Optimized thermoelectric performance driven by A-site deficient in SrxLa0.1TiO3/Ti composites

doi:10.21203/rs.3.rs-2959967/v1

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Optimized thermoelectric performance driven by A-site deficient in SrxLa0.1TiO3/Ti composites

https://doi.org/10.21203/rs.3.rs-2959967/v1

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For thermoelectric applications at high temperatures, perovskite SrTiO₃ is a promising contender as an n-type oxide thermoelectric (TE) material. In this work, Sr_xLa_0.1TiO₃/Ti composite ceramics were successfully synthesized by solid state reaction with a combination of the A-site cation vacancy and fabricating composite. Through XRD and SEM analysis, a complex microstructure composed of two kinds of titanium oxides was formed after reductive sintering and annealing. In addition, controlling the content of Sr vacancies in the Sr_xLa_0.1TiO₃ matrix increased the thermoelectric power factor by optimizing the weighted mobility in the composite, resulting in the highest power factor PF of 425.8 µWm^− 1K^− 2 and the maximum ZT of 0.16 for the Sr_0.875La_0.1TiO₃/Ti composite through the synergistic optimization of electrical and thermal transport properties. The strategy proposed in this study of design A-site-deficient composites with nano-sized metal paved a new way to fabricate high performance oxide thermoelectric materials.

SrTiO3 thermoelectrics

A-site deficiency

Composites

Thermoelectric performance

Thermoelectric generators (TEG) are thought of as energy conversion devices that use thermoelectric materials as their primary building blocks to directly transform waste heat into electric energy [1–3]. Thus, it is considered as an attractive solution to alleviate environmental pollution and energy crisis. The dimensionless figure of merit, ZT = S²σT/κ, can be used to calculate the energy conversion efficiency of thermoelectric materials, where S, σ, κ and T stand for the Seebeck coefficient, electrical conductivity, thermal conductivity, and absolute temperature, respectively. The key to achieving a high ZT is to obtain a large S and σ as well as a small κ. Among traditional thermoelectric materials, oxide thermoelectric materials have the advantages of excellent high-temperature thermal stability, environmentally friendly, being nontoxic and economically feasible compared with alloy materials [4]. SrTiO₃ has a cubic perovskite crystal structure, which is a typical representative of potential oxide thermoelectric materials. In order to improve the thermoelectric performance of SrTiO₃ as much as possible, a large number of researchers have made efforts through various strategies, such as single/double doping [5–8], forming composites [9–14], the introduction of in situ vacancy [15–19], entropy entropy engineering [20–24], etc.

It is known that the improvement of the ZT value of SrTiO₃ could be realized by doping rare earth element La or transition metal element Nb. In addition, compositing SrTiO₃ with the high conductivity second phase of micro-nano precipitates, such as metal particles and graphene, etc., is also one of the effective adjusting strategies [25–27]. The conductive second phase of composite serves as an electrical connection that can improve the electrical performance. In the meantime, the existence of the second phase can reduce the thermal conductivity by enhancing interfacial phonon scattering at the grain boundary [28]. In our earlier research, the additional role of metal particles is considered to promote the volatilization of lattice oxygen to produce more oxygen vacancies [29]. SrTiO₃ can accommodate a variety of substituted cations and defects due to its excellent structural stability. Interesting, the A-site stoichiometry of perovskite matrix has an important impact on the electrical transmission performance, as reported in many literatures [30–31]. Similar to the role of oxygen vacancy, the presence of the A-site cation vacancy causes an improvement in mobility of carriers resulting in enhanced electron conductivity.

In this work, the composition of Sr_xLa_0.1TiO₃/Ti was designed by a combination of introducing A-site cation vacancy and forming composites strategy. Sr_xLa_0.1TiO₃/Ti (x = 0.9, 0.875, 0.85, 0.825) composite ceramics were prepared using the solid-state reaction technique. The microstructure, electrical properties and thermal properties of ceramics were investigated. The inherent effects of A-site deficiency on thermal and electrical conductivities were analyzed. As a result, the highest power factor of 425.8 µWm^− 1K^− 2 was obtained in Sr_0.875La_0.1TiO₃/Ti sample, which can be attributed to the highest conductivity as well as relatively high Seebeck coefficient.

