H_1.07Ti_1.73O₄-Derived Porous Plate-like TiO₂ as High-Performance Bifunctional Anodes for Lithium- and Sodium-Ion Batteries

Abstract

Porous plate-like anatase TiO₂ particles were synthesized through a direct calcination approach using layered titanate H₁.₀₇Ti₁.₇₃O₄ as a precursor. By controlling the calcination temperature (400 °C, 500 °C, and 600 °C), the morphology and electrochemical properties of the TiO₂ samples were effectively tuned. When evaluated as anodes for lithium-ion batteries (LIBs), the porous TiO₂ materials demonstrated markedly improved rate performance compared to commercial nano-TiO₂ (n-TiO₂). Specifically, at a high current density of 5.0 A/g, p-TiO₂-500 and p-TiO₂-600 delivered discharge capacities of 70.5 mAh/g and 87.5 mAh/g, respectively, far exceeding the 27.7 mAh/g achieved by n-TiO₂. The corresponding capacity retention rates at this rate were 30.1% for p-TiO₂-500, 41.2% for p-TiO₂-600, and only 16.4% for n-TiO₂. The enhancement in kinetics is ascribed to the unique porous plate-like architecture, which promotes efficient ion transport and introduces significant pseudocapacitive contributions. When applied as anodes for sodium-ion batteries (SIBs), p-TiO₂-600 exhibited the most promising performance. This study underscores the potential of porous plate-like TiO₂ as a high-performance bifunctional anode material suitable for both LIBs and SIBs.

1. Introduction

Carbon materials are promising anodes for LIBs due to their tunable properties, low cost, and abundance. However, graphite suffers from low theoretical capacity, while disordered carbons exhibit poor initial Coulombic efficiency (ICE), driving the search for higher-performance alternatives [1,2]. The development of anode materials that exhibit high-rate capability and enhanced safety is imperative for driving progress in the energy sector. Success in this endeavor is paramount for optimizing the multifaceted performance metrics of energy storage devices, namely energy density, cycle life, cost-effectiveness, and safety [3,4,5,6]. TiO₂ has become a widely studied anode material for lithium/sodium-ion batteries due to its intrinsic safety brought by high working voltage (~1.75 V for Li, ~0.7 V for Na), simple synthesis method, and low cost [7,8]. As a semiconductor, TiO₂ has a wide bandgap (3.2 eV), which leads to extremely low intrinsic electron/ion conductivity, severely restricting ion diffusion kinetics and electrochemical performance, resulting in poor rate performance and cycling stability [9]. In order to solve this problem, various strategies have been proposed to improve the lithium-ion diffusion pathway of TiO₂ to increase its energy density [10,11]. Common methods include designing composite materials, morphology control, surface modification engineering, metal element doping, etc. [12]. Composite fabrication through the application of a conductive coating on TiO₂ nanoparticles has been shown to facilitate rapid ion diffusion [13,14]. Notwithstanding this advantage, the high surface energy and subsequent agglomeration of these nanomaterials adversely impact their rate performance and long-cycle stability [15].

Constructing porous anatase TiO₂ is also a significant strategy for enhancing the properties of TiO₂ [16,17]. It is reported that the preparation of plate-like TiO₂ is an effective method for enhancing the kinetics of electrochemical reactions [18,19]. Huang. H et al. [20] coated TiO₂ with carbon nanotubes (CNTs), which improved its electrical conductivity. Jipeng Wang et al. [21] used carbon spheres as templates and synthesized TiO₂ with a hollow-sphere morphology structure by the sol–gel method. The excellent morphology effectively improved its lithium storage performance. In addition, some metals, such as germanium (Ge) and nickel (Ni), have higher capacity and ion diffusion rates than carbon-based ones. Kim et al. prepared TiO₂/Ge nanocomposites through the sol–gel process, which provided better lithium storage capacity [22]; Xu et al. [23] synthesized highly ordered 3D Ni-TiO₂ core–shell nanoarrays (NTNAs), which demonstrated excellent electrochemical performance as the anode of sodium-ion batteries. Additionally, numerous studies have reported the potential applications of porous TiO₂ in the field of lithium-ion and sodium-ion batteries [24,25,26,27]. However, the above-mentioned methods are cumbersome to operate and costly. There are still considerable challenges if these preparation methods are put into large-scale production.

