The effect of reduced graphene oxide on microstructure and thermoelectric properties of Nb-doped A-site-deficient SrTiO3 ceramics

Journal of Alloys and Compounds 786 (2019) 884e893

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The effect of reduced graphene oxide on microstructure and thermoelectric properties of Nb-doped A-site-deficient SrTiO3 ceramics Cao Wu a, Jia Li b, Yuchi Fan a, c, *, Juanjuan Xing b, **, Hui Gu b, Zhenxing Zhou a, Xiaofang Lu a, Qihao Zhang d, Lianjun Wang a, Wan Jiang a, c, e a

State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, College of Materials Science and Engineering, Donghua University, Shanghai, 201620, China Materials Genome Institute, School of Materials Science and Engineering, Shanghai University, Shanghai, 200444, China c Institute of Functional Materials, Donghua University, Shanghai, 201620, China d Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China e School of Material Science and Engineering, Jingdezhen Ceramic Institute, Jindezhen, 333000, China b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 5 December 2018 Received in revised form 28 January 2019 Accepted 30 January 2019 Available online 5 February 2019

As promising thermoelectric materials with great stability at high temperature, SrTiO3 based ceramics have attracted much attention, albeit the thermoelectric figure of merit remains relatively low for application. To improve the thermoelectric properties, reduced graphene oxide (RGO) is successfully incorporated into Nb-doped SrTiO3 (STNO) ceramics with Sr deficiency via a hetero-aggregation strategy followed by annealing and spark plasma sintering. The addition of RGO enhances the carrier concentration via reduction effect and increases the carrier mobility probably through lowering the double Schottky barrier at grain boundaries, resulting in highly improved electrical conductivity for the composites. Moreover, RGO also stimulates the separation of Nb rich TiO2 phase, which serves as a dispersion with low thermal conductivity to prohibit the phonon transport in the composites. Consequently, a ZT value reaching 0.22 at 800 K for the RGO/STNO nanocomposite is obtained, which is 1.8 times higher compared with the monolithic STNO. © 2019 Elsevier B.V. All rights reserved.

Keywords: Thermoelectric properties Composites Perovskite Graphene Secondary phase

1. Introduction Thermoelectric (TE) materials that can directly generate electricity from heat have attracted tremendous interests in the past decades for their great potential in recycling waste heat. However, the widespread application of thermoelectric materials has not been realized to date because of the relatively low energy conversion efficiency, which is closely related to the dimensionless figure of merit ZT ¼ sS2 T=k, where s is the electrical conductivity, S is the Seebeck coefficient, T is the absolute temperature, and k is the thermal conductivity. In the past decades, a number of TE materials

* Corresponding author. State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, College of Materials Science and Engineering, Donghua University, Shanghai, 201620, China. ** Corresponding author. E-mail addresses: [email protected] (Y. Fan), [email protected] (J. Xing). https://doi.org/10.1016/j.jallcom.2019.01.376 0925-8388/© 2019 Elsevier B.V. All rights reserved.

with high ZT values, such as chalcogenides, skutterudites and other intermetallic compounds have been developed [1e5], which are generally composed of toxic, natural rare and heavy elements whose usage should be limited. In addition, these traditional TE materials are of relatively low melting point, which is not suitable for retrieving electricity from heat sources of high temperature such as industrial furnaces and incinerators [4,6]. In contrast, oxide based thermoelectric materials show their unique advantages in low toxicity, earth abundance and stability at high temperature, which are very promising as high-temperature oriented TE materials. So far a number of p-type oxide based TE materials have been developed including NaCo2O4 and Ca3Co4O9, whose ZT values can reach ~1 for the single crystalline materials [7,8]. However, the TE performance for n-type oxide materials is still far from satisfaction [9], which could severely decrease the efficiency of TE devices that need both n- and p-type materials with high-performance. The perovskite structured strontium titanate (SrTiO3) is one of the most promising n-type oxide TE materials with great structural