The general formula of the composition in this work is Sr_xLa_0.1TiO₃/Ti (x=0.9, 0.875, 0.85, 0.825), which is labeled as S90, S875, S85, S825, respectively. Ceramic samples were synthesized though a solid-state reaction. Stoichiometric amount of SrCO₃ (＞99%), La₂O₃(＞99.9%), TiO₂(＞99.8%), and sintering aid Bi₂O₃(＞99.9%) powders were wet-milled in ethanol for 24 hours using zirconia balls. The fabrication process can be found in our previous work [23,32]. The mixed slurry were dried at 333 K overnight and then calcined in air using an alumina crucible at 1473 K for 4 hours. The calcined powders were mixed with 20 wt% Ti particles (>99.9%, 40 nm), and the mixture was then re-milled again and dried with the same techniques to produce fine powder. The final mixed powders was uniaxially pressed into different size pellets (Φ15 mm×3 mm, 22 mm×4 mm×6 mm) at 100 MPa and burned out the binder at 923 K. The pellets were sintered in a tube furnace at 1723 K for 2 hours under Argon (Ar) flow. Subsequently, the sintered pellets were annealed at 1623 K for 8 hours with a reducing atmosphere of Argon + graphite (Ar+C) .

The phase structure was analyzed using a polycrystalline X-ray diffractometer (XRD, PANalytical X’pert PRO, Holland) at room temperature. Diffraction patterns were gathered using a 2°/min scanning rate for Rietveld refinement. The sample surface was polished for microstructure analysis using the Scanning electron microscopy equipped with an Energy-dispersive X-ray spectroscopy detector (SEM, Quanta 600 FEG, USA). X-ray photoelectron spectroscopy patterns were carried out to investigate the elemental valence states using a Kratos Axis Ultra DLD spectrometer. C 1s peaks at 284.8 eV were used to calibrate the XPS binding energies, Peak fitting of all the characteristic peaks was carried out using the CASA XPS software. The electrical transport properties (Seebeck coefficient and electrical conductivity) were measured using a commercial thermoelectric measurement equipment (LSR-3/1100; LINSEIS, Germany). The thermal conductivity was determined utilizing the formula κ = ρC_pD, volume density (ρ) was measured from Archimedes method. Thermal diffusivity (D) and specific heat capacity (C_p) were obtained by a laser flash apparatus (LFA 427, NETZSCH, Germany). The lattice thermal conductivity is estimated with the use of expression κ_L = κ_tot-κ_e, where κ_e is obtained from Wiedemann-Franz law (κ_e = LσT). L is the Lorentz number, which is determined by the single parabolic band function of acoustic phonon scattering ($L=1.5+\text{e}\text{x}\text{p}[-\left|S\right|/116]$) [33]. Hall effect measurements were carried out using a Hall measurement device (ET9105-HS, China) based on the Van der Pauw method under a reversible magnetic field of 1.0 T to determine carrier concentration and mobility.

