The composition and crystal lattice of the (1-x) BaBi₄Ti₄O₁₅—x Bi₄Ti₃O₁₂ samples with a certain sintering temperature of 1020°C, the X-ray diffraction test are exhibited in Figure 1, x = 0–1. The results show that all diffraction peaks are in perfect agreement with the orthorhombic structure of BBT-BiT phase of intergrowth bismuth layered structure (BLSFs) with space group P2₁am and no obvious second phase (Kobayashi et al., 2004; Wang W et al., 2008). Figure 1B shows the variation in the XRD peaks of BBT (119), BBT-BiT (118) and BiT (117) at 2θ = 30.2–30.6°, the dashed lines represent the peaks of Lorentzian profile function fitting. BBT with space group A2₁am (m = 4, orthorhombic, JCPDS No. 35–0757) has the strongest peak (119), and BiT with space group B2cb (m = 3, monoclinic, JCPDS No. 72–1019) has the strongest peak (117). Although the position of X-ray diffraction peak shifts with the change of BiT content, the strongest peaks (118) of BBT-BiT (m ∼ 3.5, JCPDS No. 42–0053) are consistent with the crystal plane index (112m + 1) of the strongest peak in intergrowth BLSFs ceramics (Wang et al., 2010; Badapanda et al., 2019; Jiang et al., 2020a), which lie between the strongest peak (119) BBT and the strongest peak (117) of BiT. In addition, the peaks of BBT-BiT with the same Angle as that of BBT and BiT, but the peak intensity of BBT-BiT is lower than that of BBT and BiT. This indicates that the crystallinity of BBT-BiT is lower than that of BBT and BiT ceramics, and the conductivity of BBT- BiT ceramics is lower than that of BBT and BiT ceramics, which will be proved in the subsequent analysis.

Figure 2 shows the SEM images and grain sizes distribution on fracture surfaces of (1-x) BBT—x BiT ceramics with x = 0 (a), 0.2 (b), 0.4 (c), 0.6 (d), 0.8 (e) and 1 (f). Due to the high grain growth rate perpendicular to the c-axis during grain growth of bismuth layered compounds, all samples showed the typical plate-like grains typical random lamellar grain orientation and obvious anisotropy in grain structure (Suarez et al., 2001; Diao et al., 2013). The images also exhibit that most pores on the grain boundary rather than internal grain, which indicate that ceramics sintering is relatively dense. At the same time, the grain size of BiT ceramics is extremely larger than that of BBT ceramics, while BBT-BiT samples forming intergrowth structure that the two type of perovskite layers are interlaced and symbiotic along the c-axis (Luo et al., 2001). In addition, the SEM images also show that the BBT-BiT intergrowth ceramics have more holes and their grains growth are not sufficiently compared with those of the single BBT and single BiT ceramics.

Figure 3A exhibits the evolutions in bulk density (ρ _b) as x value increase in (1-x) BBT—x BiT ceramics. The bulk density (ρ _b) of (1-x) BBT—x BiT ceramics increases with the x increasing, it may be caused by the differences in BBT and BiT ceramics systems. Figure 3B shows the average grain size of (1-x) BBT—x BiT ceramics. As mentioned in the XRD analysis that BBT-BiT ceramics have low crystallinity, which is usually shown as small grain size in SEM images. The average grain sizes of (1-x) BBT—x BiT (x = 0.2–0.8) ceramics are lower than those of BBT and BiT ceramics, combined with SEM images, these results indicate that the diffusion mechanism of BBT-BiT intergrowth ceramics play a role in the sintering process and inhibit the grain growth of the ceramic samples.

