*Article* **Synthesis and Characterization of Activated Carbon Co-Mixed Electrospun Titanium Oxide Nanofibers as Flow Electrode in Capacitive Deionization**

**Gbenro Folaranmi, Myriam Tauk, Mikhael Bechelany, Philippe Sistat, Marc Cretin \* and Francois Zaviska \***

Institut Européen des Membranes, IEM, UMR-5635, Université de Montpellier, ENSCM, CNRS, Place Eugène Bataillon, CEDEX 5, 34095 Montpellier, France; gbenro.folaranmi@umontpellier.fr (G.F.); myriam.tauk@umontpellier.fr (M.T.); mikhael.bechelany@umontpellier.fr (M.B.); philippe.sistat@umontpellier.fr (P.S.)

**\*** Correspondence: marc.cretin@umontpellier.fr (M.C.); francois.zaviska@umontpellier.fr (F.Z.)

**Abstract:** Flow capacitive deionization is a water desalination technique that uses liquid carbonbased electrodes to recover fresh water from brackish or seawater. This is a potential secondgeneration water desalination process, however it is limited by parameters such as feed electrode conductivity, interfacial resistance, viscosity, and so on. In this study, titanium oxide nanofibers (TiO2NF) were manufactured using an electrospinning process and then blended with commercial activated carbon (AC) to create a well distributed flow electrode in this study. Field emission scanning electron microscope (FESEM), X-ray diffraction (XRD), Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), and energy dispersive X-ray (EDX) were used to characterize the morphology, crystal structure, and chemical moieties of the as-synthesized composites. Notably, the flow electrode containing 1 wt.% TiO2NF (ACTiO2NF 1 wt.%) had the highest capacitance and the best salt removal rate (0.033 mg/min·cm2) of all the composites. The improvement in cell performance at this ratio indicates that the nanofibers are uniformly distributed over the electrode's surface, preventing electrode passivation, and nanofiber agglomeration, which could impede ion flow to the electrode's pores. This research suggests that the physical mixture could be used as a flow electrode in capacitive deionization.

**Keywords:** flow electrode capacitive deionization; electrospinning; activated carbon; desalination

#### **1. Introduction**

One of the growing challenges of the 21st century is the availability of fresh water. Water contamination because of anthropogenic activities such as industrialization, demographic change, and global warming has resulted in a significant increase in demand for safe drinking water. Water desalination technology could help to alleviate this problem by delivering high quality, pure water. Most desalination technologies, such as multiple effect desalination (MED), reverse osmosis (RO), and others, have high capital costs (when considering plant setup) and energy consumption (when considering pre- and post-treatment of water, as in RO), necessitating the development of a new desalination technique [1].

Capacitive deionization (CDI) is a growing desalination technology attracting attention as an energy-efficient, cost-effective, and ecofriendly water treatment technology. The first and most widely used CDI form involves a pair of porous carbon electrodes separated by a space in which salt water flows as an influent perpendicular to the applied electric field direction [2]. A fundamental variation of this basic CDI form emerged with the unfolding of membrane CDI (MCDI). In this architecture, ion exchange membranes were added to the CDI cell configuration to block co-ions from carrying parasitic current, which improves charge efficiency and can increase the charge storage in the electrodes porous structure [3]. In the past decade, a new class for CDI based on MCDI was developed which introduced

**Citation:** Folaranmi, G.; Tauk, M.; Bechelany, M.; Sistat, P.; Cretin, M.; Zaviska, F. Synthesis and Characterization of Activated Carbon Co-Mixed Electrospun Titanium Oxide Nanofibers as Flow Electrode in Capacitive Deionization. *Materials* **2021**, *14*, 6891. https://doi.org/ 10.3390/ma14226891

Academic Editor: Zhenghua Tang

Received: 6 October 2021 Accepted: 9 November 2021 Published: 15 November 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

carbon flow electrodes that can be pumped through the electrode compartments. Flow electrode CDI (FCDI) is a promising second-generation water desalination method based on the principle of ions adsorption. When a specific voltage is applied, it entails using a polarized flow electrode (liquid electrode) travelling via a flow channel to adsorb ions from brackish or seawater via electrical double layer (EDL) formation (interface between electrolyte and electrode) [4].

One of the challenges in FCDI is to improve the conductivity and capacitance of carbon materials used as electrodes, as well as reducing the resistance between the electrode and the electrolyte. This can be accomplished by improving the surface charge of activated carbon for better EDL formation [5].

Because of its beneficial qualities such as high surface area, porosity, availability, and low cost, graphitic carbon, materials such as activated carbon are frequently used as flow electrodes [6,7]. However, for more effective EDL formation and better ion storage, the resistance at the electrode-electrolyte interface should be lowered. As a result, surface modification is required to raise the surface charge of activated carbon.

Metal oxides such as zinc oxide (ZnO), tin oxide (SnO), zirconium oxide (ZrO), and titanium oxide (TiO2) have been used to modify carbon in this concept [8,9]. Among all these oxides, TiO2 has a high surface charge that can be used in conjunction with AC to create new types of electrodes. The aggregation of TiO2 particles, however, is one of the additive's downsides, limiting its performance [10–12]. Small quantities of TiO2 nanofibers are used in activated carbon flow electrodes to alleviate the agglomeration problem since they have a favorable large axial ratio morphology [13,14].

Due to the high resistance of the liquid medium utilized for carbon dispersion for slurry formation, charge transfer resistance is a strong factor hindering the ease of ions diffusion into the pores of carbon electrodes in flow electrodes [15]. As a result, methods for reducing interfacial resistance by adding conducting additives or introducing oxygenated functional groups to activated carbon have been reported in the literature [15–18], but no report on the properties of TiO2 nanofibers in reducing interfacial resistance and their effect on the rheological properties of carbon-based flow electrodes has been found to date.

Nanofiberous materials of high surface area and porosity are considered recently to be used for generating sustainable and green energy resources. Electrospinning method is a low-cost method used to create nanofiberous materials suited for many energy-related applications such as supercapacitors, and Li-ion batteries [19,20].

Motivated by this, we have synthesized TiO2 nanofibers by electrospinning process and mixed it with commercial AC. The proposed electrodes which consisted of different weight percentage of TiO2NF were labelled TiO2NF-x where x represents the weight percentage of TiO2NF in the composite (x = 0.5, 1.0, 1.5, 2.0, 2.5, and 5.0 wt.% TiO2NF). The electrodes were characterized by FESEM, XRD, Raman spectroscopy, XPS, EDX, and N2 adsorption/desorption. Dynamic viscosity of the mixtures in suspension to build the flow electrodes was also investigated. Hence, the electrochemical properties of the as-prepared flow electrodes (AC and TiO2NF-x) were analyzed by cyclic voltammetry (CV) to determine their capacitance after which the flow electrodes were tested in a laboratory-made FCDI cell for desalination.

#### **2. Materials and Methods**

#### *2.1. Materials*

Activated carbon, Darco (CAS no: 7440-44-0, MW 12.01 g mol−1), titanium (IV) isopropoxide (C12H28O4Ti, CAS no: 546-68-9, MW 284.22 g mol−1, 97%), poly(vinyl pyrrolidone) (PVP, CAS no: 9003-39-8, MW 1,300,000 g mol<sup>−</sup>1), and sodium chloride (NaCl, CAS no: 7647-14-5, MW 58.44 g mol−1, 99%) were obtained from Sigma Aldrich, Steinheim, Germany. Acetic acid (CH3COOH MW 60.052 g mol−1, 99.5% assay), absolute ethanol (C2H5OH MW 46.07 g mol<sup>−</sup>1, 99.99%,) were obtained from VWR chemicals, Fontenay-sous Bois, France. Polyvinylidenefluoride (PVDF) (CAS no: 24937-79-9) was obtained from Alfa Aesar, Erlenbachweg 2, Kendel, Germany. All reagents were used without any further

purification. Cationic and anionic exchange membranes were purchased from Membranes International Inc. (Ringwood, NJ, USA) and Deionized water (18 MΩ·cm) was used to prepare standard solutions and suspensions.

#### 2.1.1. Synthesis of TiO2 Nanofibers by Electrospinning

The TiO2 nanofibers were synthesized by electrospinning technique. Briefly, solution A containing 3 mL of ethanol was added to 0.3 g of polyvinyl pyrrolidone (PVP) and then stirred until dissolution. Then, solution B consisting of 2 mL ethanol, 2 mL of acetic acid, and 3 mL of titanium isopropoxide was stirred for 45 min. This is then followed by the addition of solution B into clear solution of A and then further stirred for 50 min to obtain a sol gel solution. Electrospinning of the sol–gel solution was carried out at 22.30 kV. The distance between the collector and the syringe tip was maintained at 10 cm, and the injection speed was 0.5 mL min−1. The obtained electrospun nanofiber materials were sintered in air atmosphere at 400 ◦C for 4 h with a ramp rate of 1 ◦C min<sup>−</sup>1.

#### 2.1.2. Preparation of AC and ACTiO2NF–x Flow Electrodes

For flow electrode preparation, certain amount of powdered commercial AC and ACTiO2NF-x (in which x corresponds to 1.0, 1.5, 2.0, 2.5, and 5.0 wt.% of TiO2NF) were weighed and dispersed in deionized water as presented in Table 1. The mixtures were sonicated for 3 h and stirred for 1 h before being fed into the cell. The tank containing the slurry electrode was continuously stirred on a magnetic stirrer during the course of the experiment.


**Table 1.** Composition of flow electrodes.

Note: FE: Feed electrode; DH2O: deionized water.

#### *2.2. Physical Characterization*

FESEM was used to analyze the morphology of the samples (Hitachi S4800, Tokyo, Japan). The structural properties were studied by using Raman spectroscopy (HORIBA Xplora, Tokyo, Japan) and XRD (Pan Analytical X'pert Phillips, Almelo, The Netherlands). XPS (ESCALAB 250 Thermo Electron, Montigny Le Bretonneux, France) and EDX (X-Max, Oxford, UK) were used to investigate the atomic composition and chemical moieties of the materials. For the XPS analysis, the excitation source was a monochromatic source Al Kα anode with photoenergy that was observed at 1486.6 eV. The analyzed surface has a diameter of 500 μm. The photoelectron spectra were calibrated in terms of bond energy with respect to the energy of the C=C component of carbon C1s at 284.4 eV. Surface area was obtained by using N2 adsorption/desorption at 77 K (Micromeritics ASAP, Verneuil en Halatte, France). Dynamic viscosity was measured using Anton Paar Rheometer Physica MCR 301 (Anton Paar, Graz, Austria).

#### *2.3. Electrochemical Characterizations*

Solid electrodes for electrochemical characterization of pristine AC and ACTiO2NF were made by combining activated carbon powder (0.32 g), carbon black (0.04 g), and poly (vinylidene fluoride fluoride PVdF, 0.04 g) in 3 mL N-Methyl-2-pyrrolidone (NMP). To establish homogeneity, the mixture was agitated for 2 h and then sonicated for 40 min.

After that, the slurry was applied on a graphite sheet. The coated electrode was dried in an oven for 1–2 h at 80 ◦C. The as-synthesized titanium oxide nanofibers were added (0.5, 1.0, 1.5, 2.0, 2.5, and 5.0 wt.%) to AC containing carbon black and PVDF, and the mixture was agitated for several hours in 3 mL NMP for ACTiO2NF synthesis. After that, the mixture was sonicated for 40 min. The slurry was subsequently immobilized by depositing it on a graphite sheet. It was then dried in an oven at 80 ◦C for 1–2 h to produce solid electrodes.

The electrochemical properties of the prepared electrodes were examined by using CV. CV tests were performed using a three-electrode system. The mixtures were deposited on a graphite sheet as support with an exposed surface area of 1 cm2, while a platinum mesh and a saturated 3 M KCl, Ag/AgCl electrode served as counter and reference electrodes respectively. A molar NaCl solution was used as electrolyte. Voltammetry measurements were performed with Origalys potentiostat (OGF01A, Origalys Electrochem SAS, Rillieuxla-Pape, France) at an operating window from −0.4 to 0.6 V vs ref (to ensure electrochemical stability of the electrolyte and prevent water splitting i.e., oxygen and hydrogen evolution) in 1 M NaCl electrolyte.

The double-layer capacitance was determined using cyclic voltammetry at different scan rates by considering the charging and the discharging currents at 0.1 V vs. ref. The determined double-layer capacitance of the system was the average of the absolute value of the slope of the linear plot of charging and discharging currents fitted to the data. CDL as the double layer capacitance was determined using Equation (1):

$$
\mathbf{i} = \boldsymbol{\nu} \,\, \mathbb{C}\_{\text{DL}} \,\, \text{.}\tag{1}
$$

The double-layer charging current density i (A·cm−2) is equal to the product of the scan rate <sup>υ</sup> (V·s<sup>−</sup>1), and the electrochemical double-layer capacitance CDL (F·cm<sup>−</sup>2).

#### *2.4. FCDI Measurement*

The schematic diagram of the close loop experimental set-up is shown in Figure 1 in which the cell was powered by a potentiostat and the feed solution (FS, 5 g·L<sup>−</sup>1) was made to pass through a spacer sandwiched in between cation and anion exchange membranes. The feed electrodes (FE) stored in a reservoir (were made to pass (by pumping) through flow channels and as they exit the channels, they are fed back to the reservoir and then re-circulated; this allows co-mixing of opposite charged ions outside the cell. Figure 2 shows the breakdown of the cell components with the current collector (6 mm width and 0.9 mm depth channel), the ions exchanges membranes and spacer (0.9 mm thick). The flow rate of the electrodes was operated at 40 mL min<sup>−</sup>1. Each desalination experiment was conducted for 30 min. The initial conductivity of the salt solution and that of the effluent was monitored at room temperature by ion conductivity meter (Hanna Instruments SRL) all along the 30 min of the experiment. A constant voltage difference of 1.2 V was applied to the FCDI unit cell using an Origalys potentiostat (OGF01A, Origalys Electrochem SAS) for desalination experiments. The current intensity passing through the FCDI unit cell was consequently measured by the potentiostat during the experiment.

In the present work, the following indicators defined FCDI performances:

Salt removal rate in mg·min−1·cm−<sup>2</sup> (SRR) relates to the mass of salt adsorbed (mg) per FE-FS contact area (cm2) per unit of time (min). It is calculated by Equation (2):

$$\text{SRR} = \frac{(\text{Co} - \text{Cf}) \ast \text{V}}{\text{A} \ast \text{t}} \tag{2}$$

Co and Cf are the initial and final (at t = 30 min) concentration (mol·L<sup>−</sup>1), respectively, V is the volume of the solution (L), A is the contact area between FE: FS, and t is the charging time.

The salt removal efficiency in % (SRE) was calculated using Equation (3):

$$\text{SRE} = \frac{(\text{Co} - \text{Cf})}{\text{Co}} \ast 100\tag{3}$$

Charge efficiency (CE) in % which relates to the ratio of salt adsorbed to the quantity of charge passed into the system was calculated by Equation (4):

$$\text{CE} = \frac{\text{z} \left(\text{Co} - \text{Cf}\right) \text{V F}}{\text{M} \int \text{Idt}} \ast 100 \tag{4}$$

where z is the equivalent charge of the ions, F is the Faradaic constant, M is the molar weight, and I*d*t is the integrated quantity of charge passed to the system as a function of time. CE is a good indicator for the energy efficiency of the system and will directly affect the operating cost of the system (OPEX).

**Figure 1.** Schematic diagram of FCDI setup is a figure.

**Figure 2.** Individual components of FCDI cell.

#### **3. Results and Discussion**

#### *3.1. Morphology*

The morphology of as-synthesized titanium oxide nanofibers (TiO2NF), pristine AC, and the ACTiO2NF-x electrodes are shown in Figure 3a–h respectively. From Figure 3a, it is apparent that the electrospun TiO2 showed fiber-like morphology with no beads formation. This shows the successful nanofibrous morphology of TiO2 formation by electrospinning process. Figure 3b shows the morphology of pristine AC. It is clear that it has no defined shape with rough or uneven surface characteristics while Figure 3c–h reveals the presence of TiO2NF on the surface of AC, indicating that the additive was successfully introduced by co-mixing.

**Figure 3.** FESEM top view images of (**a**) TiO2NF (**b**) AC (**c**–**h**) ACTiO2NF–x, where (x = 0.5, 1.0, 1.5, 2.0, 2.5 and 5.0 wt.% TiO2NF) respectively.

#### *3.2. Structural Properties*

Figure 4a,b shows the Raman spectra and diffractogram of the TiO2NF, pristine AC and ACTiO2NF s respectively. The Raman spectrum of the pristine AC and ACTiO2NF in Figure 4a conforms to a typical graphitic carbon with distinguishable peaks at 1350 cm−<sup>1</sup> and 1590–1610 cm−<sup>1</sup> corresponding to D and G bands respectively [21]. D band arises from a defect that is based on out of plane vibration while G band relates to the ordered structure of graphite crystals [13]. For the as-synthesized TiO2NF, Figure 4a, major peaks are observed at 142, 388, 516, and 638 cm−<sup>1</sup> indicating anatase phase characteristics [8]. Distinguishable peaks of TiO2NF were observed in some of the composites especially for those with high percentage of TiO2NF as shown in Figure 4a. This indicates the successful mixing of the nanofibers with the AC.

Structural investigation was carried out to understand the crystalline nature of the materials. Figure 4b shows the diffractogram of pure TiO2NF, pristine AC, and its composites. The TiO2NF crystals are predominantly dominated with definite and sharp diffraction peaks at 2θ = 25◦, 39◦ and 43.5◦ relating to 101, 004, and 200 planes of anatase phase respectively with the presence of rutile phase at 2θ = 27.5◦, 36◦, and 41◦ relating to 110, 101, and

111 planes respectively [8]. Typical diffraction peaks of all graphite material is observed for pristine AC and ACTiO2NF at 2θ = 26◦ and 43.5 ◦ corresponding to 002 and 100 or 101 planes of graphite respectively. The sharp diffraction peaks observed at 002 planes indicates the presence of graphite microcrystalline structure in the AC [13]. In comparison with pristine AC, peaks of TiO2NF were detected at 25◦, 27.5◦, and 48◦ diffraction peaks of TiO2NF in all the composites thus showing successful doping of the nanofibers in the AC. Furthermore, using Debye-Scherrer equation, (D = Kλ/β, cosθ) where K is the constant value of 0.9, λ is the radiation of the XRD machine (0.1541 nm), β is the full width at half maximum of the diffraction peak in radian, and θ is the diffraction angle in radian, no changes were observed in the crystallite size (10 ± 0.830 nm) of the as-synthesized TiO2 nanofibers and ACTiO2NF.