In order to verify the phase structure, the synthesized ceramics were examined by XRD at room temperature. The measured XRD spectra of Sr_xLa_0.1TiO₃/Ti composite ceramics sintered in argon and then annealed at 1350°C are shown in Fig. 1. Rietveld refinement (Fig. 2.) was carried out on all XRD spectra to confirm that the main diffraction peaks can be indexed to Sr_0.9La_0.1TiO₃ cubic perovskite phase (PDF#79–0179, space group $\text{P}\text{m}\overline{3}\text{m}$). With the increase of A-site cation vacancy content, the diffraction angle of the XRD diffraction peak does not shift significantly, and the lattice parameters of cubic phase calculated by Rietveld analysis also show unconspicuous changes overall, as shown in Table 1. This reflects the subtle and complex effect of defects formed by charge compensation mechanism on lattice expansion/contraction. Generally, the reduction from Ti⁴⁺ ions to Ti³⁺ ions and oxygen vacancies, can promote the expansion of perovskite lattice, while A-site cation vacancies can lead to lattice contraction. The lattice parameters of samples containing A-site cation vacancies significantly increase compared to pure STO (3.905 Å). Considering that the ionic radius of La (1.36 Å) is smaller than that of Sr (1.44 Å), the lattice expansion is mainly due to the generation of oxygen vacancies. However, for the defect formation process under reductive atmosphere conditions, the A-site cation vacancy promotes the formation of oxygen vacancies to a certain extent, while combining with each other to form Schottky defect pairs, which can be been as self compensation for fluctuations in crystal volume. All XRD spectra show several small additional diffraction peaks, which are at ~ 30°, ~ 42° and ~ 51°, respectively. The detected secondary phase peaks can be marked as ${\text{T}\text{i}}_{1-\text{m}}^{4+}{\text{T}\text{i}}_{\text{m}}^{3+}{\text{O}}_{2-\text{m}/2}$ (Ti₃O₅ and Ti₆O₁₁) phases, which are mainly the oxidation products formed by nano-sized metal titanium after sintering. The presence of secondary phases is further confirmed by the contrast between light and dark ceramic microstructure, as shown in Fig. 3. A reasonable explanation for the evolution of phase composition is that during the Ar atmosphere sintering and Ar + C annealing stages, the titanium dioxide formed by the oxidation of nano-sized Ti powder further reacts with Sr_xLa_0.1TiO₃ to form a new perovskite phase, and the remaining titanium dioxide is reduced to the secondary phase ${\text{T}\text{i}}_{1-\text{m}}^{4+}{\text{T}\text{i}}_{\text{m}}^{3+}{\text{O}}_{2-\text{m}/2}$ under the reducing atmosphere. Additionally, it should be noticed that the observed diffraction peak intensity of the secondary phase has not changed significantly with A-site composition, indicating that the content of the secondary phase does not increase significantly.

Table 1

Crystal structural information and measured densities for Sr_xLa_0.1TiO₃/Ti composites.
Composition code	Nominal composition	Space group	Cell parameter (Å)	Cell Volume (Å³)	Measured density (g/cm³)
Pristine STO	SrTiO₃	$\text{P}\text{m}\overline{3}\text{m}$	3.905	59.5474	/
S90	Sr_0.9La_0.1TiO₃	$\text{P}\text{m}\overline{3}\text{m}$	3.90928(5)	59.7437	4.77
S875	Sr_0.875La_0.1TiO₃	$\text{P}\text{m}\overline{3}\text{m}$	3.90956(9)	59.7567	4.78
S85	Sr_0.85La_0.1TiO₃	$\text{P}\text{m}\overline{3}\text{m}$	3.90962(4)	59.7592	4.81
S825	Sr_0.825La_0.1TiO_2.975	$\text{P}\text{m}\overline{3}\text{m}$	3.90924(6)	59.7419	4.71