Figure 4 shows the TEM images and schematic crystal structure of the 0.8BBT-0.2BiT ceramics. Figure 4A exhibits its cross-sectional TEM image, Figures 4B, D are the high-resolution cross-sectional TEM image. Figure 4C is the FFT image of Figure 4B, which shows the that both BBT and BiT phases coexist in the sample. Figures 4E, F are partial larger version of Figures 4D, F is a figure with increased contrast after color adjustment of Figure 4E, and Figure 4G is a partial larger version of Figure 4E. Figure 4F shows an irregular layered structure with alternating layers of three and four layers, and the layer spacing values are very close to half the length of the c-axis of three layers (3.3 nm) or four layers (4.2 nm). BBT ceramics are composed of bismuth oxygen layer (Bi₂O₃)²⁺ and four-layer pseudo-perovskite blocks [(Ba_1/3Bi_2/3)₃Ti₄O₁₃]²⁻ regularly alternating, while BiT ceramics are composed of bismuth oxygen layer (Bi₂O₃)²⁺ and three-layer pseudo-perovskite blocks (Bi₂Ti₃O₁₀)²⁻ regularly alternating. Combine Figures 4F, G, the schematic crystal structure of 0.8BBT-0.2BiT ceramics are formed by irregularly alternate arrangement of bismuth oxygen layer (Bi₂O₃)²⁺ and pseudo-perovskite blocks of three layers (Bi₂Ti₃O₁₀)²⁻ or four layers [(Ba_1/3Bi_2/3)₃Ti₄O₁₃]²⁻.

Figure 5 shows the temperature dependence of the dielectric constant ε _r for all samples from 30°C to 700°C at five frequencies as color online. It can be seen from ε _r—T curves of all samples at these frequencies, the dielectric constants decrease with increasing frequency (that is phenomenon of frequency dispersion) and the peak of dielectric constants occur near Curie temperature T _c. The T ₁ temperature of BiT ceramics in Figure 5F indicates that the sample has undergone phase transformation in the vicinity of this temperature. The flat dielectric constant in relation to temperature of relaxation ferroelectrics is due to the presence of polar nanoregions on the structure, which makes them in ergodic relaxor state (Ni et al., 2013). Meanwhile, the values of Curie temperature change at different frequencies and T _c increase with increasing frequency generally, except in Figures 5B, C. Due to the influence of space charge and conductivity, the dielectric constant tends to increase at low frequency and high temperature (Wang et al., 2022), so 100 kHz is selected as a frequency for subsequent experimental analysis.

Figure 6 exhibits the curves of Curie temperature T _c and piezoelectric constant d ₃₃ of (1-x) BBT—x BiT ceramics with different BiT content (x), respectively. As the T _c—x curve at 100 kHz shown in Figure 6A, the Curie temperatures T _c (458–556°C) of BBT-BiT intergrowth ceramics are between BBT (408°C) and BiT (674°C), which increase with the increase of BiT content x. Lattice distortion is the main influencing factor of Curie temperature (T _c) of bismuth layered ceramics, which can be expressed by tolerance factor (t), as shown in law (Eq. 1) (Suarez et al., 2001). t = R A + R O 2 R A + R O (1) where, R _A, R _A and R _O are the ionic radii of A-site, B-site and O^2- respectively. The smaller the tolerance factor t, the higher the Curie temperature T _c of the ceramics. At this work, the general formula is (Bi₂O₂)²⁺(A_m-1B_mO_3m+1)²⁻ for bismuth-layered structure ferroelectrics (BLSFs), the B-site is occupied by Ti⁴⁺ with an ionic radius of 0.605 Å, while the A-site is Bi³⁺ with an ionic radius of 1.17Å for Bi₄Ti₃O₁₂ (BiT) ceramics and A-site is (Ba2+1/3Bi3+2/3) for BaBi₄Ti₄O₁₅ (BBT) ceramics with an ionic radius of 1.32 Å. Thus, the radius of A site ion for (1-x) BaBi₄Ti₄O₁₅ - x Bi₄Ti₃O₁₂ intergrowth ceramics decreases with the increase of BiT content (x). With the increase of x, the increase of radius of A-site leads to decrease t, so T _c gradually increases, indicating that the BBT-BiT intergrowth structure of BLSFs increase Curie temperature compared with single BBT ceramics.

