With the development of electronic devices towards miniaturization and light weight, electrical capacitors with high recoverable energy storage density and utilization are eagerly desired.^[1–3] Currently, commercial high-power capacitors with energy density about 2.0 J/cm³ are mostly made of linear dielectric polymer materials with low dielectric constant, e.g., polypropylene, polyethylene terephthalate, and polyphenylene sulfide, because of their high dielectric breakdown strengths.^[2] Recently, antiferroelectric materials, in which dipole moments are alternatively aligned in the opposite directions, resulting in the zero net polarization, have become the promising candidates for application in high energy storage due to an electric field forced phase transition from antiferroelectric into the ferroelectric states. Compared with the commercial capacitor, the electrical capacitor made of antiferroelectric material may achieve a higher energy density (more than 10.0 J/cm³)^[4,5] and fast discharge speed (less than 20 ns).^[6] Among various antiferroelectric materials, PbZrO₃ based antiferroelectric materials have received extensive attention, owing to their possible energy storage densities as high as 50 J/cm³.^[7] A lot of effort has been made to realize the theoretical prediction of energy storage density in the PbZrO₃ related antiferroelectric material by doping elements,^[8,9] local field engineering^[10] or changing electrode material.^[11,12] For example, Parui and Krupanidhi obtained a remarkable energy storage density of 14.9 J/cm³ at 600 kV/cm in La-doped PbZrO₃ thin film.^[8] Sa et al., reported a recoverable energy density as high as 17.4 J/cm³ in the α-Fe₂O₃ nanoparticles doped PbZrO₃ thin film at 600 kV/cm.^[10] Ge et al. found that the energy storage density and the energy efficiency of Pb_0.97La_0.02(Zr_0.95Ti_0.05)O₃ film can be, respectively, enhanced to 14.8 J/cm³ at 1 MV/cm and 66% by using oxide metal LaNiO₃ as top electrodes instead of Pt.^[11]

On the other hand, the antiferroelectric-to-ferroelectric phase transition occurring in the antiferroelectric material is usually accompanied by a large volume change.^[13] An external strain/stress can also effectively change the critical field of phase transition. Many researchers have investigated the influences of applied strain/stress on properties of the ferroelectric films.^[14–18] However, only a few works have systematically studied this issue in antiferroelectric bulk or films.^[19–21] Tan et al. found that the critical electric fields, i.e., E_AF for the antiferroelectric-to-ferroelectric phase transition and E_FA for the reverse transition, are shifted to high values by exerting the uniaxial or radial compressive stresses on the ceramic composition of Pb_0.99Nb_0.02[(Zr_0.57Sn_0.43)_0.94Ti_0.06]_0.98O₃,^[20] which leads to the increase of energy storage density in the antiferroelectric capacitor.^[22] Ayyub et al. found that PbZrO₃ film presents ferroelectricity when its thickness is below 400 nm and show antiferroelectric behavior for the thickness above 550 nm due to the internal stresses caused by the substrate-clamping effect and their relaxation with thickness, which may result in the different energy storage densities and efficiencies for PbZrO₃ films at various thicknesses.^[19] After removing the stress by the substrate clamping, the energy storage properties of freestanding film may be changed. So far, a lot of freestanding ferroelectric films have been prepared by different methods.^[23–25] In this paper, we fabricated antiferroelectric Pb(Zr_0.95Ti_0.05)O₃(PZT(95/5)) films on both Pt foils and platinized Si substrates by using pulsed laser deposition. In comparison with those of the substrate-clamped reference sample (C-PZT(95/5)), the E_AF and E_FA of freestanding film (F-PZT(95/5)) are remarkably enhanced while its saturation polarization is slightly reduced. As a result, the energy storage density and the energy storage efficiency of the F-PZT(95/5) film are both significantly enhanced.

In our experiment, 200-nm-thick Pt foils, which were used as both the flexible substrates and the bottom electrodes, were obtained by etching platinized Si wafers (Pt(200 nm)/Ti(50 nm)/SiO₂(500 nm)/Si) in 10 wt% HF solutions for 4 h. Due to the dissolution of the Ti layer and SiO₂ layer in HF solutions, the thin Pt foils were detached from the wafers. The Pt foil was then transferred to a thermal oxide silicon wafer. The detailed procedures for preparing Pt foils were described in our previous publication.^[17] The 300-nm-thick antiferroelectric Pb(Zr_0.95Ti_0.05)O₃ films were grown on the Pt foils by pulsed laser deposition at 560 °C and 10-Pa oxygen pressure, with a KrF excimer laser (248 nm, Coherent Inc.) running at a repetition rate of 5 Hz and an energy density of 2.3 J/cm². In order to characterize the electrical behaviors of antiferroelectric film, an array of circular Cu top electrodes (100 nm in thickness and 100 μm in diameter) was deposited on a PZT(95/5) layer with a metal shadow mask by e-beam evaporation. For comparison, 300-nm-thick C-PZT(95/5) film grown on platinized Si substrate was also prepared using the same parameters. The crystalline structures were characterized by an x-ray diffractometer (D8 Advance, Bruker) with Cu K_α radiation. The ferroelectric properties of F-PZT(95/5) and C-PZT(95/5) films were measured by a standardized ferroelectric test system (Precision Premier II, Radiant Technologies). Direct current (DC) electric field dependence of dielectric properties of the antiferroelectric films were measured by using a Keithley 4200 parameter analyzer.