In this work, porous plate-like anatase TiO₂ was successfully fabricated, and its morphology was finely tuned by controlling the reaction temperature to enhance the electrical conductivity and cyclability of TiO₂. The porous plate-like anatase TiO₂ offers the following advantages. (1) The porous plate-like structure enhances dynamic performance. (2) The porous structure induces pseudocapacitive lithium/sodium storage behavior, thereby improving the electrochemical performance. This study systematically investigated the structure and electrochemical properties of porous plate-like anatase TiO₂ electrodes. The influence of temperature on their morphology and electrochemical performance was thoroughly discussed. This also provides a new strategy for titanium-based oxide electrodes.

2. Experiment

2.1. Preparation of Porous Plate-like TiO₂

A homogeneous stirring of 5.1 g KOH, 0.6 g LiOH-H₂O, 6.9 g TiO₂, and 25 mL of deionized water yields layered K_0.81Li_0.27Ti_1.73O₄·H₂O (KLTO) under hydrothermal reaction at 250 °C for 24 h. Then, KLTO was protonated in nitric acid solution for 12 h to obtain the layered titanate H_1.07Ti_1.73O₄·H₂O (HTO) precursor; the preparation process is shown in Figure S1. Subsequently, the layered titanate HTO was heat-treated at 400 °C, 500 °C, and 600 °C for 3 h to obtain three porous plate-like TiO₂ samples (Figure 1).

2.2. Characterization

The crystal structure and phase composition of the samples were characterized by X-ray diffraction (XRD, TD3500, Dandong Tongda Technology Co., Ltd., Dandong, China. The morphology, size, and elemental distribution of the electrodes before and after cycling were examined using a field emission scanning electron microscope (SEM, FEI Apreo S, The FEI Company of the United States, Hillsboro, OR, USA) equipped with energy-dispersive X-ray spectroscopy (EDS). Transmission electron microscopy (TEM, FEI Tecnai G2 F20 S-TWIN, The FEI Company of the United States, Hillsboro, OR, USA) was performed at an acceleration voltage of 200 kV to further analyze the microstructure of the powder samples. The specific surface area was determined through N₂ adsorption–desorption measurements (ASAP 2460, Micromeritics, American Mc Instrument Company, Sebastian, FL, USA) and evaluated using the Brunauer–Emmett–Teller (BET) method. Phase transition behavior and thermal stability were investigated by thermogravimetric analysis (TGA, NETZSCH STA 409 PC, Germany NETZSCH, Selb, Germany).

2.3. Preparation of the Electrode

The p-TiO₂ powder was blended with acetylene black and polyvinylidene fluoride (PVDF) binder in a weight ratio of 7:2:1 using N-methyl-2-pyrrolidone (NMP, Comio Chemical Reagents Co., Ltd., Tianjin, China) as the solvent to form a homogeneous slurry. The resulting slurry was coated onto copper foil and dried in a vacuum oven at 70 °C for 48 h. The dried electrode film was then punched into 10 mm diameter disks. CR2032-type coin half-cells were assembled in an argon-filled glove box (LABSTAR, Braun Inert Gas Company of Germany, Munich, Germany), using a microporous polypropylene membrane as the separator and lithium metal foil as the counter/reference electrode. The electrolyte consisted of 1 M LiPF₆ dissolved in a mixture of ethylene carbonate (EC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) in a volume ratio of 1:1:1.