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tolerances for optimizing TE properties via A- or B-site doping. Park et al. [10] have shown that La-doped SrTiO3 prepared from colloidal synthesized nanoparticles exhibit a maximum ZT of 0.37 at 973 K with 9.0 at% La doping, owing to the simultaneously improved electrical conductivity and decreased thermal conductivity through tuning doping ratio and nanostructure. B-site doped SrTiO3 with Nb also has been frequently reported due to the large density of states effective mass [11]. Recently, Wang et al. prepared Nb-doped SrTiO3 with a ZT of 0.4 at 1100 K via hydrothermal method [12], and a La and Nb co-doped nanostructured SrTiO3 with a ZT value of 0.6 at 1150 K in their later work, which is the highest TE performance for perovskite at present [13]. Furthermore, it has been suggested that introducing A-site deficiency in doped SrTiO3 is beneficial to improve the electrical conductivity of doped SrTiO3 by facilitating the generation of oxygen vacancy, while depress the thermal conductivity by increased point defects [14]. Apart from doping, the ZT value can also be effectively enhanced in terms of composite, in which the second phase should be carefully selected. As the first discovered 2D material with many excellent properties [15], graphene displays great potential in optimizing the properties of TE materials. Xie et al. reported that graphene/CoSb3 nanocomposite has a decreased thermal conductivity owing to the enhanced phonon scattering upon addition of graphene [16]. Similar result has also been reported in reduced graphene oxide (RGO)/YbyCo4Sb12 composite [17]. For the perovskites, Freer et al. [18] found that the TE performance of La doped SrTiO3 could be significantly boosted by the addition of graphene. Our previous work indicated that the highly promoted oxygen vacancies by the mild reaction between RGO and SrTiO3 matrix should be the dominant reason behind the highly enhanced electrical conductivity [19]. However, the influence of graphene on the TE properties of A-site deficient SrTiO3 has not been investigated except for very few works [20]. In this work, the hydrothermally synthesized Nb doped A-sitedeficient SrTiO3 powder with nominal composition of Sr0.93Ti0.9Nb0.1O3 (STNO) was successfully incorporated with graphene oxide via a hetero-aggregation strategy, which refers to the process of mixing two types of colloid particles possessing opposite surface charge for obtaining very uniform mixture. After reduced in H2/Ar forming gas and sintered by spark plasma sintering (SPS), RGO/STNO composites with various graphene contents were obtained. Compared with the monolithic STNO ceramic, the obtained RGO/STNO composite shows not only large power factor but also extremely low thermal conductivity, which is deeply related to the composition and microstructure transition induced by RGO.

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trimethoxy-silane (APTES) was carried out by refluxing. Firstly the STNO powder (1.5 g) was added into toluene (150 ml) followed by ultrasonic vibration 30 min to get well-dispersed suspension, then 2 ml APTES was added into the suspension. The mixture was refluxed at 423 K and kept for 6 h under stirring in Ar atmosphere. Finally, the surface modified powder was filtered, washed with ethanol and deionized water followed by drying at 333 K in a vacuum oven. 2.3. Fabrication of RGO/STNO composite Graphene oxide (GO) was synthesized by the modified Hummers method [23]. The STNO (1.2 g) powders were added into 300 ml deionized water and adjusted the pH to ~ 3.3 followed by ultrasonic treatment for 2 h to get a homogeneous colloid. Various amount of GO colloid were added to the colloid of STNO under stirring, and then the obtained hybrid mixture with GO content of 0.3 wt%, 0.6 wt% and 1.0 wt% was filtered and dried at 333 K. The GO/STNO hybrid powder was reduced at 1373 K in of 10% H2/Ar forming gas for 6 h. The obtained RGO/STNO hybrid powder was densified by SPS at 1573 K under a pressure of 80 MPa in vacuum. After sintering, the RGO/STNO composites with increasing RGO contents (denoted as S0, S1, S2 and S3) were obtained. 2.4. Characterization The scanning electron microscopy (SEM) observation and electron back scattered diffraction (EBSD) were performed using JEOL JSM-6500F equipment. The transition electron microscopy (TEM) observation was performed using JEOL JSM-7100F. Elemental analysis was performed by using inductively coupled plasma spectroscopy (Prodigy-ICP). The X-ray diffraction (XRD) patterns were obtained by Rigaku Smartlab 9kw diffractometer with Cu Ka radiation source. Raman spectra (Tokyo Instruments Co.) were collected using 532 nm wavelength incident laser light. The electrical conductivity and Seebeck coefficient were measured by ZEM-3 (Ulvac-Riko) under a low-pressure helium atmosphere at 300e800 K. The carrier concentration and the mobility were measured by a Hall measurement system (Toyo ResiTest8400 series) at room temperature in the air. The thermal diffusivity was measured at 300e800 K by a laser flash method (LFA 457, Netzsch), and the specific heat was measured at 300e800 K using a differential scanning calorimeter (DSC 404F3, Netzsch). The apparent density was measured by Archimedes method for all bulk samples.