To evaluate the microscopic morphology of the annealed sample, the SEM back-scattered electron (BSE) images of Sr_xLa_0.1TiO₃/Ti ceramics are present in Fig. 3. No visible pores were observed in the samples, which revealed that uniform and well dense ceramics were obtained. Two kinds of grain phases can be found in all ceramics, gray grain clusters as intergranular phase and lighter colored grains as matrix phase. In terms of size and shape, these grains have no significant difference between different component microstructures. The lighter colored grains are irregular polygon shape with an average size range of 4.5 ~ 6 µm. The concentration of A-site cationic vacancies has no significant regular effect on the size of the grains. The illustration of histograms in Fig. 3 shows the grain size distribution of four component samples. Compared with S85 and S825 samples, the grain size distribution of S90 and S875 is more concentrated. Furthermore, it is observed that the sizes of different gray grains in the matrix of Sr_xLa_0.1TiO₃/Ti show noticeable difference. The apparent density of ceramic samples first increases and then decreases with the increase in content of A-site vacancies, indicating a similar trend in relative density, as shown in Table 1. This is because the large-sized ions with 12 coordination occupy the A-site in the ABO₃ perovskite structure, where the generation of cationic vacancies facilitates ion transfer and enhances ion diffusion during the ceramic densification process [34]. However, excessive cation vacancies can also lead to an increase in pores in the microstructure, which results in a decrease in density. The distribution of all elements by EDS mapping for secondary phase of the S875 sample showed an evident aggregation of Ti and O and the absence of Sr, as shown in Fig. 4. To identify light and dark gray grains, element composition analysis measured from the corresponding area is performed and listed in Table 2. Therefore, it can be further inferred that the dark gray phase is a non-stoichiometric titanium oxide phase. The concentration of A-site vacancies in S875 sample normalized to Ti level is 16.2%, which can be inferred from the measured concentration of (Sr + La + Bi). EDS quantitative elemental analysis shows that the (Sr + La + Bi)/Ti ratio is lower than the nominal ratio of S875 sample. It should be attributed to the further reaction between TiO₂ generated by nano-sized metal titanium and Sr_0.875La_0.1TiO₃, which further increases the concentration of actual A-site vacancies [28].

Table.2. EDS point analysis in the secondary phase and matrix grains of S875 composite sample.

Composition	Element	Content(atom%)
Point 1#	Sr	17.55
	Ti	25.55
	O	53.04
	La	2.48
	Bi	1.38
Point 2#	Ti	39.67
Point 2#	O	60.33

The degree of reduction of Ti in four components was assessed by X-Ray photoelectron spectroscopy to further reveal their electronic structure characteristics. The high-resolution deconvoluted Ti 2p spectra are presented in Fig. 5. The asymmetric Ti 2p peak indicates the simultaneous presence of Ti⁴⁺and Ti³⁺. The ratios of Ti³⁺ compared to Ti⁴⁺ were estimated by the deconvolution characteristic peaks and the calculated relevant peak areas [35]. The [Ti³⁺] content for Sr_xLa_0.1TiO₃/Ti samples decreases with increasing Sr vacancy content (from 21.75–4.68%), as shown in Table 3. It is widely known that the concentration of [Ti³⁺] is closely correlated with the carrier concentration in SrTiO₃-based ceramics [17]. It suggests that A-site vacancy has an effect on carrier concentration, which prohibits the reduction process from Ti⁴⁺ to Ti³⁺. As a matter of fact, the detected Ti³⁺ in the sample may come from both the titanium oxide intergranular phase and the matrix phase according to the above experimental results. Meanwhile, the generation of Ti³⁺ in the matrix phase can be divided into two stages, which can be illustrated by taking the S875 sample as an example. The first stage of reduction is electronic defect charge compensation caused by La³⁺ doping, which is described as:

$$\text{0.875SrC}{\text{O}}_{3}+0.05\text{L}{\text{a}}_{2}{\text{O}}_{3}+\text{T}\text{i}{\text{O}}_{2}\underrightarrow{\text{c}\text{a}\text{l}\text{c}\text{i}\text{n}\text{a}\text{t}\text{i}\text{o}\text{n}}0.875{\text{S}\text{r}}_{\text{S}\text{r}}^{\text{x}}\text{+0.1}{\text{L}\text{a}}_{\text{S}\text{r}}^{{\bullet }}\text{+}$$

$$\text{0.9}\text{5}{\text{T}\text{i}}_{\text{T}\text{i}}^{\text{x}}+{0.025\text{V}}_{\text{S}\text{r}}^{{\prime }{\prime }}+0.05{\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }} \left[{\text{T}\text{i}}^{3+}\left(\text{d}\right)\right]\text{+3}{\text{O}}_{\text{O}}^{\text{x}}\text{+}\text{0.0125}{\text{O}}_{2}\text{↑}\text{+}\text{0.875C}{\text{O}}_{2}\uparrow$$