2.4. Electrical Measurement

Galvanostatic charge–discharge measurements were conducted within a voltage window of 0.01–3.0 V (vs. Li/Li⁺) at various current densities ranging from 0.1 to 5.0 A/g using a LAND battery test system (Wuhan Blue Power Electronics Co., Ltd., Wuhan, China). Cyclic voltammetry (CV) and pseudocapacitance analysis were carried out on a CHI660E electrochemical workstation (Shanghai Chenhua Instrument Co., Ltd., Shanghai, China) over the same voltage range, with scan rates varying from 0.1 to 1.0 mV·s⁻¹. Electrochemical impedance spectroscopy (EIS) was also performed using the same CHI660E workstation.

3. Results

The XRD curves of the p-TiO₂ prepared at different reaction temperatures are shown in Figure 2a. The samples were directly heat-treated from layered titanate HTO and were able to obtain pure phase anatase TiO₂ (p-TiO₂-400, p-TiO₂-500, and p-TiO₂-600) after heat treatments at 400 °C, 500 °C, and 600 °C. During the heat treatment of HTO powder, firstly, with the gradual increase in the muffle furnace temperature, the interlayer H₃O⁺ and H⁺ began to be detached from the interlayer, accompanied by the weakening of the peak intensity of the characteristic peaks of the laminar structure near 2θ = 9.3, and shifted to a small angle, which signified that the laminar structure of lamellar titanate began to collapse, and the characteristic features of the laminar structure completely disappeared at 200 °C., and the formation of anatase TiO₂ and rutile TiO₂ began to appear. As the temperature continued to rise, the content of anatase TiO₂ gradually increased, and when the temperature was higher than 400 °C, HTO was completely transformed into anatase TiO₂. p-TiO₂-400, p-TiO₂-500, and p-TiO₂-600 showed an increase in the crystallinity of samples with increasing temperature. The commercial nano TiO₂ (n-TiO₂) was also examined by XRD, and the characteristic peaks of the four groups of samples corresponded well to the (101), (0 0 4), (20 0), and (211) crystalline surfaces of anatase TiO₂, which were consistent with the standard card Anatase TiO₂ Jade 21-1272 [28], and all of them were anatase TiO₂. The crystal structure of anatase TiO₂ is shown in Figure 2b, where each titanium atom forms an octahedral coordination (TiO₆) with six oxygen atoms, which has four formulaic units, where each TiO₆ octahedron is connected to eight octahedra through four common vertices and four edges, thus constituting the tetragonal anatase crystal structure of I₄1/amd.

We have carried out the morphological analysis of three porous plate-like TiO₂ and commercial nano TiO₂, and the SEM images are shown in Figure 3. The layered titanate precursor is a smooth plate-like structure with an overall particle size of about 1–2 μm (Figure 3a); the three porous plate-like TiO₂ obtained after heat treatment at different temperatures as a titanium source are shown in Figure 3b–d. It can be seen that p-TiO₂-400 obtained at 400 °C maintains the plate-like structure, the overall particle size does not change significantly, and the porous structure does not appear on the plate surface. p-TiO₂-500 obtained by heat treatment at 500 °C has an overall particle size of about 1.5 μm, and a large number of uniformly distributed pores on the plate surface. p-TiO₂-600, obtained by heat treatment at 600 °C, has the largest number of pores and the largest pore diameter on the particle surface, and has the largest number of holes and the largest pore size; the plate-like structure appears to collapse and fracture under the influence of temperature, and the edges of the plate-like particles are slightly smooth. Figure 3e shows commercial nano-TiO₂ (n-TiO₂), and the powder is 50–80 nm nanorods with obvious agglomeration.

Figure 4 shows the TEM images of the four samples, from which it can be seen that the TiO₂ prepared with HTO as precursor maintains the plate-like morphology after heat treatment at different temperatures. With the increase in temperature, the micropores increase, and the particle size increases. n-TiO₂ is 50~80 nm nanoparticles with an obvious agglomeration phenomenon. The EDS analysis graph of p-TiO₂-500 is shown in Figure S2 in the Supporting Information.