As shown in Figure 6B, the piezoelectric constants d ₃₃ increase first and then decrease with the increase of x of BBT-BiT intergrowth ceramics, show an opposite trend in changes, compared with that of T _c. The intergrowth structure of BLSFs with m/m+1 layers are composed of m layers and m+1 layers, two-phase lattice distortion occurs due to the mismatch of lattice parameters, resulting in the formation of new intergrowth structural phases to match the lattice (Yang et al., 2018). The intergrowth structure, which affects the symmetry of ceramic c-axis direction and the properties of material. And the intergrowth structure construction also effects oxygen vacancy concentration, which will affect the deflection of ferroelectric domains, thus affecting the piezoelectric properties of BBT-BiT ceramics. At the same time, the maximum value of d ₃₃ for BBT-BiT ceramics is obtained at x = 0.2, the piezoelectric constant d ₃₃ is 16 pC/N and its corresponding Curie temperature T _c is 458°C. Combined with the SEM results, it can be found that the BBT-BiT intergrowth ceramic (especially 0.8BBT-0.2BiT sample) are with smaller grain size and more pores than the BBT and BiT samples. Therefore, the oxygen vacancy concentration can be controlled through many methods, such as ions doping and vacuum sintering, to further improve the grain growth and electrical properties.

Figure 7 shows the XPS survey spectra and O 1s scan of (1-x) BBT—x BiT ceramics. As shown in Figure 7A, the resolution spectrum of Bi 3 days has 3 days 3/2 and 3 days 5/2 doublets in the binding energy range of 800–770 eV, and the change of peak area reflects the decrease of Ba 3 days content with the increase of BiT content (x). Bi 4f, Bi3d and Ti 2p spectra increase first and then decrease in the peak areas as shown in Figures 7B, C. For (1-x) BBT - x BiT intergrowth ceramics (x = 0.2–0.8), the binding energy of Bi 4f, Bi3d and Ti 2p shift not significantly with the x increasing, compared with those of BBT and BiT ceramics, which also shows that the intergrowth bismuth layered ceramics did not generate other impurity phases.

To further confirm the sources of oxygen vacancy in the samples, the peaks of O 1s scan are fitted into three peaks as shown in Figure 7D. Take BBT ceramics for example, the Lattice oxygen ions (O_Lattice) are represented as O 1s (I) peak (529.29 eV), the None-lattice oxygen ions (O_Non-lattice) are represented as O 1s (II) peak (530.38 eV) induced by oxygen deficiency from deviation of BaBi₄Ti₄O₁₅ stoichiometric ratio in oxygen defect region and the O 1s (III) peak (531.68 eV) generated by the organic oxygen from the circumstance of the samples’ surface (Yoo et al., 2012; Jaim et al., 2017; Reddy et al., 2019). Table 1 shows the values of the integral Area and Area percentage content of oxygen peaks from the above different sources, and the ratio of O_Non-lattice/O_Lattice.

Bismuth volatilizes easily at a high sintering temperature of 1000–1100°C for BLSFs, resulting in forming Bi vacancies ( V B I ‴ ) and oxygen vacancies ( V O • • ) that can be expressed by Kröger–Vink formula (2) (Qiao et al., 2016): 2 B i B i × + 3 O O × ⇔ B i 2 O 3 g a s + 2 V B I ‴ + 3 V O • • (2) where B i B i ×   a n d   O O × present electrically neutral vacancy of Bi and O, respectively. According to the O_Non-lattice/O_Lattice values of the (1 − x) BBT − x BiT ceramics in Table 1, there are more oxygen vacancy defects in single BBT (2.75) and BiT (2.49) BLSFs, indicating that the intergrowth structure of BLSFs can inhibit the generation of oxygen vacancies. Oxygen vacancies promote ions to diffusion and transfer during sintering, improving the sintering speed of ceramics to promote the growth of grains (Tewari et al., 2014; Tang et al., 2020), which are also consistent with the SEM images (Figure 2) and average grain size (Figure 3B) results.

Figure 8 shows Z′-Z″ complex impedance spectra collected under various temperatures with the frequency of 100 Hz–10 MHz, the impedance spectra selected the temperature range of the impedance circle that can be obtained by the sample. These Z′-Z″ spectra are fitted with the equivalent circuit as shown in Figure 8A, which can be composed of approximately two superimposed semicircles, reflecting the impedance contribution of the grain (high-frequencies arc) and grain boundary (low-frequencies arc) to the materials (Dhak et al., 2008). Figures 8A–E show that the Z′-Z″ spectra of ceramics are approximately single semicircles, which mainly represent the dominant role of grain boundary characteristics, and the interaction between grain, grain boundary and ceramic electrode interface (Lin and Chen, 2009). Moreover, the BiT ceramic exhibits grain and grain boundary characteristics as shown in Figure 8F.