**Figure 4.** (**a**) Raman spectra of Titanium oxide nanofibers TiO2NF, pristine AC, and its composites (ACTiO2NF–x) and (**b**) diffractograms of titanium oxide nanofibers (TiO2NF), pristine AC and its composites (ACTiO2NF–x where x = 0.5, 1.0, 1.5, 2.0, 2.5, and 5.0 wt.% TiO2NF).

#### *3.3. EDX and XPS Studies*

Figure 5a,b shows the EDX spectra that were obtained in order to identify the composition of the pristine AC and the composite electrode. Evidently from EDX, the elements detected at highest percentage in our materials are C and O. Titanium was detected among other elements in little quantity in the composite as shown in Figure 5b. The presence of fluorine was also detected due to the addition of PVDF (binder) added during electrode fabrication. Moreover, EDX mapping was further used to investigate the distribution of the additive in the mixture. From Figure S1a,b it can be seen that the additive distribution at lower concentration (1 wt.%) is homogenous (even dispersion) while at higher concentration (5 wt.%), its distribution is concentrated within a particular region. The results confirm the formation of well-dispersed nanofibrous within the carbon structure at low percentage concentration.

Further investigation was carried out using XPS in order to verify any change in the chemical composition of the pristine AC and its composites. Figure 5c shows two prominent peaks at 458.69 and 464.44 eV belonging to Ti 2P3/2 and Ti 2P1/2 respectively [22,23]. From Figure 5c, it is shown that Ti (IV) is present in normal state in the composites due to the observed spin orbital splitting corresponding to 5.76 eV that is obtained between Ti 2P1/2 and Ti 2P3/2. [23]. As shown in Figure 5d, no titanium was detected in the XPS spectra survey scan of the pristine AC on comparison with ACTiO2NF; a complementary result with that of EDX. The O1s peak in the composite increased a little bit when compared to pristine AC due to the influence of oxygen content of the additive. According to XPS, the

atomic composition of the ACTiO2NF-1.0 consisted of C1s 90.6 ± 0.11%, O1s 7.9 ± 0.12%, and Ti 2p 1.0 ± 33.33% while that of pristine AC is C1s 94 ± 0.84% and O1s 5 ± 3.84%.

**Figure 5.** EDX spectra of (**a**) pristine AC (**b**) ACTiO2NF–x and XPS spectra of (**c**) Ti 2p for TiO2NF (**d**) the composite pristine AC and ACTiO2NF where x = 1.0 wt.% TiO2NF.

#### *3.4. Rheology Study*

Rheology property reveals the flow nature of the slurry used under applied force. The viscosity nature of the slurry electrode was measured as a function of shear rate. Here, rheology property is in terms of dynamic viscosity, which is used to describe resistance to flow of liquid while shear rate describes the speed of deformation of the slurry under applied force. The dynamic viscosity was determined at a constant concentration of 10 wt.% carbon content in the slurry. From Figure 6a, it is obvious that the slurry follows a non-Newtonian fluid (shear thinning effect) in which the viscosity of the slurries decreases with increasing shear rate. To understand the effect of the additive, the viscosities of both pristine AC and the ACTiO2NF mixture were measured and compared as shown in Figure 6a. It is obvious from the rheogram curves that viscosity increases as the additive content increases in the composite. However, at low percentage of additive, there seems to be no significant difference in the viscosities of both the pristine AC and the ACTiO2NF mixture but at i.e., TiO2NF ≥ 2.0 wt.%, a sharp increase in viscosity was observed. The increase in viscosity can be attributed to the high specific surface area (increase frictions) of titanium nanofibers as well as its surface charge. The surface charge of the nanofibers (whether positive or negative) leads to creation of repulsive forces existing within the nanoparticles and as a result, they tend to move further apart due to strengthened force

of repulsion; consequently, viscosity increases with increasing repulsive forces [24–28]. Therefore, as the concentration of the additive in the electrodes increases, a distinguishable jump in their viscosity was observed due to increase in force of repulsion as shown in Figure 6b. This implies that the presence of additive at high content level could lead to potential clogging of the feed electrodes in the cell.

**Figure 6.** (**a**) Rheological properties of the flow-electrodes. (**b**) Schematic diagram of the ACTiO2NF mixture behavior under repulsive forces.

#### *3.5. Electrochemical Properties*

Electrochemical behavior of the pristine AC electrode and ACTiO2NF electrodes were carried out using CV at different scan rates in a potential window from −0.4 to 0.6 V vs. Ref (to ensure electrochemical stability of the electrolyte and prevent water splitting i.e., oxygen and hydrogen evolution). The experiment was conducted in 1 M NaCl aqueous solution to investigate the influence of the additive at different ratios. CV is an important technique to probe the capacitive nature of EDL [29,30].

Based on the cyclic voltammetry cycles, the EDL capacitance of the electrodes was calculated [31] and is shown in Figure 7.

**Figure 7.** Electrical double-layer capacitance of the system for AC and ACTiO2NF-x composite electrodes (ACTiO2NF–x where x = 0.5, 1.0, 2.0, 2.5, and 5.0 wt.% TiO2NF).

Thus, incorporating TiO2NF at low rate into our carbon material enhances its double layer capacitance due to the formation of a uniform network distribution of TiO2 nanofibers between AC particles. However, as the presence of TiO2NF increases in the

composite, there seem to be exhibition of poor electrochemical behavior possibly due to the increase in resistance to easy flow of ions into the pores of the electrode because of TiO2NF agglomeration [14]. Our results correlate to the findings reported in literature [8,14].

#### *3.6. Desalination Performance*

As reported in literature, the electrosorption performance of carbon-based materials is linked to their capacitive properties amid other factors [9,29]. As explained earlier, TiO2NF was added at different wt.% to influence the capacitance of commercial AC and this factor was verified through the electrochemical characteristics of the electrodes. Furthermore, in order to confirm the link between electrosorption performances and capacitive properties, desalination experiments were carried out under a semi-continuous system in a flow channel. Desalination was conducted at an operational cell potential ΔE = 1.2 V for 30 min using 5000 mg·L−<sup>1</sup> NaCl as the feed solution. Difference in conductivity was monitored and recorded during the experiment. FCDI performance indicators such as desalination efficiency (DE) in Figure 8a, salt removal rate (SRR), and charge efficiency (CE) in Figure 8b were used to evaluate the performance of the flow electrodes.

**Figure 8.** (**a**) Desalination efficiency; (**b**) salt removal rate and charge efficiency of AC and ACTiO2NF−x, where (x = 0.5, 1.0, 1.5, 2.0, 2.5, and 5.0 wt.% TiO2NF) respectively.

Important increase in DE is noticed in Figure 8a. Notably, the ACTiO2NF electrodes show higher DE than pristine AC and among all the electrodes, ACTiO2NF-1.0 exhibited the best DE, which implies fast ion mobility to the pores of the flow electrode thus leading to quick salt removal; a consequent effect on CE and SRR.

Low desalination behavior of ACTiO2NF-x (x ≥ 2.0 wt.% TiO2NF) at high percentage could be due to the fact that at this ratio, the nanofibers tend to agglomerate among themselves (not well dispersed or less uniform) and as such, making the pores of the electrode (surface hindrance) not easily accessible for ions adsorption. In addition, electrode passivation is likely to occur due to the high presence of the nanofibers [32]. This will make the electrode surface not easily permeable for ions (impermeable layer formation on electrode surface) thus significantly affecting the electrosorption performance of the electrodes as observed in our materials, (whitish layers on the surface of flow electrodes). It has to be mentioned that only a small fraction of titanium nanofiber (≤1% weight) is sufficient to significantly increase FCDI performances without increasing much viscosity.

#### **4. Conclusions**

To summarize, we used an electrospinning procedure to create titanium nanofibers. Without any further post-treatment, the as-spun nanofibers were rationally mixed with AC using a simple agitating procedure to generate hybrid composites of ACTiO2NF-x (x = 0.5, 1.0, 1.5, 2.0, 2.5, and 5.0 wt.% ACTiO2NF). The composites were subsequently described and tested for the first time as electrodes in FCDI. Introduction of TiO2 nanofibers into AC improved the electrochemical characteristics of the material. The addition of TiO2 nanofibers to the composite electrodes increased their performance, with ACTiO2NF-1.0 demonstrating distinguishing and remarkable properties with the best desalination behavior. The composites' nanofibrous shape allows for greater anchoring within the AC network, allowing for improved ion transport and migration to the pores. Finally, the technology described has the potential to produce carbon-based flow electrodes with better shape and desalination performance in the FCDI methodology.

**Supplementary Materials:** The following are available online at https://www.mdpi.com/article/10 .3390/ma14226891/s1. Figure S1: EDX mapping of (a) ACTiO2NF-5.0 and (b) ACTiO2NF-1.0.

**Author Contributions:** Conceptualization, G.F.; methodology, G.F. and M.B.; software, G.F. and M.T.; validation, M.B., F.Z., P.S. and M.C.; investigation, G.F.; resources, F.Z., M.B., P.S. and M.C.; data curation, G.F. and M.T.; writing—original draft preparation, G.F. and M.T.; writing—review and editing, G.F., M.T., F.Z., M.B. and M.C.; visualization, G.F., M.T., F.Z. and M.B.; supervision, M.B., F.Z., P.S. and M.C.; project administration, M.B., F.Z. and M.C. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by "Axe H2O" and "Axe Energie" in European Institute of Membranes (IEM). And the Federal Government of Nigeria through Tertiary Education Trust fund (TETFUND) and Campus France (CF) for the Ph.D. funding of Gbenro Folaranmi with the funding number CAMPUS FRANCE 914886H.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding authors.

**Acknowledgments:** Special thanks to Fida Tanos (IEM) for guidance in titanium nanofibers synthesis and Mahmoud Abid (IEM) for furnace operation.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


### *Review* **BiFeO3-Based Relaxor Ferroelectrics for Energy Storage: Progress and Prospects**

**Bipul Deka 1,2,3,\* and Kyung-Hoon Cho 1,2,\***


**\*** Correspondence: bipul.deka@kumoh.ac.kr (B.D.); khcho@kumoh.ac.kr (K.-H.C.)

**Abstract:** Dielectric capacitors have been widely studied because their electrostatic storage capacity is enormous, and they can deliver the stored energy in a very short time. Relaxor ferroelectricsbased dielectric capacitors have gained tremendous importance for the efficient storage of electrical energy. Relaxor ferroelectrics possess low dielectric loss, low remanent polarization, high saturation polarization, and high breakdown strength, which are the main parameters for energy storage. This article focuses on a timely review of the energy storage performance of BiFeO3-based relaxor ferroelectrics in bulk ceramics, multilayers, and thin film forms. The article begins with a general introduction to various energy storage systems and the need for dielectric capacitors as energy storage devices. This is followed by a brief discussion on the mechanism of energy storage in capacitors, ferroelectrics, anti-ferroelectrics, and relaxor ferroelectrics as potential candidates for energy storage. The remainder of this article is devoted to reviewing the energy storage performance of bulk ceramics, multilayers, and thin films of BiFeO3-based relaxor ferroelectrics, along with a discussion of strategies to address some of the issues associated with their application as energy storage systems.

**Keywords:** energy storage; BiFeO3; relaxor ferroelectrics; domain engineering; polymorphic nanodomain

#### **1. Introduction**

Global warming poses potential threats to the planet Earth's future. The continuous burning of fossil fuels has increased the concentration of CO2 and other greenhouse gases in the Earth's atmosphere, leading to a warmer atmosphere and climate change. In addition, the depletion of fossil fuel resources, the dominant candidate in the energy market, poses the risk of an energy crisis in a world where the number of consumers is increasing day by day. Considering the serious threats to the lives on Earth and the risk of an energy crisis posed by the use of fossil fuel resources, the transition to clean energy is a serious consideration. However, renewable energy sources such as solar, wind, tides, and geothermal energy are intermittent by nature. Therefore, harnessing and storing renewable energy for future access is a challenging task.

Electrical energy harvested from renewable sources offers enormous opportunities for meeting future energy demands and the feasibility of the transition to clean energy. However, the usefulness of the electrical energy generated depends on its efficient storage, which is necessary for around-the-clock use. An efficient electrical energy storage (EES) system is the heart of the commercial and residential grid-based utilization of electrical energy. Therefore, the development of advanced EES systems is critically important for meeting the growing energy demands and effectively leveling the cyclic nature of such energy sources [1]. For over 200 years, batteries have been widely used in EES systems and are still being widely used. Solid oxide fuel cells (SOFCs), electrochemical capacitors (ECs), superconducting magnetic energy storage (SMES) systems, flywheels, and electrostatic capacitors (dielectric capacitors) are common current energy storage technologies [2].

**Citation:** Deka, B.; Cho, K.-H. BiFeO3-Based Relaxor Ferroelectrics for Energy Storage: Progress and Prospects. *Materials* **2021**, *14*, 7188. https://doi.org/10.3390/ ma14237188

Academic Editors: Marc Cretin, Sophie Tingry and Zhenghua Tang

Received: 14 October 2021 Accepted: 23 November 2021 Published: 25 November 2021

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**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

A perfect energy storage device is characterized by high energy and power densities. A comparison of the storage efficiency of the technologically relevant candidates for EES systems can be realized from the Ragone plot shown in Figure 1, which displays the status of EES systems according to their energy and power densities. As can be seen, SOFCs and batteries exhibit a high energy density with low power density, while dielectric capacitors exhibit the opposite behavior, that is, high power density and low energy density. The ECs, SMES, and flywheel have medium power and energy densities. In addition to the high power and energy density, the charge/discharge rate is a deciding factor for EES systems. The energy storage and delivery in SOFCs/batteries are based on the chemical reaction process and may take 1–100 h of time. Dielectric capacitors typically exhibit fast charge/discharge rates, between μs and ms, whereas those for ECs, flywheels, and SMES are between 1 s and 1 h. The fast charge/discharge rate of dielectric capacitors is associated with the separation of comparatively fewer heavy bound charges under the influence of the electric field to be stored. Thus, to summarize, no individual EES candidate possesses both high power and energy densities simultaneously. Therefore, the technological relevance of each candidate as an EES may be determined by the final requirements. However, among the various potential candidates for EES, dielectric capacitors have the advantage of withstanding high-voltage and large-scale applications because of their lower cost [3,4].

**Figure 1.** Ragone plot of various energy storage devices: electrostatic capacitors, electrochemical capacitors, SMES, flywheels, batteries, and SOFCs. The straight dashed lines and the associated times correspond to the characteristic times. Reused with permission from [2]. © 2021 Elsevier Ltd.

Dielectric capacitors are one of the key components in modern communication technology, with applications in electronic circuits, warfare, distributed power systems, hybrid electric vehicles, clean energy storage, high-power applications, etc. [5–11]. Among the various EES systems, dielectric capacitors exhibit the fastest discharge speed; therefore, they can generate intense pulse power [11–14]. The fast discharge speed and high fatigue resistance of dielectric capacitors enable their potential applications in various electronic systems. This includes medical equipment such as defibrillators, pacemakers, surgical lasers, and X- ray units; scientific research equipment such as high-power accelerators and high-intensity magnetic fields; commercial system devices such as camera flash, underground oil and gas exploration, avionics, transportation (hybrid cars, space shuttle power systems), transversely excited atmospheric lasers, and advanced electromagnetic systems [15].

The current review discusses the recent progress on the development of high-energy storage dielectric capacitors based on the relaxor ferroelectric (RFE) of BiFeO3. The two important figures of a capacitor that determine its energy storage performance are the recoverable energy density (*U*rec) and energy efficiency (*η*), which depend on the saturation polarization (*P*max), remnant polarization (*P*r), and breakdown strength (BDS) of the materials. Linear dielectric (LD), ferroelectric (FE), and anti-ferroelectric (AFE) materials are widely used for the fabrication of ceramic capacitors. Although the LDs possess excellent values of BDS, their *U*rec values are quite low due to weak polarization. FE materials, on the other hand, possess quite a large *P*max and *P*r. As can be seen in Section 2, the large values of both *P*max and *P*r result in a low *U*rec value. Moreover, the ferroelectric materials suffer large hysteresis loss, which has a significant detrimental effect on the energy efficiency of the capacitors. AFE materials behave like LDs in the low field regime, undergoing a field-induced FE state, yielding a high *P*max at a high electric field with large hysteresis loss. Therefore, the issues associated with the AFE for their applications in the high-energy storage application are essentially similar to that of LDs at low electric fields and to that of FEs at high electric fields. However, the problem of hysteresis loss inherent to FE and AFE materials is found to be minimized in RFEs while maintaining a significantly large value of *P*max. This motivates the scientific community to turn their heads towards the RFEs in search of high-energy storage capacitors.