Subsequently, we tested the specific surface area of four groups of samples, which were 9.7 cm²/g, 40.6 cm²/g, 42.6 cm²/g, and 13.3 cm²/g, respectively, as shown in Figure S3. The results showed that the specific surface area of TiO₂ with two porous structures, p-TiO₂-500 and p-TiO₂-600, was higher than that of n-TiO₂. p-TiO₂-400 samples had a larger particle size and did not form a porous morphology, so the specific surface area was the lowest. The specific surface area of the samples affects their electrochemical performance to a certain extent, and a larger specific surface area is favorable for adsorbing more ions during the electrochemical reaction, which makes the reaction occur more fully and improves the performance of the cells.

Then, we analyzed the formation process of nanostructures of two samples, p-TiO₂-500 and p-TiO₂-600. The precursor HTO layer is intercalated with H₃O⁺, H⁺, and water molecules. The structure is a two-dimensional nanosheet layer composed of [TiO₆] octahedra. As the temperature rises, the physically adsorbed water and some bound water between the layers first desorb. Subsequently, the H₃O⁺ and H⁺ ions between the layers began to detach. From the XRD graph, it can be observed that the characteristic peak intensity of the HTO-layered structure weakens and shifts to a smaller angle, and eventually disappears completely. It indicates that the layer spacing changes due to the loss of water molecules. Between 200 and 400 degrees Celsius, the -OH and H⁺ in the precursor begin to condense and then escape, leaving oxygen vacancies and unsaturated bonds in the [TiO₆] structure. Subsequently, Ti and O atoms undergo migration and rearrangement, forming Ti-O-Ti covalent bonds and undergoing topological transformation. At this point, the macroscopic plate-like morphology is retained, but the crystal structure of the surrounding area has changed. When the temperature exceeds 400 °C, HTO completely transforms into rutile-type TiO₂. When the temperature is above 500 °C, the rapid escape of water vapor and the internal stress generated by structural reorganization are released by forming uniformly distributed pores. The structure undergoes local collapse, fracture, and reconstruction, forming a porous plate-like structure. When the temperature reaches 600 °C, atomic migration is too rapid and sintering takes the dominant role. Some plate-like structures collapse and fracture excessively, as shown in Figure 5.

3.1. Lithium-Ion Battery Performance Testing and Analysis

The multiplicity performance of the lithium-ion half-cells was first tested (Figure 6a). p-TiO₂-500 and p-TiO₂-600 both exhibited superior multiplicity performance to n-TiO₂, while p-TiO₂-400 had a slightly lower capacity than n-TiO₂ at small current density, but gradually outperformed n-TiO₂ as the current was increased. Overall, the multiplicative performance of all three groups of porous plate TiO₂ is better than that of n-TiO₂, and p-TiO₂-500 obtains the most excellent multiplicative performance due to its optimal morphology. The details of the charging/discharging capacities of all the samples at 0.1 A/g, 0.2 A/g, 0.5 A/g, 1.0 A/g, 2.0 A/g, and 5.0 A/g recovered to 0.1 A/g current density and can be seen as shown in

Table 1. Comparison table of the rate performance of four TiO₂ samples.

Material	Discharge/Charge	The Capacity (mAh/g) at Current Density (A/g) of:
		0.1	0.2	0.5	1.0	2.0	5.0	0.1
p-TiO2-400	Discharge	152.3	122.1	92.3	76.9	66.7	54.2	159.4
	Charge	154.2	122.1	92.3	76.9	66.7	54.2	159.4
p-TiO2-500	Discharge	228.8	196.3	153.9	127.5	101.6	70.5	226.9
	Charge	234.5	201.2	155.7	127.6	101.6	70.5	228.6
p-TiO2-600	Discharge	208.6	178.1	149.0	126.4	109.5	87.5	217.8
	Charge	212.3	178.3	149.3	126.4	109.5	87.5	218.6
n-TiO2	Discharge	166.6	142.3	90.2	65.6	47.6	27.7	162.0
	Charge	169.2	143.9	91.6	65.8	47.6	27.7	162.2

The cycling performance at a current density of 0.1 A/g was subsequently tested, as shown in Figure 6b. The three porous plate TiO₂ exhibited higher electrochemical performance than n-TiO₂. Among them, p-TiO₂-500 has the best performance, maintaining a reversible capacity of 350 mAh/g after 100 charge/discharge cyclings. p-TiO₂-600, p-TiO₂-400, and n-TiO₂ can maintain a reversible capacity of 263 mAh/g, 207.5 mAh/g, and 169 mAh/g after 100 cyclings, respectively.