2. Experimental

3. Results and discussion

2.1. Fabrication of STNO

3.1. Structural characterization

For synthesizing Sr0.93Ti0.9Nb0.1O3 powder, 9 mmol of Ti(OBu)4 and 1 mmol of NbCl5 were mixed in 30 ml of ethylene glycol, and followed by adding 15 ml NaOH solution to adjust the pH under mechanical stirring. The strontium nitrate (9.3 mmol) dissolved in 15 ml of deionized water was then added to obtain a microemulsion. The microemulsion was transferred into a 100 ml polytetrafluoroethylene (PTFE) vessel. The vessel was sealed and placed inside a stainless steel autoclave, which was heated in an oven for 24 h at 455 K. After cooling down to room temperature, the products were washed by centrifuge with ethanol followed by drying at 333 K.

The elemental composition of the as-prepared STNO ceramic and composites was determined by ICP at first. It is found that the actual ratio of Sr, Ti and Nb in S0 bulk sample is Sr: Ti: Nb ¼ 0.91: 0.915: 0.085, which is close to the nominal composition but with more Sr deficiency. Fig. 1a shows the XRD patterns of as-sintered pure STNO (S0) and RGO/STNO composite with various filler content (S1, S2 and S3). The dominant diffraction peaks of all samples are in good consistence with cubic perovskite STO (Pm-3m space group, JCPDF 35-0734), which is unchanged upon the addition of graphene. Meanwhile, the lattice parameter calculated from the XRD data for the RGO/STNO composite slightly decreases with increasing content of RGO (Table 1), suggesting a subtle composition change in the matrix material. No diffraction peaks of RGO in composite can be identified due to the low graphene content (at most 1.0 wt% for S3). However, small peaks that are indexed to

2.2. Surface modification of STNO The surface modification of STNO powder with (3-Aminopropyl)

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Fig. 1. XRD patterns for STNO samples. (a) Wide-range patterns for S0, S1, S2 and S3. (b) Details of the diffraction peak in (a). (c) Low-angle patterns for GO/STNO powder, RGO/STNO powder and sintered RGO/STNO composite, respectively.

Table 1 Basic information for STNO and RGO/STNO composites. Samples

Space group

Cell parameter (Å)

Unit cell volume (Å3)

Apparent density (gcm1)

Effective mass¶ (m*)

S0 S1 S2 S3

Pm-3m Pm-3m Pm-3m Pm-3m

3.91981 3.91998 3.91972 3.91825

60.2275 60.2353 60.2234 60.1557

5.054 4.993 4.990 4.950

1.23m0 1.18m0 1.49m0 1.75m0

m0 is the mass of free electron.

