower frequencies, another plateau is observed, with values around 10⁻⁶F, being assigned to the sample-electrode interface (blocking of mobile ions in the electrodes). These assign- ments are corroborated when performing electrical measurements with non-blocking electrodes (Fig. S6), where the high impedance contribu- tion at low frequency is considerably reduced and the impedance at high frequency (assigned to grain boundary) is similar to that obtained with blocking electrodes. Consequently, the sintered ceramics present total impedance (grain boundary +bulk) values very low (around mΩ⋅cm), in agreement with reference [45], making them ideal to be used as an electrolyte in an all-solid-state battery.

Fig. 7 presents the trend of the total ionic conductivity as a function of temperature, whose values are taken from the high frequency impedance arcs shown in Fig. 7c and 7d. An equivalent circuit, which is used to examine impedance spectra, is utilized to match the experi- mental data (Fig. 7e). It is associated with series of different resistances (R) and constant phase elements (CPE) in it. The inclined line observed at low frequencies is related to the constant-phase element attributed to the blocking interfaces between the gold-ink-covered NASICON elec- trolyte [65]. Similar kind of equivalent circuit is reported in literature in case of NASICON electrolytes [66,67]. The diameter of the middle- frequency semicircle related to the interfacial resistance (R2), and the X-axis intercept at the high-frequency semicircle associated to the total resistance (R0 +R1) [68,69]. The ionic conductivity of the solid-state NASICON electrolyte increases as temperature does, with values ranging from 2.1x10⁻² mS/cm at − 30 ^◦C to 1.8 mS/cm at 90 ^◦C. The trend of the ionic conductivity as a function of temperature follows an Arrhenius behavior (Fig. 7f), with activation energy estimated as 0.28 eV, in range with NASICON electrolytes previously reported [70–72].

The conductivity of our synthesized solid-state electrolyte is comparable with previously reported literature as shown in Table 1 [73–94].

Fig. 6. Selected Rietveld-refined operando PXRD data at different charge–discharge regions.

3.5. All-solid-state battery performance

The excellent electrochemical performance of the NTPC2 electrodes in half cells encouraged us for the development of all solid-state sodium- ion cells to validate the potential use of the electrolyte. The first char- ge–discharge voltage patterns at C/10 are shown in Fig. S7. The applied pressure utilized to manufacture the cells is reduced from 800 to 600 psi by considering for the thickness of the solid-state electrolyte. Surpris- ingly, the formed solid-sate cell has a first discharge capacity of 95.1 mAh/g, which is nearly equal to the value for the liquid-electrolyte cell.

A minor drop in conductivity results in a little increase in polarization.

The dQ/dV curve of first discharge process shows clear redox peak at roughly 2.1 and 2.15 V (vs. Na⁺/Na) (Fig. S8). The second and third discharge dQ/dV curves are equivalent to the first discharge curve, showing that the solid-state based half cells act similarly to the liquid electrolyte-based cells.

As illustrated in Fig. 8a, b, the galvanostatic charge–discharge of the electrodes is evaluated at various current rates. Specific capacities decrease as current density increases; for example, at densities of C/5, C/

2, and C, the reversible capacity could reach 56.6, 38, and 28 mAh/g, respectively, which could be caused by the large polarization due to high current density [95]. Polarization in case of solid-state electrolyte is also

greater than in previously tested NTPC1 and NTPC2 electrodes tested in 1 M NaClO4/PC electrolyte (Fig. S9). Furthermore, as shown in Fig. 8a, the cell has outstanding capacity recovery capabilities when the current density is switched to C/10. The coulombic efficiency remained close to 100 % throughout the cycle experiments. It also shows that cell keeps a high discharge capacity even after charging/discharging at high current rates, indicating the extraordinary stability of NaTi2(PO4)3/C and its interface with the solid electrolyte.