For all samples, the radius of curvature decreases with increasing temperature for the Z′-Z″ complex impedance spectra, indicating that the resistivity of the sample has a negative temperature coefficient (Fujii et al., 2002). Meanwhile, Z′ decreases gradually with the increase of BiT content at the same temperature. Asymmetric in all characteristic wide peaks at measured temperatures for all samples indicates non-Debye type relaxation processes (Pradhani et al., 2018), and the values of resistance, capacitance and relaxation time for the bulk and grain boundary was listed in Table 2.

To confirm the relaxation mechanisms of the ceramics, relaxation frequencies (ω _max (T)) in relation to temperature ln ω - 1/T are plotted in Figure 9. The relaxation activation energy (E _relax) is calculated by Arrhenius equation according to the slopes obtained from the linear fitting of the plots, this law (Eq. 3) as follow (Qiao et al., 2016): ω max T = ω 0 exp − E relax k B T (3) where T is the temperature in Kelvin scale, ω ₀ is the pre-exponential factor, and k _B is the Boltzmann constant.

Figure 9 shows ln ω - 1/T relaxation frequency Arrhenius diagram of all ceramics at different temperatures, and the values of activation energy E _relax (E _g and E _gb are relaxation activation energy of grain and grain boundary, respectively) of the relaxation process are summarized in this figure and Table 2. The relaxation activation energy values of grains of (1-x) BBT - x BiT (x = 0–0.8) ceramics are no longer given here, because of the conduction mechanisms of these samples are mainly from the grain boundary contribution. In the thermal excited state, the oxygen vacancy will undergo single and double ionization and generate free electrons, as shown in Eq. 4 and Eq. 5 (Wu and Wang, 2010): V O ⇔ V O • + e ′ (4) V O • ⇔ V O • • + e ′ (5) Where V O is unionized oxygen vacancies, e ′ is free electron, and V O   •   a n d   V O • • present singly and doubly ionization oxygen vacancy, respectively. In addition, the activation energy of the double ionized oxygen vacancy is approximately 0.6–1.2 eV, while that of the single ionized oxygen vacancy is less than 0.3–0.5 eV (Ma et al., 2021).

The BBT sample above the Curie temperature T _c (408°C) with an activation energy E _gb of 1.50 eV between 425 and 575°C, and the BBT sample below T _c with an activation energy E _gb of 1.39 eV between 375 and 400°C. This figure exhibits different conduction modes before and after ferroelectric phase transition for the (1-x) BBT - x BiT (x = 0.2–0.8) intergrowth ceramics. The activation energy values of these ceramics are lower at low temperatures and low frequencies before the ferroelectric phase transition, while the values are higher of these ceramics at high temperatures and high frequencies after the ferroelectric phase transition. These values of E _relax are much larger than that of oxygen vacancy activation energy (more than 1.0 eV), indicating that the movement of oxygen vacancy is no longer the main conduction mode at high temperature and the reduction of oxygen vacancy content (Steinsvik et al., 1997).

The BiT ceramics undergo phase transformation near temperature T ₁ (Figure 5F) to exhibit various values of E _relax. At the temperature of 400–450°C, it amounts to 0.43 eV at lower frequency and 0.89 eV at higher frequency, illustrating that the conductivity mainly comes from the conductivity of single ionized oxygen vacancy at lower frequency, and the carrier migration at higher frequency is mainly manifested by oxygen vacancy overcoming potential barrier transition from high energy level to low energy level (Chan et al., 1982; Raymond et al., 2005). At higher temperature of 475–550°C, BiT ceramics show great values of relaxation activation energy, and the dominant modes of conduction are cationic vacancies rather than oxygen vacancy at higher temperature (Guiffard et al., 2005).

Compared with BiT ceramics, the intergrowth structure of BBT-BiT ceramics improves the activation energy of BLSFs and reduces the oxygen vacancy concentration of BLSFs. The more oxygen vacancies can increase the leakage current of the ceramics, resulting in the incomplete polarization of the ceramic and the decrease of the piezoelectric constant, which in keeping with the electric properties results (Figure 6).