The p-TiO₂-500, p-TiO₂-600, and n-TiO₂ were tested for 5000 charge/discharge cyclings at a high current density of 1.0 A/g, as shown in Figure 6c. The results show that the p-TiO₂-500 and p-TiO₂-600 samples can still maintain a reversible capacity of 112.3 mAh/g and 111.8 mAh/g, which are both higher than that of n-TiO₂ (61.6 mAh/g), with a performance enhancement of 82.30% and 81.49%, respectively.

Figure 7 shows the capacity/voltage curves and CV curves for the n-TiO₂, p-TiO₂-500, and p-TiO₂-600 samples. The capacity/voltage curves of the three groups of samples at 0.1 A/g current density for the first three laps, the fifth lap, and the tenth lap are shown in Figure 7a,c,e. The first discharges were 493.1 mAh/g, 612.6 mAh/g, and 518.8 mAh/g, and the Coulomb efficiencies were 41.53%, 38.96%, and 39.86%, respectively. The capacity decay of n-TiO₂ is significantly more pronounced, as it can only sustain a capacity of 164.1 mAh/g by the 10th cycle. In contrast, p-TiO₂-500 and p-TiO₂-600 are able to maintain capacities of 205.3 mAh/g and 202.6 mAh/g at the same point, respectively, exhibiting no significant signs of capacity decay; their curves nearly overlap. All three samples exhibit an embedded lithium platform near 1.7 V, which belongs to the redox reaction of Ti⁴⁺/Ti³⁺ brought about by the intercalation reaction of anatase TiO₂ with lithium ions at this working potential [29]. The electrochemical reaction equation is shown in Equation (1) [30]:

TiO₂ + xLi ⁺ +xe⁻→Li_xTiO₂

(1)

Figure 7b,d, and f illustrate the cyclic voltammetry (CV) curves of n-TiO₂, p-TiO₂-500, and p-TiO₂-600 at a scan rate of 0.1 mV/s, respectively. The characteristic peaks observed at 1.7 V and 2.0 V correspond to the reversible conversion between Ti⁴⁺ and Ti³⁺, which aligns with the lithium intercalation/de-intercalation plateau evident in the capacity/voltage profiles. The enhanced intensity of these characteristic peaks for both porous TiO₂ variants indicates a more efficient intercalation reaction during charge and discharge cyclings; consequently, p-TiO₂-500 and p-TiO₂-600 exhibit superior reversible capacities.

It has been shown that when the electrode material is porous, pseudo-capacitive Li⁺ storage behavior occurs during charging and discharging. Voltametric cycling tests were conducted on four distinct TiO₂ species at varying scan rates, as illustrated in Figure 8. The correlation between current and scan rate was analyzed to confirm the presence of pseudocapacitive behavior during both charging and discharging processes. For a redox reaction, the peak current obeys the following power law [31]:

i = av^b

(2)

In the formula, i, v, and a represent current (mA), scanning rate (mV/s), and an arbitrary constant, respectively, while b can determine whether there is pseudocapacitance behavior in the lithium storage process. As shown in Figure 8d, the logarithm of both sides of the equation gives a linear relationship between log (i) and log (v), and this linear equation has a slope of b value. According to the pseudocapacitance judgment theory, if b = 0.5, it means that only the intercalation reaction exists in the battery. If b = 1, it means that only pseudocapacitance lithium storage behavior exists. The pseudocapacitance was calculated as a function of current and scanning rate, and the b-values of the three TiO₂ samples were 0.79, 0.63, and 0.67, respectively, indicating that all three samples have a lithium storage mechanism that is controlled by both the intercalation reaction and the pseudocapacitance during the charge/discharge process.