rutile TiO2 can be seen in all composite samples, reflecting the influence of RGO on the matrix phase. Detailed study shows that the rutile TiO2 was generated during the H2/Ar reducing process. After sintering by SPS, the normalized intensity of TiO2 diffraction peaks shows no noticeable change, implying that the amount of second phase was not altered after SPS (Fig. 1c). Since the second phase of TiO2 cannot be identified in STNO ceramic without RGO from XRD result but can be seen in RGO/STNO composite without annealing (Fig. S1), it can be concluded that the addition of GO greatly promoted the separation of TiO2 under reducing atmosphere. In order to clarify the mechanism of phase separation, the microstructure of monolithic STNO ceramic and composite was investigated by SEM (Fig. 2). It is found that very small amount of TiO2 has been formed in monolithic STNO after reducing in 10% H2/ Ar gas for 6 h and SPS (Fig. 2a), though the amount is too low to be

detected by the XRD equipment. More interestingly, elemental analysis shows that the composition of TiO2 in STNO ceramic and RGO/STNO composites is different. It is observed that the separated TiO2 grain in the composite contains much higher concentration of Nb (Fig. 2b) in comparison to the TiO2 precipitate in monolithic STNO, which reflects the different reducing effect with or without RGO. Quantitative analysis by EDS further reveals that the Nb/Ti ratio in the second phase is roughly as same as that in the matrix (Table S1). Furthermore, either increasing the reduction time or raising the relative concentration of H2 can lead to increased amount of TiO2 precipitates in monolithic STNO ceramics (Fig. S2), implying that the separation of TiO2 is due to the reduction and closely related to the reducing condition. Apparently, RGO exhibits stronger reducing ability in the phase separation process compared with H2/Ar forming gas, which results in much higher TiO2 content

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Fig. 2. SEM images of polished surfaces and corresponding EDS mapping of Sr, Ti, Nb and O for (a, b) S0 and (c, d) S2, showing different composition of TiO2 in S0 and S2.

in the STNO composites compared to the monolithic STNO, as shown in Fig. 3. According to XRD result, the TiO2 content increases with increasing RGO fraction in the composites, although the S3 sample shows decreased amount of TiO2 probably owing to the less homogenous dispersion of RGO in the matrix, as reflected by the distribution of secondary phase. Therefore, a possible mechanism of the phase separation can be proposed as follows:

Sr1z Ti1x Nbx O3 /Sr1z Tið1xÞyð1xÞ Nbxyx O32y þ yTi1x Nbx O2

(1)

It can be seen that the reducing agents including RGO and H2/Ar

gas are not directly involved in the reaction, and the separation of TiO2 is mainly ascribed to relatively high concentration of Sr vacancies in the pristine STNO. The similar phenomenon has also been observed in reduced A-site-deficient perovskite with nominal composition of Sr0.93Ti0.9Nb0.1O3 [21]. However, the reducing agent can stimulate the generation of oxygen vacancy that may accelerate the phase separation. In fact, the secondary phase of Nb rich TiO2 has also been found in Sr0.8La0.067Ti0.8Nb0.2O3 ceramic prepared via solid reaction in 5% H2/Ar gas [22], which was ascribed to the result of reduction in the study. Since RGO is more efficient in generating oxygen vacancy compared with the H2/Ar forming gas [19], it is understandable that the STNO composites incorporated with RGO

Fig. 3. SEM images of polished surface for (a) S0, (b) S1, (c) S2 and (d) S3. The white circles indicate the representative area for separated TiO2 phase in each sample.