Ceramic dielectric capacitors based on BiFeO3 have recently gained interest in the field of energy storage applications because of the high polarization (~90 μC cm−2) predicted in BiFeO3, along with its high ferroelectric Curie temperature (*T*C) (~830 ◦C) [16]. The advantage of having a high *T*<sup>C</sup> is that the materials do not lose their ferroelectric (FE) nature at such high temperatures, which is essential for applications in the high-temperature regime. Temperature stability is an important issue that needs to be addressed while designing a capacitor for operation in the high-temperature regime. For example, ceramicpolymer composites have excellent storage performance; however, their performance degrades very rapidly as the temperature approaches 100 ◦C [15]. Pb-based ferroelectrics often have the disadvantage of adverse harmful effects on the environment and human beings. Pb-oxide, which is a main component of Pb-based ferroelectrics, is highly toxic and volatile at high temperatures, causing environmental pollution during the fabrication process. Disposal and recycling of Pb-based materials and devices at an industrial scale also creates atmospheric problems due to the difficulty of Pb removal. Exposure to the heavy metal Pb causes detrimental effects such as kidney and brain damages, and chronic exposure may lead to damages to the central nervous system and affect blood pressure, vitamin D metabolism, etc. Young children are more vulnerable to Pb exposure, as the absorption of Pb in children's bodies is 4–5 times higher than in an adult body. Therefore, to reduce the use of hazardous materials such as Pb, various countries have adopted different restrictions on hazardous materials [17,18]. Although there are a few reviews on dielectrics for energy storage in general, to the best of our knowledge, there has been no such review for BiFeO3-based relaxor ferroelectrics. Here, we present a review of the recent progress on BiFeO3-based relaxor ferroelectric for energy storage, discussing various issues to meet practical applications. We first discuss the fundamentals of energy storage in dielectrics and the pros and cons of various nonlinear dielectrics with respect to their applications in energy storage. We then discuss the characteristics of relaxor ferroelectrics and their importance in energy storage, followed by a brief discussion of the basic properties of BiFeO3. Following this, we present the recent progress in energy storage studies on BiFeO3 and strategies for further enhancement.

#### **2. Fundamentals of the Energy Storage Mechanism in Dielectrics**

The energy storage mechanism of dielectrics is based on their polarization under the application of an electric field. A dielectric under an applied electric field is polarized such that equal amounts of positive and negative charges accumulate at the surfaces of the dielectrics. In other words, an electric field opposite to the applied field is induced inside the dielectric. The strength of the induced field grows exponentially with time until its magnitude is equal to that of the external field. This process is known as charging the capacitor. Thus, the induced electrostatic energy is stored in the dielectric and can be used for application upon discharge through a load. The amount of stored energy (*U*) can be obtained from the potential difference (*V*) across the dielectrics and the charge (*q*) induced at the electrode on the surface of the dielectrics using the following equation:

$$
\mathcal{U}I = \int\_0^{q\_{\text{max}}} Vdq\tag{1}
$$

where *qmax* is the maximum amount of charge accumulated at the electrode when the capacitor is fully charged, and *dq* is the increment of charge during charging. A figureof-merit (FOM), which signifies the energy storage performance of a capacitor, is represented in terms of energy storage density (*U*st), defined as the energy stored per unit volume. Mathematically,

$$dL\_{st} = \frac{\int\_0^{q\_{\max}} Vdq}{Ad} = \int\_0^{D\_{\max}} EdD \tag{2}$$

where *A* is the electrode area of the capacitor, *d* is the distance between the electrodes (thickness of the dielectric layer), and *D* is the electric displacement of the capacitor. For a weak electric field, *D* is related to the external electric field (*E*) and polarization (*P*) as follows:

$$D = P + \varepsilon\_0 E \tag{3}$$

where *ε*<sup>0</sup> represents the permittivity of a vacuum. Materials obeying Equation (3) are classified as linear dielectrics. For a linear dielectric, *P* is assumed to be a linear function of *E*:

$$P = \varepsilon\_0 \chi E \tag{4}$$

where the quantity *χ* is termed the linear dielectric susceptibility. At a high electric field, it is necessary to consider the nonlinear contribution of susceptibility, and Equation (4) takes the most general form as:

$$P = \varepsilon\_0 \left( \chi E + \chi^{(2)} E^2 + \chi^{(3)} E^3 + \dots \right) \tag{5}$$

where *χ*(2) and *χ*(3) are higher-order susceptibilities, giving rise to nonlinear effects. Using linear approximation, the stored energy density of a dielectric material with a high dielectric constant (*D* ≈ *P*) can be calculated as follows:

$$
\delta L\_{\rm st} = \int\_0^{p\_{\rm max}} EdP \tag{6}
$$

Equation (6) indicates that the electric polarization as a function of the electric field should be measured to calculate *Ust*. In other words, it is necessary to measure the polarization-electric field (*P*−*E*) hysteresis loop to obtain the stored energy density, as shown in Figure 2. Therefore, the shape and size of the *P*−*E* loop and the nature of the dipole/domain structures determine the energy storage performance of dielectric materials. However, the dynamics of the polarization vector, growth of domains, and domain wall movements for *E* = 0 → *E*max and *E* = *E*max → 0 directions in the *P*−*E* measurement protocol are different from each other. This leads to a non-zero value of polarization, even

at *E* = 0, known as remanent polarization (*P*r). As a result, a part of the stored energy is lost, which appears as the hysteresis of the *P*−*E* loop. In other words, it is impossible to recover the stored energy density to its fullest amount when it is discharged. The loss part of the stored energy or energy loss density (*U*loss) is given by the area of the loop. The recoverable energy density is calculated as follows:

$$dJ\_{\rm rec} = \int\_{P\_r}^{P\_{\rm max}} EdP \tag{7}$$

Another FOM signifying the energy storage performance is the efficiency (*η*), which represents the amount of stored energy density available for use as recoverable energy density. It is defined as the ratio of the recoverable energy density to the total stored energy density:

$$\eta = \frac{lL\_{\rm rec}}{lI\_{\rm st}} \times 100\% = \frac{lI\_{\rm rec}}{lI\_{\rm rec} + lI\_{\rm loss}} \times 100\% \tag{8}$$

Equation (7) suggests that a combination of high *P*max, low *P*r, and high breakdown strength are necessary to obtain a high *U*rec value. In addition, Equation (8) requires a dielectric with low hysteresis loss to obtain a large efficiency value. Therefore, Equations (7) and (8) are often considered the governing equations for designing dielectric materials for high-performance energy storage. However, a dielectric with high *ε* usually features high dielectric loss, leading to heat generation during electric field cycling and the possibility of thermal breakdown during operation.

Typical *P*–*E* loops of LDs, FEs, AFEs, and RFEs are shown in Figure 2. LDs are characterized by very low values of polarization and a high BDS. Some of the widely studied LDs are CaTiO3 [19,20], SrTiO3 [21], and CaTiO3-CaHfO3 [22]. Because of the low value of polarization, the recoverable energy density of LD is quite low. Therefore, LDs are not suitable for application in the field of high-energy storage application. Over the years, FE and anti-ferroelectric (AFE) materials have been extensively studied for application in energy storage systems, and efforts to enhance their performance have surged. FE materials exhibit spontaneous polarization, a large value of *P*max, and a coercive field (*E*c). A typical *P*–*E* loop of FE materials is shown in Figure 2b. On the microscopic scale, FEs are composed of a large number of domains separated by domain walls. The dipoles in a domain are oriented in the same direction, and the directions of the domain polarizations can be switched by applying an electric field. However, the energy loss density is quite high in the FEs because of their high coercivity. Moreover, the *P*r and *P*max values have the same order of magnitude, resulting in a very small value of *P*max−*P*r. Therefore, the *U*rec and *η* of the FEs are not promising. Because of this, single-phase FEs have not gained much interest in energy storage devices. Among the various FEs, representative compositions studied for energy storage are based on (Bi,Na)TiO3 [23–26], Ba(Zr,Ti)O3 [27–29], BaTiO3 [30,31], and (K,Na)NbO3 [32]. Unlike FEs, AFE materials lack a net polarization because of the anti-parallel alignment of the spontaneous polarization vectors in their domain. The typical *P*−*E* loops of AFE materials are shown in Figure 2c. The electric dipoles align anti-parallel to each other in the AFE domain, as shown in the inset of Figure 2c. At a low electric field, the polarization of the AFE materials varies linearly with the applied field. At a sufficiently high electric field, the electric dipoles in a domain rotate to align in the parallel direction, and the AFE behaves similarly to an FE with a further increase in the field strength. This is known as the field-induced AFE-FE transition. Once the electric field is removed, the induced FE phase reverts to the AFE state, thereby producing double hysteresis in the *P*−*E* loop. The high electric field for the AFE-FE phase transition (*E*AFE−FE) coupled with the high *P*max and low *P*<sup>r</sup> indicates the possibility of achieving high storage capacity in AFE materials. The most intensively studied AFE systems are based on (Pb, Zr)O3 [33–36], (Bi, Na)TiO3 [37–39], and AgNbO3 [40–42]. However, the *E*AFE−FE for some AFEs is higher than their BDS at room temperature, signifying a breakdown before the transition to the highly polarized FE phase. Moreover, AFE materials cannot withstand

large charge-discharge cycles, which is an important aspect in practical operations because such cycling leads the materials to undergo several alternate AFE-FE transitions, leading to physical cracks [43]. Moreover, the high-field FE phase often suffers from severe energy loss, which is particularly observed in AgNbO3-based AFEs [44].

**Figure 2.** Typical *P*–*E* loops for (**a**) LD (**b**) FE, (**c**)AFE, and (**d**) RFE. Insets are schematics of domains with the alignment of polarization vectors (arrowheads). In RFEs, polar nanoregions (shown in circular patches) are sparingly distributed in a non-ferroelectric matrix (green area in the inset).

#### **3. Relaxor Ferroelectrics**

Relaxor ferroelectrics (RFEs) are an important class of materials that have attracted significant interest in energy storage applications. RFEs exhibit nanosized polar regions embedded in a nonpolar matrix. The polar nanoregions (PNRs) exhibit spontaneous polarization; however, the inter-PNR interaction is very weak [45]. The typical size of a PNR is 2–10 nm. PNRs are highly dynamic and sensitive to external stimuli. Because of the lack of inter-PNR interactions, PNRs under an electric field evolve independently of nearby PNRs. Therefore, the polarization state can return to its initial state after the electric field is removed. The *P*−*E* loop for a typical RFE is shown in Figure 2d, which features *P*r ~ 0, a considerably high *P*max, and a small hysteresis loop.

RFEs feature (i) a broad maximum in *ε* around *T*<sup>m</sup> (maximum temperature in the *ε*-*T* curve); (ii) strong frequency dispersion of *ε* and loss tangent (tan *δ*) peaks, i.e., shifting of the peaks toward higher temperatures while measured at lower to higher frequencies; and (iii) low *P*r [46,47]. Therefore, RFEs are FE materials that simultaneously exhibit dielectric relaxation and ferroelectricity. However, unlike normal FEs, where the paraelectric-FE phase transition can be explained by Curie's law, the temperature dependence of *χ* in RFEs in the paraelectric phase obeys the following:

$$\chi = \frac{\mathbb{C}}{(T - T\_{\varepsilon})^{\gamma}} \tag{9}$$

where the parameter *γ* (1 < *γ* < 2) represents the broadness of the dielectric peak. For a normal FE, *γ* = 1. Several models have been proposed to explain the peculiar characteristics of RFEs, such as the diffuse phase transition model [48], super paraelectric model [49], dipolar glass model [50], random-field model [51], random-site model [52], bi-relaxation model [53], and spherical random-bond-random-field model [54]. However, the underlying mechanism of RFEs is yet to be clearly understood.

#### **4. Energy Storage Performance of BiFeO3-Based Relaxor Ferroelectrics**

BiFeO3 exhibits a distorted perovskite structure, as shown in Figure 3a. It possesses a rhombohedral structure (point group: R3c) at room temperature with an *a¯a¯a¯* tilt system, in which the neighboring oxygen octahedra rotate anti-clockwise about the [111] direction [55,56]. The rhombohedral unit cell is described with lattice constants *a* = *b* = *c* = 3.965 Å and *α* = *β* = *γ* = 89.3–89.4◦ [57]. There are two formula units of BiFeO3 in the rhombohedral cell, with three atoms in its asymmetric unit occupying Wyckoff positions: 6*a* (Bi3+ and Fe3+) and 18*b* (O2−) [58]. in a hexagonal frame of reference with the hexagonal *c*-axis parallel to the diagonals of the cubic perovskite with the lattice constants *a*hex = 5.58 and *c* = 13.90 Å [57].

The *R3c* symmetry permits long-range FE order in BiFeO3 along the threefold axis [111]. Various experiments have confirmed the ferroelectricity in BiFeO3 below *T*<sup>C</sup> = 1143 K [55,59]. The constituent atoms Bi, Fe, and O are displaced from their centrosymmetric positions along the threefold axis, and Bi ions have the largest displacement with respect to O ions [55]. The lone-pair-active Bi ions in BiFeO3 are displaced to a large extent in comparison with other FE compounds with non-lone-pair-active cations. Therefore, a large value of spontaneous polarization, on the order of 90 μC cm−2, has been predicted in BiFeO3 from ab initio calculations [16,60]. However, a polarization value close to the calculated values could not be obtained until recently [61], after a series of initial experimental failures to achieve spontaneous polarization of BiFeO3, as predicted by theory [62–67]. Lebeugle et al. [61] measured a very large saturated polarization (approximately 60 μC cm<sup>−</sup>2) in a high-quality single-crystal BiFeO3 at room temperature, as shown in Figure 3b.

**Figure 3.** (**a**) Structure of R3c BiFeO3. (**b**) P-E loop of BiFeO3 bulk single crystal. Figures reproduced with permission from [55] and [61], respectively. © 2021 American Physical Society (**a**) and 2007 American Institute of Physics (**b**).

BiFeO3 possesses the highest values of spontaneous polarization and *T*<sup>C</sup> among Pbfree FEs. As discussed in Section 2, high *P*max and large *P*max –*P*<sup>r</sup> are among the most important factors for obtaining a high storage capacity. Therefore, BiFeO3 with a large *P*<sup>r</sup> in its naturally occurring FE phase is not suitable for energy storage applications, which is a drawback of all FE materials in general. One way to obtain a small *P*<sup>r</sup> is to break the long-range FE order such that it becomes an RFE. Many researchers have reported this method to create PNRs embedded in a non-FE matrix and obtain a significant reduction in the *P*<sup>r</sup> value. In addition, there are always issues pertaining to the leakage current in pure BiFeO3, which eventually limits its high breakdown strength. Various successful methods to enhance the resistivity in phase-pure BiFeO3, such as doping at the A-site and the addition of Mn, have been discussed in the literature. In the following sections, considering these issues, the energy storage performance of BiFeO3-based materials, with special emphasis on the RFEs, are reviewed.

#### *4.1. BiFeO3-Based Binary System*

BiFeO3−BaTiO3 solid solutions have been widely investigated as promising candidates in the field of ceramic dielectrics-based energy storage materials. A remarkable feature of BiFeO3-based solid solutions is the morphotropic phase boundary (MPB), where the solid solution displays a composition-driven structural transition in its phase diagram. The crystal structure changes abruptly across the MPB, and various physical properties, such as piezoelectric coefficients and polarization, are maximal at the MPB. The MPB of the BiFeO3−BaTiO3 system is shown in the phase diagram in Figure 4, where the rhombohedral and tetragonal phases coexist in the MPB region [68–70]. Careful optimization of the BaTiO3 content produces excellent ferroelectric and piezoelectric properties in MPB compositions.

**Figure 4.** Phase diagram of (1-x)BiFeO3-xBaTiO3 solid solution. Reused with permission from [68]. © 2021 Willey-VCH Verlag GmbH & Co.

However, BiFeO3-BaTiO3 also possesses a high *P*<sup>r</sup> [71] and high dielectric loss, which are detrimental for energy storage. Previous studies have shown that doping small amounts of La2O3, MnO2, and Nb2O5 in BiFeO3 can significantly enhance the electrical resistivity and energy loss density [72,73], which is beneficial for energy storage. Wang et al. [74] reported a large enhancement of resistivity in Nb2O5-modified BiFeO3- BaTiO3, i.e., (1-x)(0.65BiFeO3-0.35BaTiO3)–xNb2O5 (x = 0, 1, 3, 5 mol%) by several orders (~1010–1014 Ω cm) compared with the undoped BiFeO3-BaTiO3 (~108 Ω cm). The compounds with x = 0.01 and 0.03 exhibited slimmer *P*−*E* loops, similar to RFEs, with x = 0.03 exhibiting the highest *P*max (25.21 μC cm<sup>−</sup>2) and lowest *P*<sup>r</sup> (5.53 μC cm<sup>−</sup>2). They obtained a maximum *U*rec of 0.71 J cm−<sup>3</sup> at *E* = 90 kV cm-1. Zhu et al. [75] found a significant improvement in the BDS of 0.52BiFeO3-0.48BaTiO3 ceramic up to 130–140 kV cm−<sup>1</sup> by adding La2O3 and MnO2. Under this condition, they obtained *U*rec = 1.22 J cm−<sup>3</sup> with

*η* = 58% for 0.52Bi0.98La0.02FeO3-0.48BaTiO3 + 0.3 wt.% MnO2 compound, whereas undoped 0.52BiFeO3-0.48BaTiO3 exhibited *U*rec = 1.08 J cm−<sup>3</sup> with *η* = 49%. They found that the addition of La2O3 and MnO2 increased the amount of the FE phase, reduced the grain size, and facilitated densification, which helped to induce large Δ*P* (*P*max − *P*r) as well as BDS compared with undoped compounds. This phenomenon was found to be more pronounced when Nd was substituted for Bi sites in MnO2-added BiFeO3-BaTiO3 solid solution, as shown in Figure 5. Wang et al. [76] synthesized highly dense (relative density *ρ*<sup>r</sup> = 95% to 97.6%) 0.75BiFeO3-0.25BaTiO3 ceramics by Nd substitution with 0.1 wt.% MnO2 addition that could endure a high electric field up to 180 kV cm<sup>−</sup>1, as shown in Figure 5. The solid solution with 15 mol% Nd content featured the highest value of *U*st = 4.1 J/cm<sup>3</sup> and *U*rec = 1.82 J cm<sup>−</sup>3; however, it had a low value of *η* = 41.3%. Recently, Chen et al. [77] successfully synthesized highly dense Sm-doped BiFeO3-BaTiO3 binary ceramics that can endure a very high electric field of up to 200 kV cm−1. Sm substitution significantly reduced the grain size and enhanced the density, which is believed to be the reason for the high BDS. The binary solid solutions exhibit excellent *U*st and *U*rec; however, their low efficiency limits practical device applications.

