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Graphene promoted oxygen vacancies in perovskite for enhanced thermoelectric properties
Carbon 112 (2017) 169e176
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Graphene promoted oxygen vacancies in perovskite for enhanced thermoelectric properties Xiaopeng Feng a, Yuchi Fan b, *, Naoyuki Nomura a, Keiko Kikuchi a, Lianjun Wang c, Wan Jiang b, c, Akira Kawasaki a, ** a b c
Department of Materials Processing, Graduate School of Engineering, Tohoku University, Sendai, 9808576, Japan Institute of Functional Materials, Donghua University, Shanghai, 201620, China State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, Donghua University, Shanghai, 201620, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 September 2016 Received in revised form 2 November 2016 Accepted 7 November 2016 Available online 9 November 2016
Oxygen vacancy plays an important role in optimizing the thermoelectric properties for achieving high figure of merit in perovskite oxides. However, it remains very challenging to find a facile and effective method for creating oxygen vacancy in bulk perovskite material. Using undoped strontium titanate (STO) as a model for perovskite, it is demonstrated here that incorporating graphene into STO can promote the formation of oxygen vacancy in perovskite far more effectively than the time and energy consuming hydrogen reduction method. With only 0.64 vol% of graphene content, reduced graphene oxide (RGO)/ STO composite shows highly increased electrical conductivity and power factor compared to hydrogen reduced STO. Electron energy loss spectrum confirms the high concentration of oxygen vacancy in the area close to RGO in RGO/STO composite, which suggests a mild reaction between RGO and STO during sintering. In addition, the thermal conductivity is also depressed due to greatly restricted grain growth via adding RGO. Therefore, this study provides a fast, green and effective way of preparing oxygen deficient perovskite towards improved thermoelectric properties. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction With the ability of converting heat into electricity, thermoelectric materials have been intensively studied in the past decade due to the growing demand for energy and lack of solution for sustainable powder generation. Theoretically the efficiency of fuels can be largely improved when the waste heat is recycled via the application of thermoelectric materials with large dimensionless figure of merit, ZT ¼ sS2T/k, in which s is the electrical conductivity, S is the Seebeck coefficient, k is the thermal conductivity and T is the absolute temperature. Although numbers of high ZT materials such as Te based and Se based skutterudite and intermetallic compounds have been developed so far, most of these materials are suffered from toxicity, rarity in earth and instability [1,2]. In the regard of this situation, many researchers have revisited the possibility of using oxides as thermoelectric materials, despite the fact
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Fan), [email protected]. jp (A. Kawasaki). http://dx.doi.org/10.1016/j.carbon.2016.11.012 0008-6223/© 2016 Elsevier Ltd. All rights reserved.
that oxides are of low carrier mobility and high thermal conductivity in general. The representative p-type materials are NaCo2O4 and Ca3Co4O9 which have been found very promising especially in high temperature range, benefiting from their layered structure [3e6]. For n-type thermoelectric materials, SrTiO3 (STO) has attracted tremendous interests for its large carrier effective mass, stability at high temperature and structural tolerance with respect to doping. As a typical perovskite with composition of ABO3, STO is an intrinsic insulator without doping. Electrical conductivity can be rendered either by introducing foreign dopants to A-or B-site, or by reduction to generate oxygen vacancies. The former route is much more popular because of its obvious flexibility for tailoring properties. STO ceramics with A-site partly substituted by rare earth elements like Y, La, Sm, Gd and Dy have comparatively investigated by Muta et al., whose work suggests that the influence of A-site doping on ZT value mainly comes from the changes in thermal conductivity rather than power factor [7]. Moreover, Ohta et al. have shown that in comparison to A-site doped Sr1-xLaxTiO3, B-site doped SrTi1xNbxO3 single crystal exhibit larger power factor due to even larger density of states effective mass (m*), indicating the potential
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of Nb doping [8]. In their later work, polycrystalline ceramic of SrTi1-xNbxO3 was also fabricated. Albeit same level of power factor maintained compared to the corresponding single crystal film, thermal conductivity was not reduced effectively by grain boundaries, leading to a ZT value of 0.35 at 1000 K which is probably the highest level can be reached presently for STO ceramics at high temperature [9]. In contrast to doping with foreign atoms, oxygen deficient STO has not attracted sufficient attention so far even though this phenomenon occurs commonly when samples are fired in reductive atmosphere such as in vacuum or reducing gas. However, a few research groups have already noticed that oxygen vacancy can play an important role in optimizing the thermoelectric properties of doped STO. Yu et al. reported significantly reduced thermal conductivity in oxygen deficient La-doped STO film prepared by pulsed laser deposition (PLD) under a high vacuum of ~107 Torr at 923 K. The low thermal conductivity was mainly attributed to oxygen vacancies that suppressed heat transfer effectively [10]. Similar low thermal conductivities were obtained when PLD was conducted under a vacuum of 108 Torr at 1153 K for La-doped STO. More interestingly, it seems that the introduction of oxygen vacancies can improve the electrical conductivity by increasing carrier density without significant sacrifice of Seebeck coefficient when the A-site doping level is low (5%La), highlighting the great potential of oxygen deficient STO [11]. Apart from epitaxial films, bulk ceramics of STO with oxygen deficiency were reported