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have much higher TiO2 content, which could also partly explain the Nb segregation in TiO2 precipitate since the doping level of Nb in STO is limited. Meanwhile, the GO itself should be reduced during heat treatment and sintering processes, which is revealed by Raman spectroscopy. As shown in Fig. 4a, the Raman spectrum of RGO/STNO composite shows a peak at ~1320 cm1 (D band) corresponding to vibrations of disordered carbon atoms in graphite and a peak at ~1592 cm1 (G band) that is associated with the vibration of sp2-bonded carbon in a 2D hexagonal lattice, which confirms the existence of RGO in the STNO composite. Compared with pristine GO, the Raman spectrum of RGO/STNO composite exhibits an increased D/G intensity ratio, which is usually observed in the reduction of GO when the defect concentration is high [23,24]. In addition, the reappearance of 2D band at 2700 cm1 and D þ G band at around 2900 cm1 is also indicative of reduction, since these bands become invisible when the defect concentration is too high, such as the situation for GO [25,26]. In order to further evaluate the reduction level of RGO, XPS analysis was conducted on the pristine GO and RGO/STNO composite (S2). As shown in Fig. 4b, the C1s peak for pristine GO can be fitted into four peaks, and assigned to CeC (284.6 eV), CeO (285.8 eV), C]O (287.3 eV), OeC] O (288.5 eV). It is clear that the CeC peak increased, while the CeO and OeC]O peak became weaker and the main oxygen-containing peak of C]O disappeared after heat treatment in H2/Ar and SPS, which results in a largely decreased O/C ratio in RGO (from 59 to 18%). All these observations deliver evidence that GO has been partially reduced by heat treatment in H2/Ar and SPS, which is in agreement with the Raman spectra. The detailed microstructure of RGO/STNO ceramic composite was further investigated by TEM and STEM, as shown in Fig. 5. A fully dense microstructure with gain size varying from submicron to several microns can be seen by the bright field image. According to HRTEM observation (Fig. 5b), a RGO flake with thickness around 4 nm at the triple junction of STNO gains can be observed. Compared with the pristine GO (~700 nm), the size of the RGO in composite is relatively small (~100 nm), which can be ascribed to the reaction between RGO and STNO matrix. The second phase of Nb rich TiO2 is also confirmed by EDS mapping (Fig. 6) of S2 sample. The selected area electron diffraction clearly indicates a rutile structure of Nb doped TiO2, which is consistent with the XRD results. However, we also found that the presence of some fine SrOx phase which is not revealed by XRD and SEM images, though these precipitates may be beneficial to enhance the phonon scattering and reduce the lattice thermal conductivity.

3.2. Electrical properties of the STNO/RGO composites Theoretically, SrTiO3 is an insulator with a wide band gap (3.2 eV) [27]. Doping Nb5þ ions on Ti4þ site can donate the electrons to increase the electrical conductivity. However, reduction is generally necessary to guarantee the doping is compensated by electrons rather than Sr vacancy, which is the reason that the reduction in H2/Ar gas is included in the processing method. In fact, without the reducing process of annealing in H2/Ar gas, the direct sintered composite shows much lower electrical conductivity, since the RGO could only influence the adjacent area in the matrix (Fig. S3). The temperature dependence of electrical conductivity ðsÞ for STNO and RGO/STNO composite with different RGO contents are plotted in Fig. 7. As show in Fig. 7a, the electrical conductivity for STNO composites with different contents of RGO shows similar temperature dependence. The electrical conductivity of all samples increases with the temperature at first and turns down at a peak value between 500 and 600 K, implying different carrier scattering mechanism at low and intermediate temperature [13]. Furthermore, the electrical conductivity is significantly improved via the addition of RGO. The maximum electrical conductivity for S2 is 250.08 Scm1 at 550 K, which is 2.8 times higher than that for S0. Therefore, the first issue that should be clarified is where the improvement of electrical conductivity stems from. Apparently, in the study, both the separated TiO2 phase and RGO could influence the electrical conductivity. Similar to the STNO, the TiO2 can become conductive by doping with Nb, but most of the electrical conductivity values reported so far are much inferior to the values for the composites obtained here [28e30]. For comparison, the 10 at.% Nb doped TiO2 ceramic was prepared via solid reaction followed by the completely same procedure of annealing and sintering for preparing STNO ceramics. It is found that the electrical conductivity of the Nb doped TiO2 is less than 40 Scm1 in the whole temperature range (Fig. S4), which rules out the Nb rich TiO2 phase as the origin of highly increased electrical conductivity. Therefore, it is deduced that the improvement of electrical conductivity is related to RGO rather than the separated TiO2 phase, even though the RGO phase has much less content in the composites. To further investigate the mechanism behind the enhanced s, the carrier concentration (n) and carrier mobility (m) at room temperature were determined by the Hall measurement for the monolithic STNO and RGO/STNO composites. As shown in Fig. 7b, the n for S1, S2 and S3 is ~0.9105  1020, ~1.059  1020,