**Figure 5.** Unipolar P-E loops for 0.75Bi1-xNdxFeO3-0.25BaTiO3 + 0.1 wt% MnO2 system: (**a**) x = 0.15, (**b**) x = 0.20, (**c**) x = 0.30, and (**d**) x = 0.40. Reproduced with permission from [76]. © 2021 The Royal Society of Chemistry.

### *4.2. BiFeO3-Based Ternary System*

#### 4.2.1. Bulk Ceramics

The addition of a third perovskite oxide to binary BiFeO3-MTiO3 (M = Ba and/or Sr) as the end member of the ternary system has been found to be very promising for inducing the relaxor phase and enhancing the energy storage performance [78–84]. Zheng et al. [78] reported the successful induction of the relaxor phase in BiFeO3-BaTiO3 as a result of the substitution of BaMg1/3Nb2/3O3. The relaxor 0.61BiFeO3-0.33BaTiO3-0.06BaMg1/3Nb2/3O3 exhibited a *U*rec of 1.56 J cm−<sup>3</sup> at *E* = 125 kV cm−1, with *η* ~ 75%. This compound also exhibited good temperature stability for energy storage and efficiency in the temperature

range of 25 ◦C to 190 ◦C. In a similar effort, Zheng et al. [79] reported an improved storage performance with a *U*rec of 1.66 J cm−<sup>3</sup> at 130 kV cm−<sup>1</sup> and *η* ~ 82% in highly dense (*ρ*<sup>r</sup> > 97%) 0.61BiFeO3-0.33BaTiO3-0.06LaMg1/2Ti1/2O3 ceramics. Meanwhile, Liu et al. [80] reported enhancement of relaxor characteristics in terms of broader peaks of dielectric permittivity (Figure 6) and significant enhancement of energy storage performance in (0.66 − x)BiFeO3-0.34BaTiO3-xBaZn1/3Ta2/3O3 for x > 0. They reported slim *P*–*E* loops for x > 0 with the highest BDS of *E* = 160 kV cm−<sup>1</sup> and a high *U*rec of 2.56 J cm−<sup>3</sup> for the x = 0.06 composition (Figure 7). Tang et al. [84] reported a BDS of 180 kV cm−<sup>1</sup> and *U*rec = 1.62 J cm−<sup>3</sup> for 0.85(0.65BiFeO3-0.35BaTiO3)-0.15Ba(Zn1/3Nb2/3)O3 bulk ceramics. Sun et al. [85] reported similar values of *U*rec = 2.11 J cm−<sup>3</sup> at *E* = 195 kV cm−<sup>1</sup> with *η* = 84% in a highly dense 0.56BiFeO3-0.30BaTiO3-0.14AgNbO3+5 mol% CuO system prepared by a modified thermal quenching technique.

**Figure 6.** Induction of relaxor phase characterized by diffused phase transition in (0.66-x)BiFeO3- 0.34BaTiO3-xBaZn1/3Ta2/3O3 systems. Reproduced with permission from [80]. © 2021 The Royal Society of Chemistry.

Yu et al. [86] studied the effect of the microstructure on the energy storage performance of a BiFeO3-BaTiO3-Bi(Mg2/3Nb1/3)O3 system. They prepared coarse-grained (grain size ~2 to 4 μm) and fine-grained (~0.55 to 0.9 μm) microstructures using planetary ball milling and high-energy ball milling processes, respectively. BiFeO3-BaTiO3-Bi(Mg2/3Nb1/3)O3 solid solutions with a fine-grained microstructure exhibited higher Δ*P* (~ 30 μC cm<sup>−</sup>2) and BDS (~ 110 kV cm−1) than the coarse-grained samples (Δ*P* ~ 10 μC cm−2, BDS ~ 50 kV cm<sup>−</sup>1). Under such drastic microstructural evolution, they reported a higher *U*rec of 1.26 J cm−<sup>3</sup> in fine-grained samples, compared with *U*rec = 0.16 J cm−<sup>3</sup> for coarse-grained samples. Yang et al. [87] showed that utilizing a liquid-phase sintering mechanism can significantly enhance the BDS while maintaining the relaxor characteristics and high dielectric permittivity. Using 2 wt% BaCu(B2O5) (BCB) as the low melting point additive in 0.1 wt% MnO2-added (0.67−x)BiFeO3–0.33(Ba0.8Sr0.2)TiO3-xLa(Mg2/3Nb1/3)O3 solid solution, they achieved a BDS of 230 kV cm−<sup>1</sup> for x = 0.06. With the use of such a high field, the compound exhibited *U*rec = 3.38 J cm−<sup>3</sup> with *η* = 59%. BCB formed large amounts of liquid phase at the grain boundaries during sintering and significantly reduced the average grain size down to the submicron range, as shown in Figure 8, by impeding grain growth at lower sintering temperatures. The high density of grain boundaries in the microstructure of the submicron grain size offered high electrical resistivity, resulting in enhanced BDS (Figure 8). Moreover,

low-temperature sintering with the addition of MnO2 helped to decrease Fe3+ ↔ Fe2+ valence fluctuations by minimizing Bi2O3 loss during synthesis, which was also critical for enhancing the BDS. This compound exhibited good temperature stability, with *U*rec = 1.15–1.27 J cm−<sup>3</sup> in the temperature range of 30 ◦C to 170 ◦C. Using the PVA-assisted viscous polymer process (VPP) route, Liu et al. [88] obtained high *ρ*<sup>r</sup> ~ 99% 15 mol% Sr0.7Bi0.3FeO3–modified 0.85(0.65BiFeO3-0.35BaTiO3) system with a fine grain microstructure. The ultra-high *ρ*<sup>r</sup> and uniform submicron grains significantly enhanced the BDS, with a value of 330 kV cm−1, compared with that of the sample prepared without VPP (180 kV cm<sup>−</sup>1). Under this condition, they obtain an ultra-high *U*rec of 4.95 J cm−<sup>3</sup> with *η* ~ 73%. The calculated *U*rec from the charge-discharge cycling test was 2.36 J cm−<sup>3</sup> at 300 kV cm<sup>−</sup>1. Likewise, Sm doping has been found to be very effective in reducing the grain size and increasing the density of sintered ceramics. Chen et al. [78] reported a significant decrease in the average grain size in the 0.67Bi1-xSmxFeO3-0.33BaTiO3 system with an increase in x, resulting in a BDS of 200 kV cm−<sup>1</sup> at x = 0.1. This resulted in an *U*rec = 2.8 J cm−<sup>3</sup> for x = 0.1; however, it had low efficiency, *η* = 55.8%. In another report, Li et al. [89] reported an ultra-high *U*rec = 3.2 J cm−<sup>3</sup> at *E* = 206 kV cm−<sup>1</sup> with high efficiency, *η* = 92%, in a highly dense (*ρ*<sup>r</sup> ~ 98%) solid solution of (1−x)Bi0.83Sm0.17Fe0.95Sc0.05O3-x(0.85BaTiO3- 0.15Bi(Mg0.5Zr0.5)O3) with x = 0.75. A similar *<sup>U</sup>*rec ≈ 3.06 J cm−<sup>3</sup> at *<sup>E</sup>* = 167 kV cm−<sup>1</sup> with *η* ≈ 92% associated with the improvement of BDS has been reported in (1−x)BiFeO3 x(0.85BaTiO3-0.15Bi (Sn0.5Zn0.5)O3) with x = 0.65 [90].

**Figure 7.** BaZn1/3Ta2/3O3 modification to BiFeO3-BaTiO3 leads to slimmer *P*–*E* loops. Unipolar *P*–*E* loops for (0.66-x)BiFeO3-0.34BaTiO3-xBaZn1/3Ta2/3O3 with (**a**) x = 0.04, (**b**) x = 0.08, and (**c**) x = 0.10. (**d**) The optimum recoverable energy density and efficiencies. Reproduced with permission from [80]. © 2021 The Royal Society of Chemistry.

**Figure 8.** SEM images of 0.61BiFeO3-0.33(Ba0.8Sr0.2)TiO3-0.06La(Mg2/3Nb1/3)O3 + y wt.% MnO2 + z wt.% BaCu(B2O5) solid solution sintered at their optimum temperatures: (**a**) y = z = 0, (**b**) y = 0.1 and z = 1, (**c**) y = 0.1 and z = 2, (**d**) y = 0.1 and z = 3, and (**e**) y = 0.1 and z = 4. (**f**) Variations of BDS, grain size, and relative density with y + z. Reused with permission from [87]. © 2021 Elsevier Ltd.

Designing a specific microstructure has been found to be highly effective for enhancing energy storage performance. Microstructure strongly influences the BDS of dielectric materials and their relaxor characteristics. Wang et al. [81] designed core-shell microstructure (Figure 9a) with BaTiO3-rich shells and BiFeO3-rich cores in (0.7-x)BiFeO3-0.30BaTiO3 xBi(Zn2/3Nb1/3)O3+0.1wt%Mn2O3 ceramics that could withstand an electric field as high as 190 kV cm−1. They found that the shells and cores in the solid solutions were different in structure and electrical characteristics. The shell exhibited a pseudo-cubic structure, which is paraelectric in nature, whereas the core parts had a ferroelectric R3c structure. The core-shell structure and the cationic charge disorder at the B-sites helped to establish the relaxor phase in 0.70BiFeO3-0.30BaTiO3 (Figure 9a). Substitution of Bi(Zn2/3Nb1/3)O3 further exacerbated the long-range order, thereby inducing a highly disordered RFE phase while maintaining the polarizability as high as *P*max = 36.7 μC cm−<sup>2</sup> and Δ*P* = 32.8 μC cm<sup>−</sup>2. They obtained *U*st = 3.7 J cm−3, *U*rec = 2.06 J cm−<sup>3</sup> at 180 kV cm−<sup>1</sup> and *U*st = 2.9 J cm−3, *U*rec = 1.98 J cm−<sup>3</sup> at 190 kV cm−<sup>1</sup> for 5 mol% and 8 mol% Bi(Zn2/3Nb1/3)O3-doped compounds, respectively. The solid solutions could also successfully deliver discharge energy within 0.5 μs. Wang et al. [82] reported relaxor behavior in chemically inhomogeneous, but electrically homogeneous (0.7-x)BiFeO3-0.3BaTiO3-xNd(Zr0.5Zn0.5)O3 for x = 0.05, 0.08, and 0.10, with a core-shell structure, and studied the energy storage performance at room temperature both in ceramics and multilayers prepared by a solid-state reaction route. The substitution of Nd(Zr0.5Zn0.5)O3 induced chemical inhomogeneity, revealed as a (Bi, Fe)-rich core for x = 0.05 and (Ba,Ti)-rich cores for x = 0.1, and complex multiphase microstructures with both (Bi, Fe)-rich and (Ba,Ti)-rich cores for x = 0.08. A high degree of chemical inhomogeneity and better electrical homogeneity of grains existed at x = 0.08, which led to a high Δ*P* and provided a difficult current path for electrical breakdown. As a result, *x* = 0.08 exhibited *U*rec ~2.45 J cm−<sup>3</sup> at *E* = 240 kV cm−<sup>1</sup> with *η* = 72%. Lu et al. [91] demonstrated excellent energy storage properties in a series of solid solutions composed of BiFeO3, SrTiO3, Nb2O5, and BiMg2/3Nb1/3O3 exhibiting a core-shell structure. Their study successfully demonstrated that Nb2O5 doping into BiFeO3-SrTiO3 and employing the third perovskite end member BiMg2/3Nb1/3O3 (BMN) could induce an insulating relaxor phase at room temperature without reducing the average ionic polarizability of the solid solution. It was found that the substitution of 1–3% Nb5+ for Ti4+ (B sites) in 0.6BiFeO3-0.4SrTiO3 suppressed the formation of oxygen vacancies and significantly reduced *p*-type conductivity compared with that of the undoped compound. A similar reduction in the *p*-type conductivity was observed for the 0.56BiFeO3-0.4SrTiO3-0.04BiMg2/3Nb1/3O3-xNb2O5

(x = 0–0.05) solid solution, and an enhanced BDS of 360 kV cm-1 was obtained at an optimized Nb2O5 content (x = 0.03). Such a high BDS was attributed to two factors: (i) the improved insulating character caused by Nb doping at B-sites related to the suppression of the formation of oxygen vacancies, and (ii) a core-shell microstructure with electrical homogeneity throughout the grains. The *x* = 0.03 composition exhibited a high performance of *U*rec = 6 J cm−<sup>3</sup> with *η* = 74.6%. Then, a much-improved *U*rec value with similar *η* could be achieved by optimizing the BiMg2/3Nb1/3O3 content in (0.6-y)BiFeO3-0.4SrTiO3- 0.03Nb2O5-yBiMg2/3Nb1/3O3 (y = 0.02–0.12) solid solutions. For y = 0.1, the BDS was enhanced further up to 460 kV cm<sup>−</sup>1, producing *U*rec = 8.2 J cm−<sup>3</sup> and *η* = 74.6%. A core-shell microstructure design provides a large BDS and large *U*rec, but at the same time, it has low efficiency because of the not-very-fast response of the PNRs to the electric field.

**Figure 9.** (**a**) Bright-field TEM image of a grain in (0.70-x)BiFeO3-0.30BaTiO3-xBi(Zn2/3Nb1/3)O3 with x = 0.05, illustrating a BiFeO3-rich core and BaTiO3-rich shell. (**b**) Broad dielectric anomalies associated with the core-shell structure. P-E loops for (**c**) x = 0.05 and (**d**) x = 0.08, respectively. Reproduced with permission from [81]. © 2021 American Chemical Society.

Qi et al. [92] employed a domain engineering technique to optimize the microstructure at the domain level and showed that this technique was very effective for enhancing *U*rec and *η*. Using it, they obtained large *P*max, large Δ*P*, and superior energy storage performance in 0.57BiFeO3-0.33BaTiO3-0.1NaNbO3 with the addition of 0.1 wt% MnO2 and 2 wt% BaCu(B2O5). Nanodomain engineering produced stripe-like PNRs embedded in featureless nanodomains. The stripe-like PNRs were rich in BiFeO3, whereas the matrix domains were rich in BaTiO3 and NaNbO3. An HR-TEM image of the solid solution is shown in Figure 10a. Stripe-like PNRs rich in BiFeO3 were dispersed in BaTiO3- and NaNbO3-rich featureless nanodomains. This structural heterogeneity at the domain level led to a rapid polarization response to the external *E*-field, unlike in the core-shell microstructure, and produced a hysteresis-free *P*−*E* loop, as well as a large *P*max. Moreover, NaNbO3 substitution increased the bandgap and helped to obtain a uniform and fine-grained microstructure, which was beneficial to enhance BDS up to 360 kV cm<sup>−</sup>1, producing *<sup>U</sup>*rec ≈ 8.12 J cm−<sup>3</sup> with *<sup>η</sup>* ≈ 90%. This solid solution also exhibited excellent thermal stability, with *<sup>U</sup>*rec = 8.12 J cm−<sup>3</sup> ± 10% in the temperature range −25 ◦C to 250 ◦C, and an ultrafast discharge rate (< 100 ns).

**Figure 10.** (**a**) Domain morphology exhibiting stripe-like PNRs in 0.57BiFeO3-0.33BaTiO3-0.1NaNbO3 with the addition of 0.1 wt.% MnO2 and 2 wt.% BaCu(B2O5). (**b**) Room-temperature *P*–*E* loops and *ε* measured under various electric fields for the same compound. *U*rec and *η* measured at (**c**) various electric fields and (**d**) temperatures. Figures reused with permission from [92]. © 2021 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

#### 4.2.2. Thin Films

In the past few decades, thin films have gained enormous scientific importance in a plethora of applications. The obvious advantages thin films possess over their bulk counterparts are that they save material and have reduced weight. In addition, it is well known that material properties drastically change when they are deposited in thin film forms, which paves the way for industrial applications. Dielectric thin films with thicknesses on the nano- or submicron scale have shown promising potential in the field of energy storage for low-power small devices. This is because of their extraordinarily high BDS (1 MV/cm) and energy density compared with bulk dielectrics.

Correia et al. [93] deposited a thin film of 0.4BiFeO3-0.6SrTiO3 onto a SrRuO3-electroded (100)-SrTiO3 substrate using the pulsed laser deposition (PLD) method, which can endure an electric field as large as 0.972 MV cm−1. Under such a high electric field, the film exhibited *U*rec = 18.6 J cm−<sup>3</sup> with *η* > 85%. It also featured a small temperature coefficient of capacitance (TCC < 11%) over a wide range of temperatures up to 200 ◦C. Pan et al. [94] demonstrated that the BDS of the same material could be enhanced more than by three times, up to 3.6 MV cm−1, when doped with Mn and deposited onto a Nb-doped SrTiO3 (001) substrate by the PLD technique. The high BDS of the Mn-doped 0.4BiFeO3-0.6SrTiO3 thin film was attributed to (i) a dense, uniform, and crack-free microstructure, (ii) highquality epitaxial growth, and (iii) low leakage current by reducing the Fe3+/Fe2+ valence fluctuation with Mn substitution. They obtained a colossal *U*rec = 51 J cm−<sup>3</sup> and *η* = 64%, which are comparable to those of lead-based thin films (*U*rec = 61 J cm−<sup>3</sup> and *η* = 33%) [95]. The thin film also featured high fatigue endurance quality over 2 × 107 cycles and good thermal stability over a wide range of temperatures, from −40 ◦C to 140 ◦C. In another

report, Pan et al. [96] utilized a domain engineering method to obtain strong relaxor behavior in (1−x)BiFeO3-xSrTiO3 thin films (x = 0.3–0.75) deposited by the PLD technique and found that the BDS of the films could be enhanced further, up to 4.46 MV cm−1, by increasing the SrTiO3 content. Atomic-scale microstructure analysis based on the STEM of BiFeO3-SrTiO3 films revealed that SrTiO3 disrupted the long-range FE order, and this disruption cascaded with an increase in SrTiO3 content. The incorporation of SrTiO3 could transform the micrometer-scale FE domains into nanoscale PNRs. Paraelectric SrTiO3 acts as a matrix for embedded PNRs and separates the PNRs such that the inter-PNR interaction almost vanishes. Because there are no inter-PNR interactions, the PNRs are very dynamic under an external electric field and produce slim *P*–*E* loops with very small *P*r values. The size of the PNRs continued to decrease with increasing SrTiO3 content, forming an almost domain-less feature for x = 0.75. This domain evolution reduced the domain switching energy, producing a slim *P*–*E* loop, while maintaining a large *P*max value and a colossal BDS. In addition to the domain perspective, SrTiO3 substitution enhanced the insulating character of the films by stabilizing the Fe3+/Fe2+ valence fluctuation and reducing the formation of oxygen vacancies, leading to further enhancement of BDS. For example, the BDS value of (1-x)BiFeO3-xSrTiO3 films increased from 2.77 MV cm−<sup>1</sup> to 4.46 MV cm−<sup>1</sup> as the SrTiO3 content increased from for *x* = 0.3 to *x* = 0.75. At such a high electric field, the films with *x* = 0.6 and 0.75 exhibited a giant *U*rec ~ 70 J cm<sup>−</sup>3.