Figure 1(a) shows the XRD patterns of both F-PZT(95/5) and C-PZT(95/5) films. The XRD results illustrate that both samples possess the pure polycrystalline tetragonal PZT(95/5) phase. The splitting diffraction peaks of PZT(95/5)(001)/(100) and PZT(95/5) (002)/(200) appear around 22° and 44° for both films, respectively, as shown in Figs. 1(b) and 1(c). Similar observations have also been reported in PbZrO₃ bulk ceramic and Pb_0.99Nb_0.02(Zr_0.85Sn_0.13Ti_0.02)_0.98O₃ films, which is ascribed to the tetragonal distortion of the orthorhombic lattice of antiferroelectrics.^[26–28] The intensities of (100) and (200) diffraction peaks for F-PZT(95/5) films are much higher than those of (001) and (002) diffraction peaks, but for the C-PZT(95/5) films the intensities of (100) and (200) peaks are lower than those of (001) and (002) peaks. The variations in the intensity of diffraction peaks suggest that the C-PZT(95/5) films are c-axis-preferred oriented, but after removing the rigid substrates the F-PZT(95/5) films change to the preferred orientation of a axis. As displayed in the inset of Fig. 1(a), the (111) diffraction peak of F-PZT(95/5) film shifts about 0.12° toward the high angle in comparison with that of C-PZT(95/5) film, which suggests that the C-PZT(95/5) film suffers an out-of-plane tensile strain of 0.27% and an in-plane compressive strain of −0.13%. However, the strain caused by the substrate clamping effect is relaxed in the F-PZT(95/5) film.

Fig. 1.

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	Fig. 1. (a) X-ray diffraction patterns of both free-standing F-PZT(95/5) film and substrate-clamped C-PZT(95/5) film. The inset shows a magnified view of PZT(95/5) (111) peak. The magnified views of (b) PZT(95/5) (001) and (100) peaks, and (c) PZT(95/5) (002) and (200) peaks.

Such a change in structure of PZT(95/5) thin film should have a significant influence on its electrical properties. To bridge the structural and electrical properties in the antiferroelectric materials, polarization hysteresis loops of F-PZT(95/5) and C-PZT(95/5) films are measured at 10 kHz and at room temperature as shown in Fig. 2. Both samples exhibit the distinct features of double P–E hysteresis loops and small remanent polarizations (2.7 μC/cm² for F-PZT(95/5) and 1.7 μC/cm² for C-PZT(95/5)), which indicates that both films are antiferroelectric at room temperature. For the C-PZT(95/5) film, the antiferroelectric-to-ferroelectric phase transition occurs at E_AF = 360 kV/cm and the ferroelectric-to-antiferroelectric phase transition appears at E_FA = 160 kV/cm, which are determined by the peak value of the differentiation of the P–E loop. After removing the substrates, the E_AF and E_FA for F-PZT(95/5) films change to 770 kV/cm and 410 kV/cm, respectively, which are significantly increased as compared with those of C-PZT(95/5) film. The saturation polarizations of both F-PZT(95/5) and C-PZT(95/5) films are 31.3 μC/cm² and 32.9 μC/cm², respectively. It should be noted that the breakdown electric field of C-PZT(95/5) film is much lower than that of F-PZT(95/5) film. Thus the different maximum electric fields are used for the electrical characterizations of C-PZT(95/5) and F-PZT(95/5) films. The changes in the E_AF and E_FA can be understood by taking into account the rotation of dipole moments under the external stress. For antiferroelectric C-PZT(95/5) film which suffers an in-plane compressive stress caused by the rigid substrate, the c axis, i.e., the ⟨001⟩ direction of PZT(95/5), is easy to align out of plane.^[29] During the antiferroelectric-to-ferroelectric phase transition, the c axis of the PZT(95/5) unit cell expands while the a-axis shrinks.^[30,31] Therefore, the relaxation of in-plane compressive stress/strain in F-PZT(95/5) films acts against the antiferroelectric-to-ferroelectric phase transition. As a result, after removing the substrates, the E_AF and E_FA of F-PZT(95/5) films become larger than those of C-PZT(95/5) films.