Fig. 8c displays the cycle performance of a solid-state cell. Cells are tested 950 cycles at a current rate of C/5. The solid-state has almost the same initial discharge capacity as previously developed standard half cells with liquid electrolyte. Despite a capacity drop during the first couple of cycles, it displays an excellent and stable capacity retention of 28.3% up to 950 cycles. Even the voltage profiles are sustained after 950 cycles as well (Fig. S10). The well-synthesized solid-state electrolyte and the solid–solid interface are most likely responsible for the exceptional performance. As a result, sodium ion may intercalate and deintercalate more quickly. NaTi2(PO4)3 has a potential future as an electrode in sodium-ion batteries due to its excellent performance. Coulombic effi- ciency is depicted in Fig. 8c, and it is almost always close to 100%.

Because the SEI formation on the electrode surface takes time to form and activate, the initial coulombic efficiency may be low and gradually Fig. 7. Impedance and ionic conductivity obtained for the NASICON electrolyte at different temperatures. (a) Nyquist impedance plot at 30 ^◦C. (b) Real component of conductivity plot at 30 ^◦C. (c, d) Nyquist impedance plots as a function of temperature. (e) Equivalent circuit. (e) Arrhenius-like curve.

increase [96]. Coulombic efficiency is a measure of a cell’s cycling ability. The coulombic efficiency is still at 100 % after 950 cycles, indicating that the solid-state cell is stable.

The differences in the electrical, electrochemical and mechanical properties of a solid electrolyte compared to the more familiar liquid electrolytes are key to the challenges in all-solid-state batteries. At room temperature, the conductivity of a solid electrolyte is usually at least two or three orders of magnitude lower than that of a liquid electrolyte [97,98]. This can result from the solid electrolyte’s intrinsic properties or from existing grain boundaries. The sintered ceramics present total impedance (grain boundary +bulk) values very low (around mΩ⋅cm), in agreement with reference [45], making them ideal to be used as an electrolyte in an all-solid-state battery. Nevertheless, the ionic

conductivity is, in general, lower than conventional liquid electrolytes.

These are the reasons for excellent electrochemical performance in case of solid-state electrolyte in comparison with liquid electrolyte. The performance of our solid-state battery is compared with the literature (Table T2). It shows our solid-state battery shows good performance, especially longer cyclic performance at room temperature as compare to reported articles.

4. Conclusions

In summary, a solid-state technique was used to create a series of carbon-coated NaTi2(PO4)3 samples for use as electrode materials in a sodium-ion battery. The carbon covering shell significantly increases the electrochemical performance of the material. The carbon coating has no influence on the samples’ fundamental atomic structure, which are all pure phase NaTi2(PO4)3. It also has high cycle performance and low polarization. By speeding the movement of electrons and sodium ions, the carbon coating technique increases the sodium storage performance of NaTi2(PO4)3. This phenomenon will persist until all solid-state cells are made. All-solid-state cells constructed using Na3.16Zr1.84Y0.16Si2PO12

as the solid electrolyte and NaTi₂(PO₄)₃/C as cathode exhibit, at room temperature, excellent Coulombic efficiency, high specific capacity, and reasonable cycle stability. This is probably due to (i) low total ionic conductivity of the electrolyte (ii) the effective ion exchange in the interface and (iii) to a minimized stress caused by the less volume change of electroactive material.

Through the use of operando PXRD, it is demonstrated for the first time that the reversible structural evolution of NaTi2(PO4)3 ⇋ Na3Ti2(PO4)3 during discharge and charge follows typical two-phase reaction path with small solid solution behavior in both phases close to the fully discharged and charged states, respectively. In addition, the Na-poor NaTi2(PO4)3 initially forms as tiny nanocrystals, which can form in large numbers but are too small to diffract X-rays. As a result, at 50 state-of-charge, the change in wt% for the two phases are neither symmetrical or linear. This work, which is based on the NASICON design, effectively presents a proof of concept for a prospective all-solid- state Na-ion battery as a high-performance energy storage device.

Table 1

Conductivities of various reported sodium solid-state electrolytes.

Electrolyte Conductivity (S/cm) at room

temperature Ref.