Fig. 4. The reduction level of RGO in the composite (S2) compared with GO. (A) Raman spectra (B) High-resolution C 1s XPS spectra.

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Fig. 5. TEM and HRTEM images of S2 sample. (a) TEM images at low magnification. (b) HRTEM images showing RGO at the triple junction from the area of white square in (a), inset: the magnified area of dash lined square in (b).

Fig. 6. TEM images and EDS mapping of the area containing separated TiO2 phase in S2. (a) dark filed TEM image. (b) selected area electron diffraction from the spot indicated by white circle in (a). (cef) EDS elemental mapping for Sr, Ti, Nb and O, respectively.

~0.7497  1020 cm3 respectively, which is much higher compared with S0 (1.866  1019 cm3). Meanwhile, the m for S1, S2 and S3 is also magnitude higher than that for S0 (0.3159cm2V1s1). Moreover, both n and m increase with the increasing RGO content from sample S0 to sample S2, and then decrease for sample S3. The increased carrier density can be ascribed to the increased reduction level in the matrix upon addition of RGO. The reduction in STNO leads to the loss of oxygen and generation of electrons in compensation, which can be described by the following equation [31,32]:

  0 0 1 0 2TixTi þ Oxo 42TiTi þ V::o þ O2 2TiTi ¼ 2e 2

(2)

where the product should be CO2 and H2O instead of O2 in the presence of RGO and H2. Therefore, beside the effect on phase separation, RGO can also enhance the carrier density by promoting donor-type oxygen vacancy. Meanwhile, it should be noted that another consequence of adding RGO is the inhibition of grain growth (Fig. 8), which created large amount of grain boundaries

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Fig. 7. Transport properties for STNO and RGO/STNO composites. (a) Electrical conductivity. (b) Carrier concentration and mobility. (c) Seebeck coefficient. (d) Power factor.

Fig. 8. EBSD analysis showing grain size distribution and average grain size for (a) S0, (b) S1, (c) S2 and (d) S3.

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that may act as double Schottky barriers (DSB) to trap electrons in STNO [33,34]. As a result, when RGO content surpasses certain amount, the reduction effect of RGO reaches a limit due to the limited contact surface between RGO and STNO, while the trapping effect of increasing grain boundaries could start to prevail, leading to decreased n in the composite. Concerning the increased mobility, it should be related to the effect of RGO on the grain boundary as well. It has been reported in RGO/Al doped ZnO composite that the addition of RGO could highly boosting the carrier mobility to the level of single crystal, due to the significantly weakened boundary barrier and proper band alignment between RGO and ZnO [35]. Likewise, the reduction effect from RGO can also greatly lower the DSB adjacent to the RGO, which leads to highly improved carrier mobility. However, the increased amount of grain boundary with increasing RGO content will finally result in decreased mobility due to the strengthened scattering effect at grain boundaries, which is in consistence with the observation for sample S3. It is worth noting that the similar transportation behavior has been observed in RGO/Zn0.98Al0.02O composite, in which the n and m first increase and then decrease with increasing filler fraction, although the authors improperly attribute the decreased carrier density to the ptype conduction of RGO [36]. Fig. 7c shows the temperature dependence of the Seebeck coefficient (S) for the STNO and RGO/STNO composite with different GO contents. The S for all samples is negative in the whole temperature range, indicating that electron is the majority charge carrier. The absolute value jSj increases gradually with increasing temperature from 300 to 800 K, which is in agreement with the behavior of degenerate semiconductor [37]. The Seebeck coefficient of degenerate semiconductor closely relate to the carrier concentration (n) and effective mass m* at determined temperature,