Domain engineering techniques have been found to be more fruitful in thin films that exhibit polymorphic nanodomains, e.g., rhombohedral (*R*) and tetragonal (*T*) domains in a cubic paraelectric matrix. If these polymorphs have competitive free energy, Landau phenomenological theory predicts the weakening of polarization anisotropy and lowering of the energy barrier between the *R* and *T* polarization states [97]. It facilitates a flatter energy profile for polymorphic nanodomain RFEs compared with classic FEs and nanodomain RFEs, which minimizes the hysteresis while maintaining a high polarization (Figure 11). Pan et al. [97] demonstrated that the polymorphic domain engineering technique can produce ultra-high energy density with high efficiency in thin films of a (0.55−x)BiFeO3 xBaTiO3-0.45SrTiO3 solid solution (x = 0 to 0.4). Here, BiFeO3 and BaTiO3 were the hosts for the *R* and *T* domains, whereas SrTiO3 provided a cubic paraelectric matrix. The relaxor nature and energy storage performance of the (0.55−x)BiFeO3-xBaTiO3-0.45SrTiO3 solid solutions are shown in Figure 12. The incorporation of BaTiO3 gradually enhanced the relaxor nature, as can be seen from the wider peaks in the ε–T plots (Figure 12a), as well as the BDS for higher BaTiO3 contents. The BDS increased up to 4.9 MV cm−<sup>1</sup> for x = 0.3 and 5.3 MV cm−<sup>1</sup> for x = 0.4, compared with 3.2 MV cm−<sup>1</sup> for the x = 0 compound. This resulted in a maximum *U*rec of 112 J cm−<sup>3</sup> and 110 J cm−<sup>3</sup> for x = 0.3 and 0.4, respectively, with *η* > 80% (Figure 12). Moreover, the enhancement of the relaxor nature resulted in good temperature stability for x = 0.3 and 0.4 regarding their energy storage performance over a wide range of temperatures. Kurusumovic et al. [98] employed a combined defect engineering method to explore the energy storage performance of relaxor thin films of BiFeO3-BaTiO3 solid solutions doped with Mn. The combined approach of defect engineering consisted of an interval mono-layer by mono-layer deposition (LLD) and Mn addition. The LLD produced highly stoichiometric and perfectly crystalline films compared with standard deposited films, while the addition of Mn reduced the leakage current by creating vacancy trap centers. Using this method, they obtained an ultra-high value of *U*rec = 80 J cm−<sup>3</sup> at a BDS of 3.1 MV cm−<sup>1</sup> with *η* = 78% in 2.5 mol.% Mn-doped 0.25BiFeO3-0.75BaTiO3 thin films.

**Figure 11.** Design of RFEs via polymorphic nanodomain design. Comparative display of Landau energy profiles and *P*−*E* loops of an FE with micrometer-sized domains, an RFE with nanodomains, and an RFE with polymorphic nanodomains. The *P*<sup>R</sup> represents the polarization states along the rhombohedral (*R*) directions, and *P*<sup>T</sup> is along the tetragonal (*T*) direction. The shadowed area in the *P*-*E* loops represents the recoverable energy density. Figures reproduced with permission from [97]. © 2021 The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science.

**Figure 12.** Characteristics of (0.55−*x*)BiFeO3-*x*BaTiO3-0.45SrTiO3 (x = 0.0–0.4) films. (**a**) Temperature-dependent dielectric permittivity and loss tangent at a frequency of 1 MHz, (**b**) first-order reversal curve (FORC) *P*−*E* loops, (**c**) energy density and efficiency values with respect to applied electric fields up to breakdown fields, and (**d**) comparison of the energy storage performance at breakdown fields. Figures reproduced with permission from [97]. © 2021 The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science.

#### *4.3. Multilayered Structure*

Although the energy density of thin films is superior to that of bulk ceramics, the usability of thin films is limited for low-power applications because of their small volume. In this context, multilayered structures have received great attention because the technology behind them is well known and inexpensive. A multilayered ceramic consists of a number of thin ceramic layers with thicknesses on the micrometer scale and internal electrode layers stacked in parallel and connected through terminal electrodes. It features both high BDS (on the MV cm−<sup>1</sup> level) and large volume; therefore, it is very promising for practical high-power energy storage applications.

Wang et al. [76] reported a large improvement in the energy storage performance of a multilayered ceramic composed of 0.75(Bi0.85Nd0.15FeO3)-0.25BaTiO3 + 0.1wt.% MnO2 and Pt internal electrodes, compared with its bulk ceramic counterpart, as shown in Figure 13. With nine active layers having a total thickness of ~ 0.78 mm, they obtained a significant enhancement of the BDS to 540 kV cm−<sup>1</sup> with *U*st ~ 8.75 J cm<sup>−</sup>3, *U*rec~ 6.74 J cm<sup>−</sup>3, and *η*~77%. Yan et al. [99] reported similar energy storage properties in (1-x)(0.67BiFeO3-0.33BaTiO3) xNa0.73Bi0.09NbO3 multilayered ceramics with A-site cation vacancies. For an optimized content of Na0.73Bi0.09NbO3, i.e., x = 0.12, the multilayer exhibited *U*rec = 5.57 J cm−<sup>3</sup> at *E* = 410 kV cm−<sup>1</sup> with *η* = 83.8%. A multilayer ceramic with 7-μm-thick 0.61BiFeO3- 0.33Ba0.8Sr0.02TiO3-0.06La(Mg2/3Nb1/3)O3 dielectric layers was reported to exhibit a much higher BDS, > 740 kV cm−<sup>1</sup> [100]. This device featured a high *U*rec of 10 J cm−<sup>3</sup> with *η* ~ 72% at *E* = 730 kV cm−1. Wang et al. reported a large Δ*P* of ~ 34 μC cm−<sup>2</sup> and a high BDS of 700 kV cm−<sup>1</sup> in a multilayered ceramic composed of 16-μm-thick electrically homogeneous 0.62BiFeO3-0.3BaTiO3-0.08NdZr1/2Zn1/2O3 ceramic layers with a core-shell microstructure [82]. With seven ceramic layers having an active electrode area of 33 mm2, *U*rec as high as 10.5 J cm−<sup>3</sup> with *η* = 87% was obtained. Lu et al. [91] obtained BDS > 1 MV cm−<sup>1</sup> in a multilayered ceramic with 10 mol% Bi(Mg2/3Nb1/3)O3 and 3 mol% Nb2O5-doped 0.6BiFeO3-0.4SrTiO3 with a core-shell microstructure. The device, with 8-μm-

thick dielectric layers, exhibited *U*rec = 15.8 J cm−<sup>3</sup> under *E* > 1 MV cm−<sup>1</sup> with *η* = 75.2%. Wang et al. [101] reported a high BDS of 953 kV cm−<sup>1</sup> in a multilayered ceramic with 8-μm-thick 0.57BiFeO3-0.3BaTiO3-0.13Bi(Li0.5Nb0.5)O3 layers (Figure 14) having a coreshell structure and a 5-mm<sup>2</sup> active electrode area, which was greatly improved compared with that of the bulk ceramic counterpart (260 kV cm-1). While the bulk ceramic featured a *U*rec ~ 3.64 J cm−<sup>3</sup> at *E* ~ 260 kV cm−<sup>1</sup> with *η* ~ 75%, the multilayered ceramic exhibited a *U*rec = 13.8 J cm−<sup>3</sup> with *η* = 81% because of the large enhancement of the BDS. The multilayered ceramic showed good temperature stability (< 10%) and frequency independence (< 5%) of *U*rec from 0.01 to 1 Hz, as well as fatigue resistance (< 5%) during 104 cycles of unipolar *<sup>P</sup>*−*<sup>E</sup>* loop tests in the temperature range from RT to 100 ◦C at *<sup>E</sup>* = 400 kV-cm<sup>−</sup>1.

**Figure 13.** Energy storage performance of 0.75(Bi0.85Nd0.15)FeO3-0.25BaTiO3 multilayers. (**a**) P–E loops at various electric fields at room temperature, (**b**) P–E loops at various temperatures. *U*st, *U*rec, and *η* with (**c**) electric field and (**d**) temperature. An SEM image of the multilayers is shown in the inset of (**a**). Figures reproduced with permission from [76]. © 2021 The Royal Society of Chemistry.

**Figure 14.** (**a**) Room temperature unipolar *P*–*E* loops at various electric fields and (**b**) calculated energy storage properties of 0.57BiFeO3-0.3BaTiO3-0.13Bi(Li0.5Nb0.5)O3 multilayers. Figures reused under Creative Commons Attribution 3.0 Unported Licence [101]. © 2021 The Royal Society of Chemistry.

#### **5. Critical Issues and Strategy**

#### *5.1. Leakage Current Control*

High discharge energy density and energy efficiency are the primary requirements of an EES system. However, most BiFeO3-based RFEs suffer from a low efficiency (<85%) despite having ultra-high discharge energy densities. Leakage-related high conductivity in BiFeO3-based RFEs produces apparent large values of *P*<sup>r</sup> and wide hysteresis loops, which lead to a higher loss density. The conductivity usually increases at high temperatures, limiting their potential for high-temperature applications.

The high leakage current in BiFeO3-based compounds is associated with the loss of Bi2O3 during its synthesis at elevated temperatures (950–1050 ◦C), which creates ionic vacancies and Fe3+/Fe2+ valence fluctuation [102–105]. Aliovalent doping in BiFeO3-based compounds has been effective for reducing the leakage current density by several orders of magnitude [104,105]. However, these compounds are fired at high temperatures for synthesis in both bulk ceramics and multilayered structures, exposing them to Bi2O3 loss. Adopting a low-temperature sintering method is expected to reduce the Bi2O3 loss and the associated leakage current densities by a few more orders. The addition of appropriate glass additives and/or low melting point compounds is expected to significantly reduce the sintering temperature [106], thereby reducing the possibility of Bi2O3 volatilization. Low-temperature sintering is also beneficial for the enhancement of BDS, as discussed in Section 5.2.

#### *5.2. Microstructure Engineering*

In Figure 15, we plot *U*rec as a function of BDS for BiFeO3-based binary and ternary solid solutions fabricated in bulk ceramics, ceramic multilayers, and thin films. A high BDS usually provides a higher recoverable energy density. Therefore, enhancing the BDS is one of the most effective ways to increase the energy storage performance of RFEs. Microstructure engineering, such as the design of core-shell structures and domain engineering, has been found to be very effective for enhancing the BDS and realizing a dynamic relaxor ferroelectric phase. The BDS depends on the grain size (BDS <sup>∝</sup> 1/√*<sup>G</sup>* (*G* = average grain size)) and size distribution. Therefore, it is imperative to design a fine-grained microstructure with a uniform size distribution to obtain a high BDS. Another way to improve the BDS is to prepare a highly dense microstructure. Pores or cavities present in microstructures are usually filled with gaseous or liquid phases with lower permittivity than those of solid dielectrics. As a result, the voltage across the cavities or pores (*V*c) is enhanced as per the following relation [107]:

$$V\_c = \frac{V\_{app}}{\left[1 + \frac{\varepsilon\_c}{\varepsilon\_r} \left(\frac{d}{I} - 1\right)\right]} \tag{10}$$

where *V*app is the applied voltage, *ε*<sup>c</sup> and *ε*<sup>r</sup> are the permittivity of the cavity and dielectric, respectively, and *t* and *d* are the size of the cavity and the thickness of the dielectric, respectively. A small pore can create a large electric field across it and cause local breakdown and internal discharge even at a low external voltage. Therefore, a highly dense microstructure is essential. Because BiFeO3-based materials are prone to the loss of Bi and Fe valence fluctuations during sintering at high temperatures, resulting in high electrical conductivity, sintering at low temperatures could be very beneficial. However, a detailed investigation of the microstructure control for the purpose of enhancing the energy storage performance employing the low-temperature sintering technique in BiFeO3-based dielectrics is yet to be conducted. The two-step sintering (TSS) method modified by Chen and Wan [108] is one of the most cost-effective and simple methods to produce ultra-high-density materials and fine grains. In this technique, high-temperature heating is performed for a few minutes, and then the material is allowed to cool to the sintering temperature, where the sample is sintered for a prolonged time. However, the sintering temperature should be wisely chosen so that densification without further grain growth occurs.

**Figure 15.** Recoverable energy density as a function of breakdown strength for BiFeO3-based bulk ceramics of binary solid solutions and bulk ceramics, microstructure-designed bulk ceramics, and multilayers and thin films of ternary solid (TS) solutions.

#### *5.3. Band Gap Engineering*

Electronic breakdown is a crucial intrinsic mechanism for the breakdown of a dielectric in a large electric-field regime. Above a certain electric field, electrons in the valence band gain sufficient energy to jump to the conduction band. This results in an increase in the electron density in the conduction band and leads to a large current discharge. Therefore, an insulator exhibiting a wide forbidden energy gap can withstand a high electric field. The empirical relation between the BDS and energy band gap is given by the following relation [109]:

$$E\_B = 24.442 \exp\left(0.315 \sqrt{E\_\mathcal{g} \omega\_{\text{max}}}\right) \tag{11}$$

where *E*<sup>g</sup> is the energy bandgap, and *ω*max is the maximum phonon frequency. Pure BiFeO3 has a direct bandgap of 3 eV [110]. Only a few studies have attempted to enhance the BDS of BiFeO3-based RFEs via modulation of the energy bandgap. Qi et al. [92] reported the enhancement of BDS by doping of high-band-gap materials such as NaNbO3 (*E*<sup>g</sup> ~ 3.28 eV) in 0.67BiFeO3-0.33BaTiO3 solid solutions. They demonstrated that the bandgap of the solid solutions (0.67-x)BiFeO3-0.33BaTiO3-xNaNbO3 (0 ≤ x ≤ 0.15) increases monotonically with increases in x from ~ 2.6 eV for x = 0 to ~ 2.95 eV for x = 0.15. The enhancement of BDS followed a similar trend with an increase in NaNbO3 content, where the average BDS increased from 230 kV cm−<sup>1</sup> for x = 0 to 420 kV cm−<sup>1</sup> for x = 0.15. Further studies on bandgap engineering can be explored to enhance the BDS and energy storage performance of BiFeO3-based relaxor ferroelectrics.

#### *5.4. Electromechanical Breakdown*

Another issue that has not gained much attention but is very critical in BiFeO3-based relaxor ferroelectrics is the electromechanical breakdown. Because a high electric field is required to obtain a large recoverable energy density, the dielectrics are under extreme electrostrictive strain, increasing the mechanical breakdown. This issue should be taken very seriously from the device viewpoint, as the devices undergo a large number of chargedischarge cycles. Although high recoverable energy density with high efficiency is obtained at high electric fields under laboratory conditions, the same samples might not be suitable for application purposes where they totally collapse because of mechanical failure under a large number of charge-discharge cycles. However, systematic studies on the issues of electromechanical breakdown in BiFeO3-based dielectrics are lacking. Rare-earth-ion doped BiFeO3 compounds with their composition lying across the morphotropic phase boundary could be interesting in this regard as these compounds feature high electromechanical strain. For example, Walker et al. [111] reported cycle-dependent large electromechanical strain in Sm-doped BiFeO3 polycrystalline samples, associated with the electric field-induced phase transition and ferroelectric/ferroelastic domain switching.

#### **6. Summary**

BiFeO3-based relaxor ferroelectrics are projected to be potential Pb-free candidates for application in the field of high-energy-density storage and high-power-delivery systems. They are mostly fabricated in the form of ceramics, multilayers, and thin films from binary or ternary solid solutions with other perovskite oxides. BiFeO3-BaTiO3 and BiFeO3- SrTiO3, with their compositional ratios lying in the morphotropic phase boundary, are the two most widely studied BiFeO3-based binary solid solutions for high-energy-density storage. However, the typical value of recovering energy density of bulk ceramics of BiFeO3-based binary solid solutions is well below 3 J cm−3, with low efficiency (typically below 50%). Ternary solid solutions exhibiting specific microstructures, such as core-shell structures and stripe-like nanodomains, have shown significant enhancement of the energy density properties. For example, 0.5BiFeO3-0.4SrTiO3-0.03Nb2O5-0.1BiMg2/3Nb1/3O3 bulk ceramic with a core-shell microstructure exhibited excellent energy storage properties, with *U*rec ~ 8.2 J cm−<sup>3</sup> and *η* ~ 74%. The nanodomain-engineered bulk ceramic 0.57BiFeO3- 0.33BaTiO3-0.1NaNbO3 with the addition of 0.1 wt% MnO2 and 2 wt% BaCu(B2O5) also featured a similar value of *U*rec (~ 8.12 J cm<sup>−</sup>3); however, it had a higher value of *η* ~ 90%. The microstructure design approach has been found to be attractive for enhancing the energy storage properties of thin films and multilayers. The BDS values of thin films and multilayers can be as high as MV cm-1, which is very helpful for further increasing the recoverable energy density. Thin films of a (0.55−*x*)BiFeO3-*x*BaTiO3-0.45SrTiO3 solid solution (*x* = 0–0.4) containing polymorphic nanodomains showed a recoverable energy density of ~ 110 J cm−<sup>3</sup> at a BDS ~ 5 MV/cm with an efficiency of ~ 80%, which is the highest among BiFeO3-based systems. The achievement of such excellent energy storage properties is very encouraging for applications in low-power, small-size devices. However, for large-scale applications, it is necessary to focus on the fabrication of multilayers, as they deliver high power and are easy to fabricate. Although relatively few studies have been conducted on BiFeO3-based multilayers, their energy storage performance is very encouraging (*U*rec ~ 15 J cm−3), with a high BDS of ~ 1 MV/cm. More studies should be conducted in the field of BiFeO3-based multilayers to optimize structural parameters such as the thickness of dielectric layers and number of layers so that a better energy storage device that can deliver high output power can be fabricated at a low cost.