Table 2 exhibits the fitting values of resistance (R _g and R _gb), capacitance (C _g and C _gb), relaxation time determined (τ _g and τ _gb) and relaxation activation energy (E _g and E _gb) for bulk and grain boundary at various temperatures of 450°C, 500°C and 550°C, respectively. The resistance and relaxation activation energy values of the grain boundary are greater than those of the grain (R _g > R _gb, E _g > E _gb), and the resistance and relaxation time of ceramics gradually decrease with the increasing of BiT content x and test temperature.

The real part of impedance responses for BBT (a), 0.8BBT-0.2BiT (c), and BiT (e) BiT, as well as imaginary part of impedance responses for BBT (b), 0.8BBT-0.2BiT (d), and BiT (f) BiT are presented in Figure 10. The amplitude of Zʹ decreases with increasing temperature, indicating an increase of the AC conductivity of ceramics with increasing temperature, as shown in Figures 10A, C, E (Rafiq et al., 2015). With the temperature increasing, the maximum values of Zʺ are shifted and broadened at a higher frequency, which indicates the presence of relaxation processes with temperature-dependent, as shown in Figures 10B, D, F (Macdonald and Jhonson, 2005). There are two relaxation peaks at the frequency range of BiT ceramics as shown in Figure 10F, the lower and higher frequencies regions correspond to the impedance response of the grain-boundary and grain respectively.

To study further on the conduction mechanism of ceramics, the AC conductivity (σ _ac) (Majhi et al., 2009) and conductance activation energy (E _con) (Gupta et al., 2018) of samples conform to the following Eq. 6 and Eq. 7, respectively. σ ac = Z ′ Z ′ 2 + Z ″ 2 d A (6) Where d and A are the diameter and area of the sample, respectively. And the AC conductivity σ _ac diagrams of all samples are shown in Figure 11. σ ac = σ 0 exp − E con k B T (7) Where σ ₀ is the pre-exponential term. The values of conductance activation E _con for all samples are summarized in Figure 12.

Figure 11 shows the variation of AC conductivity and frequency of all ceramics at different temperatures. The AC conductivity of the (1-x) BBT - x BiT ceramics increases with x increasing, because the conductivity of BiT ceramics itself is higher than that of BBT ceramics. The values of conductivity for 0.2BBT-0.8BiT ceramics with the optimal performance are in the range of 5.6 × 10^–5–6.8 × 10^–3 S/m below the curie temperature of 475°C (at frequencies of 10²–10⁶ Hz), which is lower than that of single BiT ceramics (0.85–22.74 S/m). Combined with the above analysis results, it can be concluded that the decrease of oxygen vacancy concentration has the main contribution to the decrease of ceramic conductivity. Meanwhile, the AC conductivity increases with the increase of temperature, indicating that conductance is related to thermally activated carriers (mainly oxygen vacancies and local carriers). Moreover, the conductivity of the ceramic is basically unchanged at low frequency, and the conductivity increases with increasing temperature and frequency at higher frequencies by the formation of AC conductance by the charge hopping process (Raymond et al., 2005).

Figure 12 shows the Arrhenius plots of AC conductivity of (1-x) BBT—x BiT ceramics (x = 0–1) at different frequencies, and the activation energy values of ceramics were calculated from the Ln σ—T linear fitting data. As for the (1-x) BBT—x BiT ceramics (x = 0–0.8), the values of activation energy are close to the double ionization activation energy of the oxygen vacancy at the lower frequencies of 1 kHz and 10 kHz, while the values of activation energy are close to the single ionization activation energy of the oxygen vacancy at the higher frequencies of 100 kHz and 1000 kHz. These results indicate that the conductivity mechanism is mainly related to the secondary ionization and single ionization of the oxygen vacancy at lower and higher frequencies, respectively. It also can be seen that the activation energy of BiT ceramics is very low, between 0.05 and 0.38 eV, indicating that its conductive mechanism is no longer oxygen vacancies, the conductive mechanism is probably impurity conductance (Zhang, S. J et al., 2005). Thus, the oxygen vacancies are the main carriers and the main cause of the conductivity behavior for (1-x) BBT—x BiT ceramics (x = 0–0.8), and the conductive mechanism of BiT ceramics is probably impurity conductance.