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which can be explained as the following equation:

S¼

8p2 K 2B 2

3eh

m* T

 p 23 3n

(3)

where kB is Boltzmann constant, h is Planck constant, e is the carrier charge, and T is the absolute value. Compared with the STNO, the Seebeck coefficient of the RGO/STNO composite with different GO contents is decreased, mainly owing to the increased carrier concentration. The effective mass m* is also slightly changed with increasing RGO content (Table 1), which can be understood as a comprehensive effect from the changes in lattice structure of STNO and separated TiO2 phase. Since the variance of m* for different samples is not large, there should be no fundamental change in the band structure of STNO and the carrier concentration should play the dominant role in Seebeck coefficient, which is consistent with the trend observed for jSj at room temperature. Consequently, the corresponding power factor (PF ¼ sS2 ) as a function of temperature were calculated as presented in Fig. 7d. It can be seen that although monolithic STNO has higher PF before 600 K, the RGO/ STNO composites show higher PF at temperature higher than 600e650 K, owing to the much higher electrical conductivity. The maximum PF obtained for S2 is about 0.8 mW m1K2 at 800 K, which is greatly higher than that of the monolithic STNO and RGO/ SrTiO3 composite [19]. 3.3. Thermal transport property of the RGO/STNO composites The total thermal conductivity (k) as a function of temperature for various samples is shown in Fig. 9a. The total thermal conductivity for all samples shows similar trend of decreased value with

Fig. 9. Thermal properties for STNO and RGO/STNO composites. (a) Total thermal conductivity. (b) Lattice thermal conductivity. (c) The lattice thermal conductivity of 10 at.% Nb doped TiO2 and TiO2d (The data for the TiO2d samples is from ref. 36) (d) the variation of separated TiO2 phase content obtained by the statistical analysis from SEM images.

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increasing temperature at 300e800 K. It is found that the k for the monolithic STNO is similar to the reported SrTiO3 based ceramics fabricated via hydrothermal method [12]. Remarkably, even with the lowest loading of RGO (S1), the thermal conductivity can be largely decreased in the whole investigated temperature range. As the content of graphene further increased, the thermal conductivity becomes even lower for samples S2 and S3, which is ~23% smaller than that of S0 at 300 K. After extracting the lattice thermal conductivity (kL) from k by the expression kL ¼ k-ke using WiedemannFranz law:

ke ¼ sLT

(4)