Na3.2Zr2Si2.2P0.8O12-0.5NaF 3.6 ×10⁻³ [73]

Na3Al2P3O12 3.3 ×10⁻⁷ (at 100 ^◦C) [74]

Na1.5Al0.5Ge1.5P3O12 9.27 ×10⁻⁵ (at 140 ^◦C) [75]

Na11Sn2PS12 1.4 ×10⁻³ [76]

Na3Zr2Si2PO12 1.935 ×10⁻⁴ [77]

Na3.33Zr1.67Sc0.29Yb0.04Si2PO12 1.62 ×10⁻³ [78]

Na1.8Al0.8Ge1.2P3O12 3.8 ×10⁻⁴(at 300 ^◦C) [79]

Na3.4Zr2Si2.4P0.6O12 4 ×10⁻³ [80]

Na3.5Cr0.5Ti1.5(PO4)3 8.52 ×10⁻⁴ [81]

Na3.4Sc0.4Zr1.6(SiO4)2(PO4) 4 ×10⁻³ [82]

Na3.5Zr0.5Sc1.5(PO4)3 9.47 ×10⁻⁴ [83]

Na10SnP2S12 0.4 ×10⁻³ [84]

Na3.4Sc2(SiO4)0.4(PO4)2.6 8.3 ×10^{− 4} [85]

Na2Ti2SiP2O12 1 ×10^{− 4} [86]

Na1.5Sn0.5Ge1.5(PO4)3 8.39 ×10⁻⁵ [87]

Na10SnSb2S12 0.52 ×10⁻³ [88]

Na3.4AlScSi0.4P2.6O12 1.1 ×10⁻⁶ [89]

Na5.1FeB0.1Si3.9O12 1.69 ×10⁻³(at 300 ^◦C) [90]

Na1.5Y0.3Ga0.2Ge1.5(PO4)3 1.15 ×10⁻⁵ [91]

Na3.1Zr1.95Mg0.05Si2PO12 3.5 ×10⁻³ [92]

Na2(Ga0.1Ge0.9)2Se4.95 ~10⁻⁵ [93]

Na3P0.62As0.38S4 1.46 ×10⁻³ [94]

Na3.16Zr1.84Y0.16Si2PO12 0.202 ×10⁻³

1.8 ×10⁻³(at 90 ^◦C) Present work

Fig. 8. (a) Rate capability of solid-state cell between C/10 and 10C. (b) Galvanostatic charge/discharge profiles at various current rates. (c) Cycling performance and Coulombic efficiency measured at C/5 for 950 cycles.

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.

Data availability

Data will be made available on request.

Acknowledgements

Bidhan Pandit acknowledges the CONEX-Plus programme funded by Universidad Carlos III de Madrid (UC3M) and the European Commission through the Marie-Sklodowska Curie COFUND Action (Grant Agreement No 801538). This work has been also supported by the Madrid Gov- ernment (Comunidad de Madrid-Spain) through two projects: 1) the Multiannual Agreement with UC3M (“Fostering Young Doctors Research”, CIRENAICA-CM-UC3M), and in the context of the VPRICIT (Research and Technological Innovation Regional Programme); 2) DROMADER-CM (Y2020/NMT6584). Authors would like to thank the Agencia Espanola de Investigaci˜ ´on/Fondo Europeo de Desarrollo Regional (FEDER/UE) for funding the projects PID2019-106662RBC43.

The Danish Ministry for Research and Education is thanked for funding travel expenses related to the work at MAXIV through the Instrument Centre DanScatt. We acknowledge MAX IV Laboratory for time on Beamline DanMAX under Proposal [20211012]. Research conducted at MAX IV is supported by the Swedish Research council under contract 2018-07152, the Swedish Governmental Agency for Innovation Systems under contract 2018-04969, and Formas under contract 2019-02496.

DanMAX is funded by the NUFI grant no. 4059-00009B. Authors also acknowledge Universidad Carlos III de Madrid (Read & Publish Agree- ment CRUE-CSIC 2023) for funding the article processing charge (APC) to make this article open access.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.

org/10.1016/j.cej.2023.144509.

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