In addition, we calculated the pseudocapacitance ratios of the three TiO₂ according to Equation (3) [32]:

i(V) = k₁v + k₂v^1/2

(3)

k₁v and k₂v^1/2 correspond to the pseudocapacitance and the current contribution of the nested process, respectively. The lithium storage behavior of pseudocapacitance brings great capacity contribution to the composite anode electrode materials at different scanning rates, and CV curves at different scanning rates are shown in Figure 9.

We have performed analytical tests related to the kinetics of electrochemical reactions; electrochemical impedance spectroscopy (EIS) was performed at a frequency in the range of 0.01 Hz~1000 kHz for three groups of samples and fitted the data to an equivalent circuit diagram as shown in the inset (Figure 10a, all the samples have been subjected to 100 charging and discharging cyclings at a current density of 0.1 A/g). It can be seen that p-TiO₂-500 has the lowest charge transfer resistance, p-TiO₂-600 has a slightly higher charge transfer resistance, and n-TiO₂ has the largest charge transfer resistance. Rs represents the internal resistance of the cell, and the diagonal line indicates the Warburg resistance (W), which is an important parameter for calculating the lithium-ion diffusion rate (D_Li+), which can be calculated by Equation (4) [33]. The details of the impedance and the lithium-ion diffusion coefficient for each part of the three sets of samples are shown in

Table 2. Comparison table of impedance of each part for n-TiO₂, p-TiO₂-500, and p-TiO₂-600, and their lithium-ion diffusion coefficients.

Sample	Rs (Ω)	Rct (Ω)	Z′ (Ω)	DLi+ (cm2/s)
n-TiO2	12.33	160.19	172.52	1.04 × 10−16
p-TiO2-500	11.98	90.51	102.49	5.99 × 10−16
p-TiO2-600	12.01	120.76	132.77	4.93 × 10−16

. As calculated, the two porous TiO₂ not only exhibit lower charge transfer impedance than n-TiO₂, but also possess a higher lithium-ion diffusion rate.

$$D_{Li+}={\displaystyle \frac{{R^{2}T}^2}{2{A^2{n}^4{F}^4{C}^2\sigma }^2}}$$

(4)

In the formula, R denotes a constant gas (8.314 J K⁻¹ mol⁻¹) and T denotes the ambient temperature (298.15 K), A denotes the surface of the electrode (0.785 cm²), and n denotes the number of electrons in the electronic transfer reactions. The Faraday constant (96,486 C/mol) is F, σ [34] is the slope of Z′, and with respect to ω^1/2 (Figure 7b), and C is the concentration of lithium-ion phase (mol/cm³). The capacity/voltage curves and dQ/dV curves of the three sets of samples are shown in Figure 10c,d. The two porous TiO₂ exhibit a more pronounced intercalation reaction plateau at 1.7 V in the discharge curves than n-TiO₂, as well as stronger intercalation reaction-embedded lithium peaks in the dQ/dV curves. p-TiO₂-500 and p-TiO₂-600 exhibit significantly improved electrochemical performance due to the higher degree of intercalation reaction.

3.2. Sodium-Ion Battery Performance Testing and Analysis

The electrochemical performance of p-TiO₂-500, p-TiO₂-600, and n-TiO₂ sodium-ion batteries is shown in Figure 11. The porous structured TiO₂ all demonstrated a performance enhancement to varying degrees compared to n-TiO₂, with the p-TiO₂-600 sample showing the most significant performance enhancement.