Fig. 2.

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	Fig. 2. Polarization hysteresis loops of antiferromagnetic F-PZT(95/5) and C-PZT(95/5) film.

As is well known, the electric displacement reflects the dipole orientation, while the dielectric constant indicates the motion of these dipoles. For a linear dielectric material, the dielectric constant can be related to the electric polarization as follows:

Fig. 3.

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	Fig. 3. Electric field dependences of (∂P/∂E)/ε0 for (a) C-PZT(95/5) and (b) F-PZT(95/5) film, obtained by polarization hysteresis loops. Electric field dependences of dielectric constant for (c) C-PZT(95/5) and (d) F-PZT(95/5) film, obtained by the capacitance–electric field measurement.

The energy storage density W of an antiferroelectric material can be obtained from the hysteresis P–E loop by using the relation

Fig. 4.

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	Fig. 4. Variations of (a) maximum polarization Pmax and (b) the calculated energy storage density with electric field of F-PZT(95/5) and C-PZT(95/5) film, measured at different maximum electric fields.

As shown in Table 1, when a saturation electric field of 500 kV/cm is applied to the C-PZT(95/5) film, the total energy, energy storage and energy loss are 11.9 J/cm³, 6.49 J/cm³, and 5.41 J/cm³, respectively. In contrast, when a saturation electric field of 1300 kV/cm is applied to the F-PZT(95/5) film, the total energy, energy storage and energy loss are 25.67 J/cm³, 17.45 J/cm³, and 8.22 J/cm³, respectively. Compared with those of the C-PZT(95/5) film, the energy storage density and the energy loss density of F-PZT(95/5) film are enhanced by 10.96 J/cm³ and 2.81 J/cm³, respectively. Consequently, the energy efficiency of F-PZT(95/5) film reaches 67.9%. The enhanced energy storage properties of F-PZT(95/5) film after the relaxation of stress can be understood from the displacements of Zr⁴⁺ and Ti⁴⁺ atoms. In an ABO₃ perovskite structure, the ferroelectric polarization originates from the displacements of B⁴⁺ ions from the equilibrium B sites. For antiferroelectric PZT(95/5) film, Pb²⁺ ions occupy the A sites and Zr⁴⁺ and Ti⁴⁺ ions occupy the B sites. During antiferroelectric-to-ferroelectric transition, the lattice parameter a decreases while c increases. The XRD patterns show that C-PZT(95/5) film is c-axis-preferred oriented, but the F-PZT(95/5) films change to the preferred orientation of a axis. An in-plane compressive strain exists in the C-PZT(95/5) film, which benefits the antiferroelectric-to-ferroelectric transition. After relaxation of the in-plane compressive stress in F-PZT(95/5) film, the coulomb interaction between two O²⁻ ions at the opposite vertices of octahedral becomes stronger, which causes the Pb²⁺ ions at A site to be hard to displace from the ab plane to ⟨111⟩ direction during the antiferroelectric-to-ferroelectric phase transition. As a result, compared with in C-PZT(95/5) film, a higher electric field is needed to induce the antiferroelectric-to-ferroelectric phase transition in F-PZT(95/5) film, which therefore increases both the energy storage density and the energy storage efficiency of the F-PZT(95/5) film after removing the substrate and relaxing the substrate clamping.

Table 1.

Table 1.

Table 1. Energy storage properties of both C-PZT(95/5) and F-PZT(95/5) membranes. . Samples Total energy density/(J/cm3) Energy storage density/(J/cm3) Energy loss density/(J/cm3) Energy Efficiency/% C-PZT(95/5) 11.9 6.49 5.41 54.5 F-PZT(95/5) 25.67 17.45 8.22 67.9

Table 1. Energy storage properties of both C-PZT(95/5) and F-PZT(95/5) membranes. .

In this work, 300-nm-thick free-standing antiferroelectric PZT(95/5) film is successfully prepared on 200-nm-thick Pt foils by the pulsed laser deposition. The intensities of the (100) and (200) diffraction peaks of the F-PZT(95/5) film are increased after removing the substrate, indicating that F-PZT(95/5) film possesses an a-axis-preferred orientation. The antiferroelectric-to-ferroelectric electric field of the F-PZT(95/5) film increases up to 770 kV/cm compared with that of the substrate-clamped C-PZT(95/5) film. Therefore, the energy storage density and the energy storage efficiency are enhanced to 17.45 J/cm³ and 67.9%, respectively. Using the wafer bonding technology, the free-standing C-PZT(95/5) films can be integrated on any kind of flexible substrates, making them available for applications in flexible electrical devices such as capacitors and actuators.