where L is the Lorenz number (2:44  188 V 2 K 2 ), it is found that the influence of electron thermal conductivity is very weak for the all samples, indicating that the total thermal conductivity mainly depends on the lattice thermal conductivity. However, since the S2 has larger electrical conductivity in comparison to S3, the S2 sample shows lowest kL among the composites (Fig. 9b). It has been shown in our previous study that the addition of RGO can effectively reduce the lattice thermal conductivity of SrTiO3 ceramic through restraining the grain growth [19]. Therefore, it is natural to infer that the depressed kL here is due to the same reason. However, according to the SEM observation and statistic result from EBSD (Fig. 8), the average of grain size for monolithic STNO is actually quite fine compared with the monolithic SrTiO3. Although the average grain size (d) is inversely proportional to the RGO content, the difference is not as large as the situation for RGO/SrTiO3 composite with respect to pure SrTiO3 (d¼558 nm for S0 and 490 nm for S1). The extremely fine microstructure in monolithic STNO can be ascribed to the reduction effect before sintering, which has been demonstrated by Chung et al. in 2.4 atom% Nb doped SrTiO3 [38]. Thus, the largely decreased kL cannot be attributed to decreased grain size alone, noting that S3 with the finest grain size shows slightly larger kL compared with S3. Similarly, the greatly depressed kL cannot be solely ascribed to RGO as well. Although the RGO induced thermal resistance could play a role in decreasing kL [39]. Based on these facts, it is deduced that the deceased kL should be mainly related to the separated Nb rich TiO2 phase. In order to clarify this speculation, the thermal conductivity of 10 at.% Nb doped TiO2 ceramic was investigated (Fig. 9c). It is found that the kL of 10 at.% Nb doped TiO2 ceramic is much lower than that of STNO and undoped TiO2 [40]. At 300 K, the thermal conductivity of 10 at.% Nb doped TiO2 ceramic is 2.27 Wm1 K1, which is about 3.0 times lower than the reported value for undoped TiO2 (7.0 Wm1 K1 ). Therefore, the Nb rich TiO2 with very low intrinsic thermal conductivity could serve as nanodispersion to inhibit the phonon transport in RGO/STNO composites. In addition, according to the statistical analysis from SEM (Fig. 9d), the content of Nb rich TiO2 phase increases with increasing RGO until 12.51% for S2 and then decreases for S3, which is in very good agreement with the trend of lattice thermal conductivity. As a result, the S1, S2 and S3 composites show remarkable enhancement of ZT compared with S0 (Fig. 10). The ZT value for S2 reaches a maximum of ~0.22 at 800 K, which is 1.8 times than that of S0. Moreover, the ZT value is also higher than the reported data for other A-site deficient and Nb doped SrTiO3 prepared by solid reaction (ZT ¼ 0.2 at 800 K in Sr0.95Ti0.9Nb0.1O3 [21] and Sr0.94Ti0.8Nb0.2O3 [41]). This improvement of thermoelectric performance can be ascribed to three factors: i) incorporation of RGO into STNO can improve carrier concentration via production of oxygen vacancies; ii) RGO at the grain boundary can depress DSB to enhance carrier mobility; iii) the thermal conductivity of RGO/ STNO composites is also highly decreased mainly due to the

Fig. 10. The temperature dependence of ZT values for S0, S1, S2 and S3.

separated Nb rich TiO2 second phase promoted by RGO. 4. Conclusions In summary, a Nb-doped A-site-deficient SrTiO3 composite incorporated with RGO has been successfully fabricated via a facile hetero-aggregation method followed by annealing in reducing atmosphere and SPS. It is found that the addition of RGO profoundly changed the microstructure of STNO ceramics, leading to separated rutile TiO2 phase with enriched Nb. Further investigation reveals that the highly enhanced electrical conductivity is attributed to the presence of RGO which increases carrier concentration by the reduction effect, and increases carrier mobility via lowering DSB at grain boundaries. Meanwhile, the Nb rich TiO2 second phases also give rise to greatly decreased lattice thermal conductivity, due to the extremely low thermal conductivity of Nb rich TiO2 and the scattering effect as nanoscale inclusion. Consequently, RGO/STNO composite shows largely enhanced thermoelectric properties compared with the monolithic STNO, leading to a ZT value of ~0.22 at 800 K. Therefore, incorporation of graphene appears to be a very effective and promising route for modulating and improving the thermoelectric properties of A-site deficient perovskite. Acknowledgements This work was funded by the Fundamental Research Funds for the Central Universities (2232017A-07), Natural Science Foundation of China (Nos.51432004, 51532006, 51774096, 51871053), Shanghai Committee of Science and Technology (No. 16JC1401800, 18JC1411200), DHU Distinguished Young Professor Program, Science and Technology Commission of Shanghai Municipality (16DZ2260601) and the 111 Project (D16002). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2019.01.376. References [1] G.J. Snyder, E.S. Toberer, Complex thermoelectric materials, Nat. Mater. 7 (2008) 105e114.

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