We tested the sodium-ion CV curves and capacity/voltage curves for three groups of samples. The CV curves of the three sets of samples for the first three revolutions at a scan rate of 0.1 mV/s show (Figure 12a,c,e) that all three sets of samples exhibit a de-embedding peak of sodium ions near 0.7 V, where an intercalation reaction occurs at this potential, corresponding to the reversible transition between Ti3⁺ and Ti4⁺ that occurs in Equation (5) [35]. The capacity/voltage profiles of the three groups of samples for the first three laps at a current density of 0.1 A/g, as well as the tenth lap, are shown in Figure 12b,d,f. The first-lap discharges of all the samples exhibit a large irreversible capacity loss:

TiO₂ + xNa⁺ + xe⁻→Na_xTiO₂

(5)

In order to determine the embedded sodium platform more accurately, we plotted the capacity/voltage curves of the three groups of samples after 100-turn charge–discharge cyclings at 0.1 A/g current density and differentiated the corresponding discharge curves (embedded sodium stage) in terms of dQ/dV, as shown in Figure 13. The results show that both porous TiO₂ exhibit higher charge/discharge capacities than n-TiO₂, with a downward-sloping line near 1.0–0 V in the embedded sodium stage for all of them. By differentiating the three discharge curves in terms of dQ/dV (Figure 13b), an embedded peak near 0.7 V is seen for all three groups of samples, which is consistent with the CV curve results. Compared with the corresponding dQ/dV curves tested for Li-ion batteries, the characteristic peak intensities of the dQ/dV curves for porous TiO₂ sodium-ion batteries are slightly higher than those for n-TiO₂, with the highest characteristic peak intensity for p-TiO₂-600 (the dQ/dV characteristic peak intensities for p-TiO₂-500 and p-TiO₂-600 Li-ion batteries, on the other hand, are significantly higher than those for n-TiO₂, and the characteristic peak intensities for p-TiO₂-500 have the highest characteristic peak intensity, Figure 10d). The difference between the two cells may be due to the lower intercalation reaction of the porous structure of the sodium-ion radius ambassador and the larger pore size of the p-TiO₂-600 surface micropores, which resulted in the highest electrochemical performance in the sodium-ion batteries.

Subsequently, we investigated the electrochemical kinetic properties of sodium ions across the three sample groups, as illustrated in Figure 14. The equivalent circuit fitting of the data indicates that the porous plate-like p-TiO₂-500 exhibits the lowest charge-transfer resistance, followed by p-TiO₂-600 with a slightly higher resistance, while n-TiO₂ demonstrates the highest resistance. These findings are consistent with those obtained from lithium-ion battery EIS tests. Additionally, we calculated the sodium ion diffusion rate (D_Na+) for all three sample groups in accordance with Equation (4).

The ion-diffusion relation coefficient Weber factor σ was obtained by fitting the linear relationship between Z′ and ω^1/2, and then the sodium-ion diffusion coefficient calculation was carried out. The details of impedance, as well as sodium-ion diffusion coefficients for each part of the three sets of samples, are shown in

Table 3. Comparison table of impedance of each part for p-TiO₂-500, p-TiO₂-600, and n-TiO₂ sodium-ion battery and their sodium-ion diffusion coefficients.

Sample	Rs (Ω)	Rct (Ω)	Z′ (Ω)	DNa+ (cm2/s)
n-TiO2	9.659	358.8	368.459	4.88 × 10−19
p-TiO2-500	8.853	143.1	151.953	1.09 × 10−18
p-TiO2-600	8.752	235.9	244.652	1.60 × 10−18

.

4. Conclusions

In this study, porous-platelet anatase TiO₂ was prepared using layered titanate as the titanium source and platelet-template heat treatment. Porous platelet TiO₂ exhibits higher electrochemical performance than n-TiO₂ in both lithium/sodium-ion batteries, with the most significant performance enhancement from the p-TiO₂-500 lithium-ion battery and the best electrochemical performance from the p-TiO₂-600 sodium-ion battery. The porous structure improves the kinetic properties and is accompanied by a large number of pseudocapacitive behaviors that work together and thus can bring a significant performance enhancement to them. The tests related to lithium/sodium-ion batteries confirm that this porous plate-like TiO₂ is a promising and versatile anode material for batteries.