Accompany by the release of oxygen, the Ti ion in the A-site nonstoichiometric samples will be reduced to electronic defects ${\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }}$ that can be recorded as Ti³⁺ (d), and the strontium deficiency partially offsets the excess charge of La³⁺. The second stage of reduction depends on the generation and ionization of oxygen vacancy defects by the reduction atmosphere treatment. The electronic defect Ti³⁺ generated during this process can be recorded as Ti³⁺ (v). The chemical reaction can be presented as:

$$\text{Ti}{\text{O}}_{2}\underrightarrow{\text{r}\text{e}\text{d}\text{u}\text{c}\text{t}\text{i}\text{o}\text{n}}(1-4{\epsilon }){\text{T}\text{i}}_{\text{T}\text{i}}^{\text{x}}+4{\epsilon }{\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }} \left[{\text{T}\text{i}}^{3+}\left(\text{v}\right)\right]+(2-2{\epsilon }){\text{O}}_{\text{O}}^{\text{x}}+{2{\epsilon }\text{V}}_{\text{O}}^{{\bullet }{\bullet }}+{\epsilon }{\text{O}}_{2}\uparrow$$

$${\text{S}\text{r}}_{0.875}{\text{L}\text{a}}_{0.1}{\text{(Ti}}_{0.95}^{4+}{\text{Ti}}_{0.05}^{3+}{\text{)O}}_{3}\underrightarrow{\text{r}\text{e}\text{d}\text{u}\text{c}\text{t}\text{i}\text{o}\text{n}}0.875{\text{S}\text{r}}_{\text{S}\text{r}}^{\text{x}}\text{+}{0.025\text{V}}_{\text{S}\text{r}}^{{\prime }{\prime }}\text{+}\text{0.1}{\text{L}\text{a}}_{\text{S}\text{r}}^{{\bullet }}\text{+(0.95-4θ)}{\text{T}\text{i}}_{\text{T}\text{i}}^{\text{x}}+$$

$$0.05{\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }}+4{\theta }{\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }} \left[{\text{T}\text{i}}^{3+}\left(\text{v}\right)\right]\text{+}{2{\theta }\text{V}}_{\text{O}}^{{\bullet }{\bullet }}\text{+(3-2θ)}{\text{O}}_{\text{O}}^{\text{x}}+{\theta }{\text{O}}_{2}\uparrow$$

Where ε and θ are respectively defined as oxygen vacancies concentrations of the matrix phase and the second phase titanium dioxide generated during the reductive process. Formula 2 explains how secondary phases ${\text{T}\text{i}}_{1-\text{m}}^{4+}{\text{T}\text{i}}_{\text{m}}^{3+}{\text{O}}_{2-\text{m}/2}$ are produced. Taking into account the contribution of A-site vacancies to charge neutrality, the charge balance equation can be written as:

$${[\text{T}\text{i}}_{\text{T}\text{i}}^{{\prime }}\left]{+2[\text{V}}_{\text{S}\text{r}}^{{\prime }{\prime }}\right]\approx \left[{\text{L}\text{a}}_{\text{S}\text{r}}^{{\bullet }}\right]+2\left[{\text{V}}_{\text{O}}^{{\bullet }{\bullet }}\right]$$

The general consensus is that, [Ti³⁺] is equivalent to the concentration of n-type charge carriers (n) under reduction conditions. Therefore, the introduction of A-site cation vacancies decreases the carrier concentration of Sr_xLa_0.1TiO₃/Ti composite samples, which is consistent with the subsequent Hall test results.

Table.3. XPS results for Ti cations in Sr_xLa_0.1TiO₃/Ti composite samples.