**Author Contributions:** B.D. and K.-H.C. equally contributed to this work. K.-H.C. conceptualized the idea, B.D. reviewed literatures, drafted and wrote the manuscript. B.D. and K.-H.C. read and edited the manuscript for final submission. K.-H.C. acquired the funding. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the National Research Foundation (NRF) of Korea funded by the Ministry of Education [NRF-2019R1I1A3A01058105, NRF-2018R1A6A1A03025761] and Ministry of Science and ICT (Grand Information Technology Research Center support program) [IITP-2021-2020-0-01612] supervised by the Institute for Information & Communications Technology Planning & Evaluation.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


### *Article* **Harvesting Energy from Bridge Vibration by Piezoelectric Structure with Magnets Tailoring Potential Energy**

**Zhiyong Zhou 1,2, Haiwei Zhang 1, Weiyang Qin 3, Pei Zhu 1,\*, Ping Wang 4,\* and Wenfeng Du <sup>1</sup>**

	- 10160091@vip.henu.edu.cn (Z.Z.); hwzhang@henu.edu.cn (H.Z.); 10160016@vip.henu.edu.cn (W.D.)

**Abstract:** Bridges play an increasingly more important role in modern transportation, which is why many sensors are mounted on it in order to provide safety. However, supplying reliable power to these sensors has always been a great challenge. Scavenging energy from bridge vibration to power the wireless sensors has attracted more attention in recent years. Moreover, it has been proved that the linear energy harvester cannot always work efficiently since the vibration energy of the bridge distributes over a broad frequency band. In this paper, a nonlinear energy harvester is proposed to enhance the performance of harvesting bridge vibration energy. Analyses on potential energy, restoring force, and stiffness were carried out. By adjusting the separation distance between magnets, the harvester could own a low and flat potential energy, which could help the harvester oscillate on a high-energy orbit and generate high output. For validation, corresponding experiments were carried out. The results show that the output of the optimal configuration outperforms that of the linear one. Moreover, with the increase in vehicle speed, a component of extremely low frequency is gradually enhanced, which corresponds to the motion on the high-energy orbit. This study may give an effective method of harvesting energy from bridge vibration excited by moving vehicles with different moving speeds.

**Keywords:** energy harvesting; bridge vibration; vehicle; mono-stable energy harvester; linear energy harvester; moving speed

#### **1. Introduction**

Bridge safety has drawn great concern in recent decades owing to the developments in the cross-sea bridge, cable-stayed bridge, and suspension bridge [1]. To ensure the safety of the bridge, a sensor network should be incorporated into it so that many node sensors are distributed to monitor the state of the bridge to prevent hazards caused by degradation [2]. For the node sensors, the replacement of the battery is a difficult challenge [3]. Frequent battery replacement is excessively expensive and nearly impossible in some dangerous and special hard-to-reach areas [4]. To tackle the difficulty, harvesting energy from bridge vibration to power sensors is emerging as a promising solution. In principle, the vibration energy of structure can be converted to electric energy through three methods, i.e., the triboelectric, electro-magnetic, and piezoelectric mechanisms [5].

The piezoelectric scenario generally scavenges energy from bridge vibration excited by travelling vehicles. Developing energy harvesting devices coupled with the bridge structure has been drawing more and more attention from scholars [6]. Assadi et al. [7] attached a piezoelectric patch to a simply supported beam with a mass moving on it. When the moving mass travelled on the beam, the piezoelectric patch could produce voltage due to deflection. The analytical and experimental results indicated that the speed of the moving mass had a

**Citation:** Zhou, Z.; Zhang, H.; Qin, W.; Zhu, P.; Wang, P.; Du, W. Harvesting Energy from Bridge Vibration by Piezoelectric Structure with Magnets Tailoring Potential Energy. *Materials* **2022**, *15*, 33. https://doi.org/10.3390/ ma15010033

Academic Editors: Marc Cretin, Sophie Tingry and Zhenghua Tang

Received: 2 December 2021 Accepted: 17 December 2021 Published: 21 December 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

significant influence on voltage output. Ali et al. [8] used a linear single-degree-of-freedom model as an energy harvester to harvest the vibration energy of a simply supported beam excited by the motion of a point load. Galchev et al. [9] designed an inertial energy harvester to convert traffic-induced bridge vibrations to electric energy. The average power of this fabricated device was measured at different positions of a suspension bridge. Peigney and Siegert [10] designed and tested a cantilever piezoelectric harvester to harvest vibration energy from traffic-induced vibrations in bridges. Erturk [11] derived analytical expressions of transient response in the time domain of piezoelectric output power. Zhang et al. [12] conducted a numerical parametric study and derived the exact analytical solution of a piezoelectric energy harvester excited by beam vibration under moving harmonic loads. Amini [13] numerically investigated the harvested power from the vibrations of a beam excited by multi-moving loads. The results showed that the efficiency of the harvester increased with an increase in moving velocity. Xiang et al. [14] investigated a pavement system under moving vehicles to scavenge traffic-induced energy using piezoelectric transducers. When the velocity of a moving vehicle was close to a critical value, the power output of the piezoelectric harvester was optimal. Romero et al. [15] analyzed the energy harvesting performance of railway bridges in different operational conditions. The results showed that the passage of trains had a major effect on energy harvesting performance. Therefore, the typical cantilevered (linear) harvester always had an enormous challenge in scavenging energy from bridge vibration since the abundant vibration energy of a bridge excited by different vehicles exists within a broad frequency band.

At present, harvesting energy from base vibrations is being widely investigated, in which the magnetic or axial loading forces are often used to enhance harvesting performance. Accordingly, the nonlinearity caused by the introduced magnetic forces was studied in depth so as to improve the enhancement in harvesting ability [16]. With the development of the study, the configurations with different kinds of stability, e.g., monostability [17], bi-stability, tri-stability [18], quad-stability [19], and even penta-stability, were proposed and proved to be able to broaden the operating bandwidth effectively. For example, Fan et al. [20] proposed a mono-stable piezoelectric energy harvester to improve the efficiency of energy extraction from low-level excitations by two stoppers and four magnets. Naseer et al. [21] theoretically investigated harvesting energy from vortex-induced vibrations by using nonlinear attractive magnetic forces. Jang and Chen [22] derived the Fokker–Plank–Kolmogorov equation of a mono-stable Duffing oscillator with piezoelectric coupling. The effects of the bandwidth, initial conditions, and cubic nonlinearity on voltage were numerically studied. Erturk and Inman [23] conducted the theoretical and experimental studies on the high-energy orbits of a bi-stable energy harvester. Moreover, the relative advantages and trade-offs of bi-stable and mono-stable harvesters were presented by Zhao and Erturk [24]. Although the mono-stable energy harvester is more robust and reliable than the bi-stable one under stochastic excitations, the bi-stable energy harvester should be designed deliberately to work more efficiently for a given excitation. Kim and Seok [25] developed the mathematical model of an energy harvester with multi-stable (mono-, bi-, and tri-stable) characteristics. The results of the simulation showed that the tri-stable energy harvester had a wider bandwidth than both the bi-stable and linear ones. Wang et al. [26] utilized the tri-stable characteristics to enhance the energy harvesting performance of a galloping piezoelectric energy harvester.

In previous studies, the linear vibration energy harvesters had been adopted to scavenge energy from bridge vibration, but it was found that they could not work efficiently under broadband excitations. Hence, nonlinearity was introduced into the design of the energy harvester under different excitation types, e.g., the harmonic, stochastic, or flowinduced vibrations. The multi-stability created by magnetic interaction has been proved to be particularly effective. Striking improvement in energy harvesting performance has been demonstrated. However, in the multi-stable structures, the energy required to cross the potential barrier and produce a big deflection is fairly large. In this study, we tried using nonlinear magnetic forces to obtain an optimal shape for potential energy. To the best

of the authors' knowledge, using magnetic nonlinearity to tailor potential energy has not been studied in the harvesting of bridge vibration energy. These techniques can produce large amplitudes by creating a desired potential energy shape rather than a multi-stable shape. In this paper, a scenario of harvesting the energy of a bridge by mono-stability is proposed, in which the magnetic forces are introduced to form a desirable potential energy and a high-energy orbit. The mono-stable piezoelectric energy harvester (MPEH) can produce a large output if it oscillates in the high-energy orbit. It was proved that the MPEH could reach a high efficiency in harvesting the vibration energy of a bridge. Firstly, for the particular mono-stable characteristic, corresponding analyses on potential energy, restoring force, and stiffness were carried out. The results could help determine the optimal shape for potential energy. Then, experimental studies were conducted, and the dynamic responses and electrical outputs were investigated under different conditions. Finally, a summary is presented and some conclusions are drawn at the end of this paper.

#### **2. Energy Harvester for Bridge Vibration Excited by Moving Vehicles**

The schematic of MPEH is shown in Figure 1, which is composed of a cantilever beam with two piezoelectric patches attached to the root, a tip magnet fixed at the beam's free end, and two fixed magnets. The three magnets are placed such that the tip and fixed magnets are magnetically attractive. By adjusting the separation distance (*d*) and gap distance (*dg)*, the potential energy can be tailored to own a flat-shaped bottom, which is beneficial for the MPEH to execute a large-amplitude vibration. In particular, the directions of the magnetic and elastic forces are opposite to each other, which could move the operating frequency band toward the low frequency, thus reducing the energy required to produce a large deflection of the harvester and improving the voltage and power outputs. The classical linear piezoelectric energy harvester (LPEH) often shows inefficiency since the abundant vibration energy excited by moving vehicles at different speeds exist in the form of a broad frequency band with small magnitude. In order to collect more electric energy, a large LPEH (especially a very long substrate layer with a length *Ls*) is required; however, it is difficult to meet the energy supply requirements of the densely distributed sensors. Therefore, scavenging energy from bridge vibration by linear structures has proved to be an enormous challenge. Compared with LPEH, the MPEH can attain the high-energy orbit under the excitation of passing vehicles and generate a high output.

**Figure 1.** Schematic diagram of MPEH.

As Figure 2 shows, the MPEH is fixed to a bridge with the length of *Lb* and the thickness of *Tb*. Then, as a moving vehicle travels on the bridge at speed *v*, its load excites the elastic bridge to oscillate with the acceleration *ab*(*t*). Thus, the MPEH oscillates and generates electric power through piezoelectric materials. According to previous related studies, the bridge can be modeled as an Euler–Bernoulli beam, which is a simply supported elastic beam with infinite degrees of freedom. The moving vehicle travelling on this beam with hinged-hinged boundary conditions can be simplified as a moving mass.

**Figure 2.** Schematic diagram of energy harvesting from bridge vibration excited by a moving vehicle.

#### **3. Analysis of Potential Energy, Restoring Force, and Stiffness**

The total potential energy of MPEH consists of two parts: the elastic potential energy of the substrate and the magnetic potential energy that originated from the magnets.

According to the Euler–Bernoulli beam theory, the strain of the piezoelectric cantilever is proportional to the second spatial derivative of deflection. Hence, the potential energy calculation in this paper has been based on the following two assumptions: (a) the flat section perpendicular to the center line of the beam before deformation still stays flat after deformation; (b) the plane of the deformed cross section is perpendicular to the deformed axis. Then, the elastic potential energy of the piezoelectric cantilever beam can be given by the following equation [27]:

$$\mathrm{d}L\_{\mathbf{c}} = \frac{1}{2} E\_{\mathrm{s}} I\_{\mathrm{s}} \int\_{0}^{L\_{\mathbf{s}}} \left( \frac{\partial^{2} w(\mathbf{x}, t)}{\partial \mathbf{x}^{2}} \right)^{2} d\mathbf{x} + E\_{p} I\_{p} \int\_{0}^{L\_{p}} \left( \frac{\partial^{2} w(\mathbf{x}, t)}{\partial \mathbf{x}^{2}} \right)^{2} d\mathbf{x} \tag{1}$$

where *Es* is the Young modulus of the substrate; *Is* is the inertia moment of the substrate given by *Is* <sup>=</sup> *WsT*<sup>3</sup> *s* <sup>12</sup> (*Ws* and *Ts* are the width and thickness of the substrate, respectively); *Ls* is the length of the substrate; *Ep* is the bending stiffness of the piezoelectric patch; *Ip* is the total inertia moment of two piezoelectric patches on both sides of the substrate, which can be given by *Ip* <sup>=</sup> *WpTp*(4*T*<sup>2</sup> *p*+6*TpTs*+3*T*<sup>2</sup> *s* ) <sup>6</sup> (*Wp* and *Tp* are the width and thickness of the piezoelectric patch, respectively); *Lp* is the length of the piezoelectric patch; *w*(*x*, *t*) is the displacement of the cantilever beam.

The magnetic potential energy can be obtained based on the assumption of a magnetic dipole. Thus, the magnetic potential energy generated by two fixed magnets upon the tip of the magnet can be expressed as follows [28]:

$$\mathcal{U}L\_{\rm m} = -\frac{\mu\_0 a\_1 a\_2}{2\pi} \left\{ \left( w(L\_\circ, t) - \frac{d\_\mathcal{S}}{2} \right)^2 + d^2 \right\}^{-\frac{3}{2}} - \frac{\mu\_0 a\_1 a\_3}{2\pi} \left\{ \left( w(L\_\circ, t) + \frac{d\_\mathcal{S}}{2} \right)^2 + d^2 \right\}^{-\frac{3}{2}} \tag{2}$$

where *μ*<sup>0</sup> is the permeability constant; *a*<sup>1</sup> and *a*<sup>2</sup> (*a*3) are the effective magnetic moment of the tip magnet and the fixed magnets, respectively; *d* is the horizontal distance from the center of the tip magnet to the center of the fixed magnets; *dg* is the gap distance between the two fixed magnets.

Since the potential energy is determined by the separation distance *d*, the total potential energy is computed with the distance *d* varying so as to illustrate the influence of *d*. The results are illustrated in Figure 3. The system parameters are given in Table 1. It can be seen from Figure 3 that as *d* decreases from 90 mm to 14 mm, the potential energy experiences the linear stable state, the mono-stable state, and then the bi-stable state. Specifically, when *d* is quite large, the system acts nearly as a linear one resulting from a very weak magnetic force. With a decrease in the separate distance *d*, the magnetic force will become large. The potential energy could own a steep bottom, implying that the MPEH may have a relatively large amplitude of vibration under external excitations. The ideal mono-stable configuration is expected to make a strong magnetic coupling between the magnetic force and the elastic force. This desirable case can provide a large beam deflection and then a high-voltage output under low-level bridge excitations. Especially if *d* is smaller than a critical value, i.e., 14 mm, the magnetic force will be larger than the elastic restoring force. With two stable equilibrium positions emerging, the MPEH becomes a bi-stable system. However, in the bi-stable state, two potential wells and a potential barrier will appear, which would need more excitation energy. Thus, for weak excitations, the bi-stable MPEH may not be the best option. In contrast, the mono-stable state near the bi-stable one is the most appropriate choice.

**Figure 3.** Potential energy for different values of separation distance.


**Table 1.** Properties for the analysis of potential energy function.

To illustrate the variation of the MPEH's characteristics with respect to *d*, the restoring force and stiffness at different separation distances are depicted in Figure 4. It is obvious that the restoring force and stiffness are greatly reduced with a decrease in *d*. The reason is that the attractive magnetic force will keep increasing when *d* decreases and will reach nearly an equal value to the elastic force at a critical distance. The MPEH will have an extremely low restoring force and bending stiffness. Therefore, even under the weak vibration of the bridge, the large deflection of the MPEH is likely to happen, thereby producing a high-voltage output. In contrast, the linear harvester only has the satisfactory harvesting efficiency while the excitation frequency matches the resonance frequency. However, in the practical environment, the speeds and weights of vehicles passing the bridge usually exhibit stochastic characteristics. Therefore, the vibration energy generated by passing vehicles distributes over a fairly broad frequency band. Thus, in terms of practical excitation, the MPEH may be superior to the LPEH.

**Figure 4.** (**a**) Restoring force; (**b**) stiffness at various separation distances.

#### **4. Experiment Setup**

In order to show the dynamic response and performance of the electric output of the MPEH, a prototype of the MPEH is designed and fabricated. The experimental setup is shown in Figure 5. The model bridge is made from an acrylic sheet with the dimensions of 1300 mm × 130 mm × 8 mm (seen in Figure 5a). The moving vehicle is simulated by a moving steel ball with the mass of 1.003 kg, which can be accelerated from different heights of the acceleration track, as shown in Figure 5b. The guardrails are installed on both sides of the bridge to prevent the steel ball from falling off the bridge. The schematic diagram of the experiment equipment is presented in Figure 5c. The MPEH is fixed in the middle of the bridge, as shown Figure 6. The related parameter values are listed in Table 1. The Piezoelectric patch pasted onto the substrate of the MPEH is connected to a resistance box. The strain sensor (120-5AA) is attached to the substrate of the MPEH near the piezoelectric patch. When it passes the bridge, the moving load will excite the bridge to oscillate. As for the three magnets, one is attached to the free end of the piezoelectric beam, and the other two magnets are fixed at the fixture. The separation distance (*d*) and gap distance (*dg*) are set to 17 mm and 45 mm, respectively. The LPEH is fabricated as well, as shown in Figure 7, which is similar to the MPEH in configuration but does not have the two fixed magnets. In the experiments, the LPEH and MPEH were put in the same testing environment for comparison. The strain response and voltage output of the harvester under bridge vibration were displayed and recorded by a digital signal acquisition device (DH5922N, Dong Hua). The value of the sampling frequency was set to 1000 Hz. In each test, to simulate different speeds of vehicles, the steel ball was released from different heights of the acrylic track.

**Figure 5.** The experimental setup: (**a**) test section, (**b**) acceleration section, and (**c**) schematic diagram of equipment.

**Figure 6.** Prototype of the MPEH: (**a**) top view and (**b**) side view.

**Figure 7.** Prototype of the LPEH.

#### **5. Results and Discussions**

In the experiment, both the LPEH and MPEH were installed in the middle span of the bridge. A strain sensor and a piezoelectric patch were bonded to the root of the substrate so we could obtain the strain response of the system and the dynamic output voltage. Figure 8 gives the variance of strain and power density (W/m3) at the load resistance of 0.9 MΩ for different moving speeds. To control the moving speed, the steel ball was put on an inclined track and released at different heights. The running interval was controlled to be between 0.50 s and 1.07 s, corresponding to the practical interval of a vehicle passing through bridges. From Figure 8, it is clear that for both the LPEH and MPEH, their variances of strain and power density increase with the moving speed. Especially for the speeds higher than 1.82 m/s, the MPEH's power density increases rapidly and outperforms that of LPEH, and so does the MPEH's strain.

**Figure 8.** (**a**) Variance of strain; (**b**) power density of LPEH and MPEH for different moving speeds.

To show the dynamics of the MPEH and LPEH, the time histories of strain are illustrated in Figure 9, in which six moving speeds (*v* = 1.07, 1.53, 1.82, 2.07, 2.30, and 2.50 m/s) are considered. In Figure 9, at *t* = 0 s, the vehicle arrives at the bridge, while the dashed line denotes the moment the vehicle leaves the bridge. The label "On bridge" represents the time interval of the vehicle moving on the bridge, while the label "Leaving bridge" represents the time interval after the vehicle leaves the bridge. In all cases, it is apparent that the strain of the MPEH is much larger than that of the LPEH when the vehicle travels on the bridge. This can be explained by the fact that an impact load acts on the bridge at the instant when the vehicle just enters the bridge. The MPEH is designed to have two magnets, with which the potential energy can be tailored to have a wide and flat bottom. This potential energy is beneficial for the occurrence of a large deflection. Thus, the small vibration of the bridge can trigger the large deflection of the MPEH, which is desirable in energy harvesting. Specifically, at a slow speed, e.g., *v* = 1.07 m/s, the maximum strain value of the MPEH can reach 3.6 × <sup>10</sup>−5, which is 164% higher than that of the LPEH (1.36 × <sup>10</sup><sup>−</sup>5). In both the MPEH and LPEH, the maximum strain increases with the moving speed. In any case, the vibration period of the MPEH is much longer than that of the LPEH, which is beneficial to harvesting weak vibration energy. After the vehicle leaves the bridge, the responses of the LPEH and MPEH exhibit a characteristic of decay vibration. Although the amount of time that the vehicle moves on the bridge is not very long due to its relatively fast speed (such as *v* = 2.07 m/s, 2.30 m/s, or 2.5 m/s), the frequency and amplitude of a large strain is satisfactory. From the comparison of strains between the MPEH and LPEH, it can be concluded that the deflection of the MPEH is much larger than that of the LPEH, particularly during the period when the vehicle is moving on the bridge.

**Figure 9.** Strain responses of the LPEH and the MPEH for the vehicle travelling on and leaving the bridge: (**a**) *v* = 1.07 m/s, (**b**) *v* = 1.53 m/s, (**c**) *v* = 1.82 m/s, (**d**) *v* = 2.07 m/s, (**e**) *v* = 2.30 m/s and (**f**) *v* = 2.50 m/s.

Figure 10 shows the open-circuit voltages of the LPEH and MPEH when the vehicle travels on and leaves the bridge. By comparing Figure 9 with Figure 10, it can be concluded that the deflection of the harvester is closely related to the output voltage. Moreover, from Figure 10, it can be seen that the voltage of the MPEH is significantly larger than that of the LPEH. In all cases, the large output voltage is generated during the period when the vehicle is travelling on the bridge, in which the vibration amplitude makes a major contribution to the voltage output. For a slower speed of 1.53 m/s, the voltage amplitudes of the LPEH and MPEH are 0.41 V and 0.57 V, respectively. The vibrations caused by the impact load and the moving vehicle are very weak, which make the LPEH execute a small amplitude vibration. For the MPEH, owing to the nonlinear magnetic forces, the vibration amplitude is relative large under weak excitations. Thus, the MPEH could improve the voltage output at slower speeds. By increasing the moving speed, the bridge vibration caused by the moving vehicle becomes violent. Thus, the amplitude of the voltage output increases accordingly, resulting in a high harvesting efficiency. It should be noted that when the speed of the vehicle reaches a certain level (such as 2.30 m/s or 2.50 m/s), there is no significant increase in the voltage amplitude. For example, the maximum voltage of the MPEH reaches 6.94 V for the speed of 2.30 m/s. However, as the moving speed increases to 2.50 m/s, the maximum voltage only has a small increase and reaches 7.49 V. This is because the MPEH's two magnets are prompted to protect the piezoelectric material from damage when the beam deflection is too large.

**Figure 10.** Voltage output of the LPEH and MPEH for the vehicle travelling on and leaving the bridge: (**a**) *v* = 1.07 m/s, (**b**) *v* = 1.53 m/s, (**c**) *v* = 1.82 m/s, (**d**) *v* = 2.07 m/s, (**e**) *v* = 2.30 m/s and (**f**) *v* = 2.50 m/s.

In order to better understand the dynamic characteristics of the LPEH and MPEH, their frequency spectra of output voltage are shown Figure 11. The peak of the LPEH is located near 14.5 Hz. In contrast, the MPEH's peak changes with the vehicle speed. For different moving speeds, unlike the LPEH, the MPEH has a different frequency spectrum distribution. More specifically, at *v* = 1.07 m/s, the frequency spectrum of the MPEH output is located at 4.90 Hz, and the voltage amplitude of the MPEH is significantly greater than that of the LPEH. As the moving speed increases to 2.07 m/s, the main peak of voltage shifts to 7.32 Hz, and the response energy mainly distributes over a wide frequency range of 0–10.0 Hz (Figure 9d). It should be noted that a peak appears at the extremely low frequency of 0.25 Hz, corresponding to the motion between the two fixed magnets or to the high-energy orbit. Then, by further increasing the moving speed, the frequency distribution of the voltage is changed, but the peak at the extremely low frequency increases steadily, implying that the response between the two fixed magnets enhances gradually. In particular, at *v* = 2.30 m/s and *v* = 2.50 m/s, from the frequency spectra of the voltage, it is apparent that the vibration energy is concentrated in the range of 0–15 Hz, with a peak of the extremely low frequency, as shown in Figures 9 and 10 (the vibration period of the MPEH is significantly larger than that of the LPEH). Therefore, for the MPEH, since the potential energy is controlled to have an optimal shape, the energy required to produce a large deflection is greatly reduced due to a comparatively strong coupling between the elastic force and the magnetic force. Especially in practice, the moving vehicles consecutively travel on a bridge, so the MPEH will be excited continuously. Then, the

MPEH can keep oscillating in a high-energy orbit and produce a consecutive large output. Thus, the MPEH is suitable for harvesting vibration energy from a practical traffic bridge.

**Figure 11.** Frequency spectra of the output voltage for a vehicle travelling on and leaving the bridge: (**a**) *v* = 1.07 m/s, (**b**) *v* = 1.53 m/s, (**c**) *v* = 1.82 m/s, (**d**) *v* = 2.07 m/s, (**e**) *v* = 2.30 m/s and (**f**) *v* = 2.50 m/s.

#### **6. Conclusions**

In summary, this paper reports a novel concept on how to improve energy harvesting from traffic bridges. The magnets are introduced to tailor the potential energy rather than to create bi-stability. By adjusting the magnets' positions, the ideal mono-stable configuration with an extremely low restoring force and bending stiffness can be obtained. This desirable configuration can make the beam oscillate in a high-energy orbit and gives a large output under weak bridge excitations. The experimental results show that the energy harvesting performance of MPEH is significantly higher than the LPEH when the vehicle speed exceeds 1.82 m/s. For different moving speeds, the MPEH has an extreme peak in frequency spectrum, corresponding to a high-energy orbit. This study may open a new way for energy harvesters designed for traffic bridges. However, some further investigations are still needed to maximize the energy harvesting output. The separation distance and gap distance are two key factors for potential energy, the restoring force, and stiffness. The performance of the MPEHs can be promoted further by optimizing the tip magnet and the fixed magnet.

**Author Contributions:** Conceptualization, Z.Z. and H.Z.; methodology, P.Z.; software, Z.Z.; validation, Z.Z., W.Q. and W.D.; formal analysis, P.W.; investigation, H.Z.; resources, P.Z.; data curation, W.D. and P.W.; writing—original draft preparation, Z.Z., H.Z.; writing—review and editing, W.Q. and P.W.; visualization, P.W.; supervision, P.Z. and P.W.; project administration, Z.Z.; funding acquisition, Z.Z. and P.Z. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was supported by the National Natural Science Foundation of China (Grant No. 52005155), China Postdoctoral Science Foundation (Grant No. 2020M673470), Key Scientific Research Project of Colleges and Universities in Henan Province (Grant No. 20A130001), and Key Research Development and Promotion Project in Henan Province (Grant Nos. 202102310249, 212102310952).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


### *Article* **Study of the Impact of Graphite Orientation and Ion Transport on EDLC Performance**

**Joseph M. Gallet de St Aurin <sup>1</sup> and Jonathan Phillips 2,\***


**Abstract:** A model study of electric double layer capacitor (EDLC)-style capacitors in which the electrodes were composed of low surface area-oriented flakes of graphite that compressed to form a paper-like morphology has suggested that ion transport rates significantly impact EDLC energy and power density. Twelve capacitors were constructed, each using the same model electrode material and the same aqueous NaCl electrolyte, but differing in relative electrode orientation, degree of electrode compression, and presence/absence of an ionic transport salt bridge. All were tested with a galvanostat over a range of discharge currents. Significant differences in energy and power density and estimated series resistance were found as a function of all the factors listed, indicating that capacitor performance is not simply a function of the electrode surface area. This simple postulation was advanced and tested against data: net ion (Na+, Cl<sup>−</sup>) 'velocity' during both charge and discharge significantly impacts capacitive performance.

**Keywords:** dielectric; oriented graphite; electrolyte; EDLC; ion flow; capacitance

**Citation:** de St Aurin, J.M.G.; Phillips, J. Study of the Impact of Graphite Orientation and Ion Transport on EDLC Performance. *Materials* **2022**, *15*, 155. https://doi.org/10.3390/ ma15010155

Academic Editors: Marc Cretin, Sophie Tingry and Carlos Manuel Silva

Received: 12 November 2021 Accepted: 22 December 2021 Published: 26 December 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

#### **1. Introduction**

The development of capacitors with high energy and/or power density will help enable the proposed switch from combustion-based energy systems to 'green' electric systems that are recharged with renewables. Ideally, and theoretically, capacitor energy density could surpass that of batteries, allowing capacitors to replace batteries. Realistically, it is unlikely that capacitors will ever match the energy density of batteries, but the superior power output and durability of capacitors ensures that capacitors will have a number of niche rolls, including load leveling in systems for which batteries are the primary source (e.g., electric vehicles), battery life extension applications such as for satellites, collecting energy from high power surge sources such as regenerative braking, and providing pulsed power for inherently high power applications, as in [1–3].

Given the increased urgency of efforts to create a fully electric power system using only renewable energy sources, resources for research to improve capacitors, particularly electric double layer capacitors (EDLC), also known as Supercapacitors, have increased dramatically in the last decade. In 'carbon only' systems, as per this report, virtually all the increased research has been directed at exploring the effect of the microstructure (e.g., graphene, CNT) and the impact of various additives on carbon energy density, power density, and electric conductivity [1,2,4–12]. There are some general caveats to much of this research. As pointed out elsewhere, the fact that there are no universal protocols for capacitor testing makes comparisons and claims of superior performance difficult to verify [1,13–16]. In addition, other factors, including ionic conductivity, fabrication method, absolute electrode dimensions, even measurement set-up, are likely to impact performance [17]. There has been relatively little work designed to quantify the impact of these factors on carbon-based capacitor performance [13,17,18].

The novelty of the present work is the test of the impact of two overlooked factors on EDLC performance: the orientation of graphite planes in the electrodes, and enhanced pathways for ionic transport. Study of these factors show, for the first time, that they dramatically impact energy storage and power production. Regarding orientation, in the absence of any other modification, the orientation of the graphite flake-based electrodes relative to each other led to a factor of >3× change in energy density. Regarding enhancing ion transport, the use of a 'salt bridge' to enhance ion transport between electrodes increased both energy density and power production significantly. These results, from the simple 'model' system tested, provide experimentally quantitative support for the general postulation that many factors impact EDLC performance.

#### **2. Materials and Methods**

EDLC Microstructure: The electrodes were created from a commercial material, Grafoil, composed of flakes of graphite compressed to create a 'paper-like' material with a very modest surface area [19,20]. In prior work, it was demonstrated that the graphite flake basal planes are strongly oriented parallel to the macroscopic surface of the paper [21].

Two types of Grafoil were employed in capacitor electrodes: (1) Grafoil provided by the manufacture directly; (2) Grafoil mechanically compressed between stainless steel platens to 2000 psi using an Instron SATEC machine (Instron Corp, Norwood, MA, USA).

One simple measure of the impact of compression was the reduction in measured thickness. The measured thickness of the 'raw' material was 0.4 mm, which was reduced to 0.36 mm by compression. To determine the effect of compression on microstructure, both surface area and relative basal plane orientation were determined. A NOVA 4200e Surface Area and Pore Size analyzer (Quantachrome Instruments, Boynton Beach, FL, USA) was used to determine the BET surface area of uncompressed (15.5 m2/g) and compressed (12.7 m2/g), material. Notably, the measured value for the uncompressed material was similar to that reported by the manufacture, 18.5 m2/g, indicating that the measurement technique yields reasonably good quantitative data. In sum, compression at 2000 psi reduces surface area moderately, ~20%.

To test for changes in orientation, anticipated on the basis of earlier studies, relative X-ray diffraction peak heights were determined with a Rigaku MiniFlex 600 benchtop X-ray diffractometer (Rigaku Analytical Devices, Wilmington, MA, USA) (Figure 1). Based on a comparison of intensity ratios between specific crystallographic planes for graphite, it was determined that compression did not modify the relative basal plane alignment of the Grafoil by a statistically significant amount. Both the compressed and uncompressed samples were found to have more than 90% of the graphitic crystals aligned with the macroscopic surface. This is not consistent with earlier work, though might reflect changes in the Grafoil manufacturing process over the decades between studies. That is, the commercial material presently available shows greater alignment of the basal planes and lower surface area than that recorded for the 'as received' material in prior decades.

EDLC Electrode Structure: To test the impact of orientation, electrodes were made of Grafoil that was folded to form an accordion structure, in which each fold was 1 cm across and 3 cm long (Figure 2).

Once the surface incisions were made, the Grafoil was folded along these incisions in an alternating pattern such that the incision formed the outer spine of the fold. The opposite side of each initial incision was folded into each "valley". Each electrode, dry, weighed approximately 1.5 g.

After the accordion-fold was complete, the inside surface of each valley fold was lightly scoured to promote liquid adhesion when the electrolyte was added. Prior to testing of each configuration, two drops of 3% NaCl electrolyte (approximately 0.05 mL/drop) were added to each valley fold on both faces of the electrode. The accordion-fold electrodes were then lightly compressed to close each of the folds.

**Figure 1.** *Compressed and Uncompressed Grafoil XRD*. Both compressed and uncompressed XRD show high orientation of [002] graphite basal planes parallel to surface. For the true powder pattern diffraction pattern arising from ground Grafoil, the [110] (77.19◦ 2θ) and [112] (81.2◦ 2θ) diffraction lines are nearly equal in intensity. In contrast to the powder pattern, for both the compressed and uncompressed materials employed in this study, the [110] peak is nearly invisible, but the [112] peak is quite strong [21]. (Note: Only ~6% of the [002] reflection is shown in order to allow the relative intensity of the other peaks to be clearly shown).

**Figure 2.** *Cont.*

(**C**)

**Figure 2.** *Accordion Electrode Structure*: (**A**) Dimensions shown. For each electrode, a section of Grafoil, 10 cm × 3 cm + a tab at one end, was cut, scored, and folded as shown. (**B**) After initial folding, ~0.01 cm3 of 3% NaCl/DI electrolyte was added to each 'valley'. (**C**) The completed electrode.

Capacitor Construction: Following the construction of two Grafoil accordion electrodes, the next step in capacitor construction was to employ an electrically insulating, commercially woven nylon blend separator (84% nylon and 16% spandex) with a measured thickness of 0.4 cm and 50% open space, employed previously in parallel plate capacitors [22], as an electrically insulating, ion transport 'friendly' media between two, well oriented, electrodes. To enable ion transport, the woven nylon was wetted with 5 drops of the 3 wt% NaCl solution. In prior work with parallel plate capacitors, this electrolyte was found to qualify as an SDM. Indeed, in earlier work [18,22–24], aqueous NaCl solution generally had dielectric constants >10<sup>8</sup> at low frequency (ca. <1 Hz).