Composition	Ti 2P (%)
Composition	Ti⁴⁺	Ti³⁺
S90	78.25	21.75
S875	84.26	15.74
S85	88.55	11.45
S825	95.32	4.68

Various electrical transport characteristics were measured for four Sr_xLa_0.1TiO₃/Ti ceramics, as shown in Fig. 6. The electrical conductivity of all samples with different A-site deficient levels are plotted as a function of temperature in the range of 300 K to 1073 K. It can be seen that the electrical conductivity shows a similar tendency that first increases substantially until 523 K ~ 573 K, then gradually decreases with a small slope with the rising temperature. The samples show semiconductor behavior at a relatively low temperature range, which indicates that the raising carrier concentration and mobility increasing the electrical conductivity, and the carriers are mainly scattered by ionized impurities (donor ion). At higher temperature ranges, the lattice vibration and acoustic phonon scattering are dominant, and the carrier mobility decreases continuously with rising temperature, leading to a decrease in conductivity [36]. Therefore, the semiconductor-metal (S-M) transition of conductivity is approximately 550 K. Besides, it is clear that Sr vacancies can improve the electrical conductivity, especially at low temperatures. At 1073 K, the electrical conductivity of S875 composite is still slightly higher than that of S90 composite.

In order to investigate the electrical transport mechanism, the electrical conductivity below the S-M transition temperature is analyzed by the thermally excited small polaron hopping (SPH) conduction mechanism, which is expressed by the following equation [37–38]:

$${\sigma }=\frac{{\sigma }_{0}}{T}\text{e}\text{x}\text{p}\left[-\frac{{E}_{\text{H}\text{o}\text{p}}}{{k}_{\text{B}}T}\right]$$

where σ₀ is a constant, T is the absolute temperature, E_Hop is the activation energy of the small polaron transition, and K_B is the Boltzmann constant. The relationship between lnσT and K_BT is depicted in Fig. 6b, indicating that the carriers of all samples followed the SPH conduction model. With increasing Sr vacancy concentration in Sr_xLa_0.1TiO₃/Ti samples, the activation energy (E_Hop) decreases, providing more opportunity for the transition of small phonon, as displayed in Table 4. Therefore, A-site vacancy for Sr_xLa_0.1TiO₃/Ti ceramics is beneficial for the electrical transport performance.

The negative Seebeck coefficients of all samples have been obtained in the whole measuring temperature range, as shown in Fig. 6c. This suggests that Sr_xLa_0.1TiO₃/Ti ceramics are n-type semiconductor materials with electrons as dominant charge carriers, which also corresponds to the negative value of the carrier concentration (Table 4). The absolute Seebeck value of all samples increases monotonously with increasing temperature from 373 K to 1073 K. The S85 sample reaches the highest absolute Seebeck coefficient among all samples, which is ~ 208 µV/K at 1073 K. In addition, as Sr vacancy concentration increasing, the absolute value |S| of all samples showed a trend of first increasing and then decreasing in the whole measuring temperature range. It is worth noting that the |S| value in 1073 K has been found to be increased by 6.7% from 195 µV/K for S90 sample to 208 µV/K for S875 sample. Obtaining a high Seebeck coefficient and conductivity simultaneously primarily depends on large carrier mobility, as shown in Table 4.

The temperature dependence of the calculated power factor (Sσ²) is shown in Fig. 6d. For all samples, the power factor (PF) value increases with increasing temperature. The PF value follows a similar trend to Seebeck coefficient, initially increasing with Sr vacancy concentration and then decreasing. It was observed that S875 sample exhibits the highest PF value through the combination of the highest electrical conductivity and Seebeck coefficient. At 873 K, the maximum PF value of ~ 426 µW/mK² is obtained for S875 sample, which is almost 20% increase compared with S90 sample (~ 356 µW/mK²).