The three 'specific orientations' of electrodes that were employed are illustrated in Figure 3. In each configuration, Grafoil surface planes are oriented in a different manner relative to the separator. In the Basal Plane Orientation (BPO), all the macroscopic surfaces of the folded Grafoil are parallel to the separator. Ions diffusing through the separator need to either climb 'up and over' a Grafoil layer, or through one, to reach the next valley. In the Edge Plane Orientation (EPO), the Grafoil surface are perpendicular to the separator. Ions diffusing through the separator could directly transport into all 'valleys' between the Grafoil sheets. In the Offset Orientation (OO), one electrode is arranged per the BPO, and one per the EPO.

**Figure 3.** *Edge View of the Grafoil Sheet Orientation.* Capacitors were constructed using accordion folded Grafoil with three 'specific orientations' relative to the separator, as illustrated.

The next step was to create a 'compressed', mechanically stable, capacitor by using rubber bands to 'attach' the electrodes to a standard sized light microscope slide. In Figure 4, a capacitor in the offset configuration is shown with a 'salt bridge'. Indeed, each capacitor was tested in two configurations: with and without a 'salt bridge'. The net was twelve capacitors built and tested, each with a unique arrangement of the materials, orientations, and 'salt bridge' level: three different orientations × two different compression levels × two salt bridge configurations (Table 1).

**Figure 4.** *Salt Bridge*. Capacitor in offset configuration with molded 'salt bridge' at one end.

The salt bridge was prepared with a gel prepared by mixing 1.8 g of fumed silica (Sigma Aldrich, 0.007 μm avg. particle size, St. Louis, MO, USA) to which a NaCl (saltwater) solution, 0.6 g NaCl and 20 g DI, was slowly added. Similar gels were previously shown to be effective for ion transport in parallel plate-type capacitors [23,24]. Its use as a salt bridge between electrodes requires the same properties. Additionally, the gel held forms such that it was capable of being shaped. As shown, it was 'molded' as a relatively thick layer on one end of the capacitors, creating an avenue for ion transport between the two electrodes.

The final step, to prevent drying, was to place the capacitor, along with a small amount of wetted paper towel, inside a gallon-sized, low density polypropylene bag. Small slits allowed the tabs on the electrodes (Figure 2) to be connected to the galvanostat with alligator clips.


**Table 1.** Summary of Construction and Testing Protocol for 12 Capacitors.

CCC—Constant current charge. CVH—Constant voltage hold. CCD—Constant current discharge. (Illustrated, Figure 4, bottom.).

Measurement Protocol: The test apparatus used in each experiment was the BioLogic VSP-300, a multichannel potentiostat/galvanostat. Each capacitor was studied using the same three step /bidirectional constant current method, a protocol similar to that used to characterize commercial capacitors [25]. Step 1: Charge to a selected positive voltage at constant current. Step 2: Hold capacitor at that voltage for a select period. Step 3: Discharge at a given constant current. A cycle is completed by repeating the Steps 1–3, except Step 1 is done to a negative voltage. Each cycle is repeated six times. Next, the discharge current in Step 3 is changed, and a second set of data collected. In total, six cycles, each at a different discharge current, were measured for each capacitor (Table 1). A typical cycle for one capacitor is shown in its entirety in Figure 5.

As noted earlier, there is no generally accepted protocol for permitting claims of superior performance, or even an absolute comparison of any capacitor properties [1,13–16); however, for this study, that is not a significant concern. This study is not about creating a superior capacitor or comparison with prior published reports, but rather designing a methodology which allows for quantitative comparisons for energy storage and power delivery between capacitors with different orientations and different ion transport paths/mechanisms. The method employed undoubtedly serves this purpose well; moreover, it is consistent with the body of research conducted at the Naval Postgraduate School [1,18,22–24]. Prediction: any and all alternative approached to capacitive property measurement would lead to the same basic conclusions regarding the impact of graphite plane orientation and ion transport significance.

All techniques for measuring capacitance are imperfect [1]. An illustrative example are the inherent limits of impedance spectroscopy. This technique is generally limited at 300 K to study capacitive behavior below 25 mV, as higher temperatures lead to nonlinear behavior. In fact, the fundamental non-linearity of capacitors is reflected in the need to model capacitor behavior with 'equivalent circuits', which are often very complex. The selection of constant current discharge for the present study was made because the results, using this approach, for low frequency discharge are the least convoluted, most

'transparent', available over wide voltage and current ranges, and comparable to a large body of literature [1,13–16,18,22–24].

**Figure 5.** *Example Full Data Set for One Capacitor:* Top: All the data from one capacitor: four discharge currents and six cycles of each (Table 1). Bottom: An expansion of data showing a single cycle. Note: Theoretically, for a capacitor of constant capacitance, the discharge is linear with time. For these capacitors, that was approximately true below ~1.9 Volts. Above that value, voltage dropped sharply, and little energy was delivered to the load.

Three primary values obtained from the data, specific energy (J/kg), specific power density (W/kg), and capacitance below 1 volt (F/g), are reported here. The reported/plotted specific energy density (J/kg) value for each discharge current was the average of energy density values obtained from the twelve discharge curves (six positive voltage, six negative voltage) collected at that current, divided by the total weight of the dry electrodes. The energy density for each individual discharge curve was computed to be the integrated area under the entire discharge curve (V × sec) multiplied by the discharge current. 'Power density' (J/s\* kg), a value appropriate for the quantitative comparison of 'power' between different capacitors, as established in earlier works [18,23,24], was obtained from each discharge curve by dividing the energy of the discharge by the total discharge time and total electrode mass. The capacitance below 1 volt was computed in the standard fashion [1] for constant current discharge, by dividing the current by the near linear value of the discharge curve (dV/dt). The ESR for each CHD sequence was determined using the industry standard method [25]: dividing the voltage drop Δ*V* that occurs within the first 10 ms by the constant current discharge rate *I*. The reported/graphed value at any given discharge current is the average obtained from twelve discharges.

#### **3. Results**

Overall, the data for the model system showed many of the qualitative results expected for capacitors, including: (1) The higher the discharge current, the higher the power density; (2) The higher the electrode surface area/kg, the higher the energy and power density. Three other results are not anticipated by standard theory, and have not been previously reported: (1) The higher the discharge current, the *higher* the energy density; (2) The geometry of the electrodes significantly impacts both energy and power density (3); A 'salt bridge' increases energy and power density. Indeed, the difference in energy density as a function of configuration is virtually eliminated by the salt bridge.

Specific Energy: The specific energies of each of the three configurations with no salt bridge, using either uncompressed or compressed Grafoil, are shown in Figure 6. The electrode configuration clearly impacts performance. BPO configurations resulted in considerably poorer performances than EPO and OO configurations. For example, at a discharge time of 100 s, the uncompressed BPO (~220 J/kg) had a measured specific energy approximately 28% that of the uncompressed EPO (~800 J/kg).

Another clear finding is that, for any particular configuration using compressed Grafoil, electrodes resulted in lower specific energy than its uncompressed counterparts. In the case of the BPO configuration, the difference is very significant. For example, at a discharge time of 1000 s, the compressed BPO electrodes have only about 25 percent the energy density of the uncompressed electrodes. The difference in performance between compressed and uncompressed Grafoil electrodes is significant, but not as dramatic, in the other two configurations. In both the OO and EPO configurations, the uncompressed energy density at a discharge time of 1000 s is <2× larger than that of the compressed electrode configuration.

It is notable that the energy vs. discharge time curves are rather 'flat', and, more significantly, have a 'negative' slope. That is, for EDLC and other capacitors, the energy delivered to the load generally increases with longer discharge times. In this work, the energy densities decreased as the discharge current was decreased, resulting in an increase in discharge time.

In Figure 7, the impact of the salt bridge on specific energy for the uncompressed Grafoil electrodes in the three configurations is shown. It is evident that the salt bridge increases the energy density in all cases. However, the most important finding is that the salt bridge nearly removes the difference in energy density as a function of configuration. With a salt bridge in place, the three configurations show very similar behavior. Indeed, with the salt bridge in place, the BPO and EPO behaviors are virtually identical and the OO performance is only ~25% better. Finally, it should be noted that the 'negative slope' as a function of discharge time is not impacted by a salt bridge.

**Figure 6.** *Summary of Specific Energy for Salt Bridge Free Capacitors.* The energy density data for the six capacitors is shown. One previously unreported trend is apparent: The energy density is a strong function of electrode orientation. It is also clear that compressing/reducing the surface area of the graphitic electrode material reduces specific energy, and energy density decreases slowly with decreasing discharge time.

**Figure 7.** *Impact of Salt Bridge on Specific Energy*. The use of a salt bridge (solid lines) improves performance for all three configurations. More significantly, it 'collapses' the differences in energy density as a function of configuration.

It is reasonable to compare the capacitance (below 1 volt), and, by inference, energy density, of these electrodes with other pure carbon materials. In earlier studies, a graph of capacitance, as a function of the carbon surface area, demonstrated a linear relationship [2,26]. The slope of the line can be employed to predict capacitance on the basis of the measured surface area. For uncompressed Grafoil with a measured surface area of 15.5 m2/g, this suggests a capacitance of 0.93 F/g. As shown in Figure 8, this value is about a factor of two better than that observed in the uncompressed material in the OO configuration. The figure also shows that comparisons of 'capacitance' are fraught with confusion, as so many parameters, such as voltage range, over which capacitance is measured, measurement frequency, ion identity, and details of construction, which are not simply carbon surface area, impact this value. In sum, the results of this study are qualitatively consistent with the anticipated capacitance based on earlier correlations. That is, the energy density per unit area of the electrodes is in close quantitative agreement with earlier findings.

**Figure 8.** *Capacitance as a Function of Discharge Time*. Consistently, the presence of a salt bridge (solid lines) improves capacitance, all other factors were unchanged.

Specific Power: The specific power (W/kg) for the six capacitors, made with the uncompressed Grafoil, all three configurations with and without salt bridge, are shown (Figure 9). For the three capacitors without a salt bridge, as with the energy density results, the power density is clearly a function of the configuration. The BPO configuration without a salt bridge delivers the least power. For example, for a 1000 s discharge, it would deliver ~2.5 × <sup>10</sup>−<sup>1</sup> W/kg, which is about 30% of the power delivery (~8.0 × <sup>10</sup>−<sup>1</sup> W/kg) for the uncompressed EPO configuration.

As with the case for energy density, the power density differences collapse for the capacitors built with salt bridges. All three configurations with a salt bridge have very similar power density curves as a function of discharge time. Finally, it is notable that the power density curves are of standard slope. That is, as the discharge times decrease, and the delivered power increases.

**Figure 9.** *Specific Power Increases with Higher Current/Lower Discharge Time*. Configuration impacts power significantly in the absence of a salt bridge, but the specific power value differences are almost independent of configuration when a salt bridge (solid lines) is added.

Estimated Series Resistance: The finding that the energy and power density values of all the capacitors studied herein increases with shorter discharges may reflect the unusual ESR behavior. As shown in Figure 10, the ESR values are not constant, but rather a function of the discharge time/current. For all the capacitors, compressed and uncompressed, with and without a salt bridge, ESR values are higher for smaller currents/longer discharges. Given that the load and the output resistance form a voltage divider, increased output resistance will concomitantly reduce the energy drop across the load. Hence, the ESR trend with the discharge time is both consistent with, and explains, the observed drop in the delivered energy at longer discharge times.

The ESR data show clear trend lines for any particular capacitor, but certainly do not suggest a clear quantifiable formula linking ESR to orientation, the state of compression, or presence of a salt bridge. Table 2 does suggest some weak qualitative correlations. First, for any given orientation, compression increases ESR, and second a salt bridge reduces ESR. It is also clear that there is not a strong orientation dependence. For example, all three configurations are the same with compressed/no salt bridge electrodes within experimental error. Uncompressed/salt bridge capacitors are all essentially identical.

Ragone Plots: Presenting the data in Ragone format helps illustrate the unique behavior of all these capacitors (Figure 11). In particular, plotting the data in the standard Ragone format shows that the 'slopes' are inverted. In general, Ragone charts show that, for both batteries and capacitors, delivering more power comes at the expense of reducing the energy delivered. Thus, the lines of conventional capacitors and batteries, plotted as per Figure 11, show a negative slope. For all capacitors studied herein, there is no trade-off. Both energy and power increase as the discharge time is reduced, yielding positive slopes.

**Figure 10.** *ESR of all capacitors as a Function of Discharge Rate.* It is notable that ESR is a strong function of discharge time for all capacitors studied.


**Table 2.** ESR as a function of orientation, compression, and salt bridge.

The power delivered curves also show a remarkably high slope. This reflects the fact that, unlike standard capacitive behavior, the energy delivered increases as the discharge time is reduced. The power delivered is the ratio of the energy delivered to the time of delivery. Conventionally, the numerator term decreases with shorter discharge time. Still, typically, this value does not decrease as rapidly as the denominator, hence the power increases with the decreasing discharge time/higher current. For the capacitors studied herein, the numerator increased, and the denominator decreased, leading to a sharp slope.

**Figure 11.** *Ragone Plot*. These plots show the same configuration and salt bridge impacts observed in the specific energy and power plots. Also notable is that the plots are all positive slopes, whereas, in general, for capacitors, the slopes on a Ragone chart are negative.

#### **4. Discussion**

This study was conducted to determine, on a fundamental level, whether either or both orientation and ion transport 'velocity' can impact the performance of EDLC capacitors. In order to minimize the number of complicating factors, the electrodes were created with low surface area, pure graphite flakes, strongly oriented with basal planes parallel to the surface.

The study, as intended, produced clear empirical findings regarding the target questions of the impact of electrode orientation and the potential impact of ion transport. Regarding electrode orientation: As shown in Figure 6, the orientation can change the energy density by a factor of more than 3×, and, as shown in Figure 8, the orientation can impact the power density by more than a factor of 2. Regarding ion transport: In all cases, the use of a salt bridge increased performance, although the magnitude varied as a function of electrode orientation. Capacitors with electrodes in orientations that yielded higher energy and power density even without salt bridges were marginally improved. The orientation with the lowest energy and power density without a salt bridge, BPO, improved significantly. In net, using a salt bridge effectively made all electrode orientations roughly equivalent performers. Postulate: Ion transport in EDLC can be performance determining. Note: It is understood that improved ion transport only improves performance to a finite limit.

Other data supports the suggestion that ion transport, and not always electron transport, can be performance limiting. For example, it was found that compression of the Grafoil led to a 25% reduction in surface area, but far larger fractional reductions in energy and power density in all cases. Indeed, the reduction was as great as an order of magnitude decrease in these values. Yet, compression, logically, should increase electron transport by enhancing contact between graphitic plates and concomitantly reduce ion transport by 'shrinking' or eliminating channels for ion transport. Thus, the large reduction in performance to Grafoil compression can logically be attributed to reductions in ion transport, but not to a reduction in the electron conductivity of the electrodes.

Can the velocity of ion transport somehow explain the inverted slope for the energy vs. discharge time curves (Figures 6, 8 and 10)? Postulate: The energy and voltage of the charged species, electrons, on the electrodes are higher if more 'ionic' dipoles (Na+ or Cl−) are present in the electric double layer. This concept is a variation on the recently postulated Theory of Superdielectric Materials (T-SDM), as discussed elsewhere [27–29]. Thus, given a constant charging period, a system with faster ion transport will allow more ions to travel and 'add' to the electric double layer than a system with lower ion transport rates. Conversely, this same ion transport 'advantage' can become disadvantageous the greater the discharge time. For lower discharge currents/longer discharge times, more ions 'retreat' and neutralize (e.g., Na+ + Cl<sup>−</sup> => NaCl). This reduces the number of electric dipoles in the boundary layer, and hence reduces the energy and voltage of the charges remaining on the electrodes. This implies that, the faster the electrons are removed, the fewer ions, which clearly move much more slowly than electrons, will neutralize. Consequently, rapid discharge should lead to more net energy falling on the load, as observed.

In order to gain additional insight into the impact of various factors on 'transport', the ESR was measured for all capacitors. The impact of a salt bridge on the ESR, an addition to the capacitor which should only impact ion transport, is revealing. In all cases the salt bridge reduced the ESR value. This trend raises this question: why should ion transport change net resistance? Is not net resistance a function of electron transport? A postulated answer: Ion and electron transport are 'coupled'. If ion transport is enhanced, so too is electron transport. The physics behind this proposed coupling is not obvious, or is at least 'complicated', and was not be considered here.

#### **5. Conclusions**

In conclusion, the key empirical finding is that, in carbon-based EDLC, enhanced ion transport in the electrolytic material improves performance. Increasing ion transport rates via the proper orientation of electrodes, the inclusion of a salt bridge, or an increased porosity also increases energy and power density. Moreover, it appears that ion and electron transport are linked. If ion transport is enhanced, electron transport is as well. The empirical findings are certain, and a simplistic model of improved performance linking to ion transport enhancement appears to be at least consistent with all observations. Yet, the underlying physics is not clear. Why are ion transport and electron transport coupled? As the ions retreat/neutralize, does this reduce the energy of the charges remaining on the electrodes?

**Author Contributions:** Concepetulalization: J.P.; methodology, J.P. and J.M.G.d.S.A.; software J.M.G.d.S.A.; validation, J.P. and J.M.G.d.S.A.; formal analysis, J.P. and J.M.G.d.S.A.; investigation, J.M.G.d.S.A.; resources, J.P.; data curation, J.M.G.d.S.A.; writing—original draft preparation, J.M.G.d.S.A.; writing—review and editing, J.P.; visualization, J.M.G.d.S.A.; supervision, J.P. project administration, J.P.; funding acquisition, J.P. and J.M.G.d.S.A. All authors have read and agreed to the published version of the manuscript.

**Funding:** No external funding supported this work.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** A great deal of the source data can be found here: Joseph Gallet de St. Aurin, M.S. Thesis: FOLDED CARBON ELECTRODES: A NOVEL APPROACH TO ENHANCING SUPERCAPACITOR PERFORMANCE Naval Postgraduate School. 2021.

**Acknowledgments:** Claudia C. Luhrs, MAE Dept., Naval Postgraduate School, assisted with XRD and BET measurements.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