Furthermore, the carrier concentration (n) and mobility (µ_H) of Sr_xLa_0.1TiO₃/Ti composites at room temperature were measured using Hall measurement technology, and it was found that the carrier concentration range of all samples is (1.3 ~ 2.3)×10¹⁹ cm³, as shown in Table 4. The Seebeck coefficient expression of degenerate semiconductors can be used to estimate the effective mass (m*) of all samples and is displayed as follows [39].

$$S=\frac{8{{\pi }}^{2}{k}_{\text{B}}^{2}}{3e{ℎ}^{2}}{m}^{\ast }T{\left(\frac{{\pi }}{3n}\right)}^{2/3}$$

where e is the electric quantity of the electron and h is the Planck constant. It can be seen from the Pisarenko diagram (Fig. 6f) that the effective mass of electrons at room temperature shows a decreasing trend with increasing Sr vacancy concentration. In order to better illustrate the significant role of electron mobility in improving the electrical transmission performance, the weighted mobility (µ_w) independent of the carrier concentration is calculated using Eq. 6, which is calculated using the data of electrical conductivity and Seebeck coefficient [40–41].

$${\mu }_{w}=\frac{3{ℎ}^{3}\sigma }{8\pi e{\left(2{m}_{e}{k}_{\text{B}}T\right)}^{3/2}}\left[\frac{\text{e}\text{x}\text{p}\left[\frac{\left|S\right|}{{k}_{\text{B}}/e}-2\right]}{1+\text{e}\text{x}\text{p}\left[-5\left(\frac{\left|S\right|}{{k}_{\text{B}}/e}-1\right)\right]}+\frac{\frac{3}{{\pi }^{2}}\frac{\left|S\right|}{{k}_{\text{B}}/e}}{1+\text{e}\text{x}\text{p}\left[5\left(\frac{\left|S\right|}{{k}_{\text{B}}/e}-1\right)\right]}\right]$$

where m_e is the electronic mass. Sr_xLa_0.1TiO₃/Ti composites have exhibited an upward trend of the weighted mobility prior to the S-M transition temperature (373–573 K), attributed to the thermal activation of carriers under the action of high temperature (Fig. 6e). Subsequently, in the temperature region of metallic conductivity behavior, driven by the carrier acoustic phonon scattering, µ_w decreases in a T^3/2 dependence with temperature. At about 480 K, the maximum weighted mobility of S85 sample is ~ 15.2 cm²V^− 1S^− 1, which is 26% higher than that of S90 sample due to the introduction of A-site deficiency.

Table.4. Activation energy (E_Hop) of the small polaron transition, room-temperature carrier concentration (n) and carrier mobility (µ_H) for Sr_xLa_0.1TiO₃/Ti composite samples.

Composition	E_Hop (eV)	n (×10¹⁹ cm³)	µ_H (cm²/V/s)
S90	-0.0739 ± 0.0038	2.24	23.28
S875	-0.0646 ± 0.0033	2.15	49.15
S85	-0.0614 ± 0.0035	1.31	109.21
S825	-0.0609 ± 0.0032	1.40	77.10

The temperature dependence of the total thermal conductivity of Sr_xLa_0.1TiO₃/Ti composite ceramics are presented in Fig. 7a. The total thermal conductivity of all samples within the measured temperature range drop monotonically as temperature rises, originating from the phonon-phonon Umklapp scattering. It is noteworthy that the total thermal conductivity increases slightly and then decreases with the increase of Sr vacancy in the whole temperature range. The electronic thermal conductivity exhibits a similar variation trend with temperature in accord with the electrical conductivity. It is notable that the proportion of κ_e/κ_tot is about 7%, indicating that the lattice thermal conductivity plays a dominant role in the total thermal conductivity. Furthermore, The κ_L reduced following the similar trend of the κ_tot with the increasing Sr vacancy content and increasing temperature in the whole temperature range. This phenomenon mainly depends on the additional phonon scattering due to the increase of crystal lattice disorder by A-site deficiency. The effect of A-site vacancy on decreasing lattice thermal conductivity is particularly remarkable at 373 K. The lattice thermal conductivity for S90 sample is 3.56 Wm^− 1K^− 1, while the lattice thermal conductivity for S825 sample decreases to 3.25 Wm^− 1K^− 1, decreasing by about 9%. As a consequence, it was discovered that with rising temperature, the dimensionless figure-of-merit ZT value for Sr_xLa_0.1TiO₃/Ti composite ceramics increased, as shown in Fig. 7d. The variation trend of ZT value with component is the same as PF value. The S875 sample obtained the highest ZT value, 0.16, which was attained at 1023 K. As can be seen, the thermoelectric properties of Sr_xLa_0.1TiO₃/Ti composites can be further enhanced by introducing an appropriate amount of A-site vacancies into the perovskite structure.

The composite ceramics of Sr_xLa_0.1TiO₃/Ti with a series of Sr-deficiency have been prepared by the solid-state reaction route. By the reductive sintering and annealing, the formation of fractional ${\text{T}\text{i}}_{1-\text{m}}^{4+}{\text{T}\text{i}}_{\text{m}}^{3+}{\text{O}}_{2-\text{m}/2}$ (Ti₃O₅ and Ti₆O₁₁) phases from the oxidation of nano Ti was observed in the microstructure. Moreover, the presence of Sr vacancies enhances the weighted mobility of charge carrier in the Sr_xLa_0.1TiO₃/Ti system. As a result, the maximum PF value of 425.8 µWm⁻¹K⁻² has been attained in S875 composite. At the same time, the presence of Sr vacancies can also increase the phonon scattering at relatively high Sr vacancy concentration. Consequently, a maximum ZT value of 0.16 (at 1073 K) was achieved for the S875 composites. The presented approach is expected to cooperatively optimize thermoelectric characteristics of composite oxide thermoelectric ceramics based on SrTiO₃.

Author contribution Zhihao Lou: Conceptualization, Methodology, Investigation, Data Curation, Formal analysis, Writing-Original Draft. Ping Zhang: Conceptualization, Investigation, Data Curation, Writing-Review & Editing. Penghui Chen: Investigation, Data Curation, Writing-Review & Editing. Ziyao Wei: Investigation, Data Curation, Writing-Review & Editing. Jianjun Gou: Data Curation, Writing-Review & Editing. Jie Xu: Data Curation, Writing-Review & Editing. Chunlin Gong: Data Curation, Writing-Review & Editing. Feng Gao: Conceptualization, Methodology, Supervision, Writing-Review & Editing, Funding acquisition.

Funding This work was supported by National Natural Science Foundation of China (52272123, 52072301), the National Key R&D Program of China (Grant Number: 2022YFB3504901), the Undergraduate Innovation and Entrepreneurship Training Program of Shaanxi Province (S202210699186), Guangxi Science and Technology Plan Project (Gui Ke AB22035043), the Open Fund of State Key Laboratory of New Ceramic and Fine Processing Tsinghua University (No. KFZD202102), and the ‘111’ Project (No. B20028). We would like to thank the Analytical & Testing Center of Northwestern Polytechnical University for the measurements of SEM, EDS and valuable discussion.

Conflict of interest The authors declare no competing interests.

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Optimized thermoelectric performance driven by A-site deficient in SrxLa0.1TiO3/Ti composites

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Composition code	Nominal composition	Space group	Cell parameter (Å)	Cell Volume (Å³)	Measured density (g/cm³)
Pristine STO	SrTiO₃	\(\text{P}\text{m}\overline{3}\text{m}\)	3.905	59.5474	/
S90	Sr_0.9La_0.1TiO₃	\(\text{P}\text{m}\overline{3}\text{m}\)	3.90928(5)	59.7437	4.77
S875	Sr_0.875La_0.1TiO₃	\(\text{P}\text{m}\overline{3}\text{m}\)	3.90956(9)	59.7567	4.78
S85	Sr_0.85La_0.1TiO₃	\(\text{P}\text{m}\overline{3}\text{m}\)	3.90962(4)	59.7592	4.81
S825	Sr_0.825La_0.1TiO_2.975	\(\text{P}\text{m}\overline{3}\text{m}\)	3.90924(6)	59.7419	4.71