as well, usually prepared by annealing in reducing gas (i.e. H2/Ar forming gas) at elevated temperature after sintering [12e15]. Nevertheless, gradient oxygen vacancy distribution in the sample cannot be avoided generally owing to the diffusion controlled process from surface to inside. Therefore more effective approach to prepare oxygen deficient STO bulk ceramics is of great importance to the enhancement of thermoelectric properties and application of STO based materials, since polycrystalline ceramics have unarguable advantage in cost in comparison to single crystals and chemically or physically deposited films. In this work, we demonstrate a novel and facile method for creating oxygen vacancy in perovskite ceramics using undoped STO as a model material. Instead of post-sintering heat treatment in reducing atmosphere, oxygen deficiency in STO was generated during spark plasma sintering (SPS) process by addition of uniformly dispersed graphene oxide (GO) in STO powders. The reduced GO (RGO) at elevated temperature served as carbon source that further reduced the STO at sintering temperature of 1473 K, leading to highly increased concentration of oxygen vacancy and residual graphene platelets as inclusion in the final STO compact. The attained RGO/oxygen deficient STO composite shows largely improved power factor and suppressed thermal conductivity compared with oxygen deficient STO prepared by long-time reduction in hydrogen, giving rise to a ZT value of 0.09 at 760 K in the absence of A- or B-site doping. 2. Experimental 2.1. Surface modification of STO The original Strontium Titanate powders (Sigma-Aldrich Co. LLC. Japan, purity of 99%, average size of 100 nm) were surface modified by (3-Aminopropyl) trimethoxy-silane (APTS, SigmaAldrich Co. LLC. Japan, purity of 97%). Firstly, STO powders were added into toluene (Wako Pure Chemical Industries, Ltd, Japan, purity of 99.99%) and dropped 4 ml APTS into the solution of STO. Then the mixture was heated to 423 K and holding for 6 h in the atmosphere of Ar with the reflux devices. The products were washed by toluene and filtered by vacuum filtration followed by
drying at 333 K. Here modified STO powders were obtained. 2.2. Preparation of RGO/STO composite Graphene oxide (GO) was prepared by the modified Hummers method [16]. STO (2.4 g) powders were poured into 500 ml deionized water and the pH was adjusted to 3.3. Through the ultrasonic treatment for 2 h, the colloid of STO formed. The colloid of GO was dropped to the solution of STO by the method of titration with stirring. The mixture was filtered and dried at 333 K. GO/STO hybrid powders were densified by Spark Plasma Sintering (SPS) apparatus (Dr. Sinter 511-S, Sumitomo Coal Mining Co. Ltd. Japan) with the graphite dies under the pressure of 60 MPa in the vacuum atmosphere. The sintering temperature was controlled from room temperature to 1473 K at the heating rate of 100 K/min, and holding for 5 min at 1473 K. During sintering, GO was reduced to RGO. Finally, the disk of RGO/STO composite with a diameter of 10 mm and a thickness of 3 mm was obtained. Al2O3/0.6 vol% RGO composite, as well as RGO/STO composite (0, 0.64 vol%) were prepared by the same method. For comparison, the pure STO ceramic was also sintered using the same condition. Oxygen deficient STO prepared by traditional method was prepared by heat treatment of pure STO ceramic in H2/Ar flowing gas at 1573 K for 10 h. 2.3. Characterization Zeta potentials of as-prepared GO colloid and STO colloid were measured using Zetasizer Nano ZS (Malvern, UK). This instrument utilizes Laser Doppler Micro-electrophoresis technique to measure zeta potential with a red laser (633 nm). The pH values of both samples were adjusted by MPT-2Autotitrator (Malvern, UK) using HCl and NaOH solutions (0.01 M). Before the measurement, both samples were sonicated for 0.5 h to achieve homogeneous dispersion. The carbon content in the composite was measured by combustion infrared detection technique (844 series, LECO, MI, US). The transmission electron microscopy (TEM) observation, electron energy-loss spectroscopy (EELS) and low-angle annular dark field (LAADF) were performed using a JEOL JEM-ARM200F electron microscope. The electron back scattered diffraction (EBSD) was performed using JEOL JSM-7100F. The field-emission scanning electron microscopy (SEM) observation was performed by a JEOL JSM-6500F scanning electron microscope. Micro-Raman spectra were measured with a SOLAR TII Nanofinder (Tokyo Instruments Co.) with 532 nm wavelength incident laser light and a 100 objective. The X-ray diffraction (XRD) patterns were collected by Rigaku Smartlab 9kw diffractometer with Cu Ka radiation source (l ¼ 0.15418 nm). The electrical conductivity was measured by a static dc method based on the slope of the voltage versus the temperature-difference curves using ZEM-3 (Ulvac-Riko) under a low-pressure helium atmosphere. The carrier density and the mobility were measured by a Hall measurement system (Toyo ResiTest8400 series) at room temperature in the air. The thermal diffusivity was measured by a laser flash method (LFA 457, Netzsch), and the specific heat was measured using a differential scanning calorimeter (DSC 404F3, Netzsch). 3. Results and discussion 3.1. Preparation of RGO/STO composites To set a model for the effect of graphene on thermoelectric properties of STO composites, the strategy adopted here for preparing the composite should meet several requirements: i) the quality of composite prepared by this strategy should be high,
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which means the uniform dispersion and absence of agglomeration of graphene in matrix are expected. ii) Though undoped STO was selected as a model material to study here, the strategy should be universal for doped STO and other types of perovskite so that the conclusions drawn here can be easily transfer to those materials as guidance for designing novel composites with improved thermoelectric properties. Based on these considerations, GO (precursor of graphene) and STO powders were homogeneously mixed by hetero-aggregation before densification. This strategy, which is especially effective when oxide particles can form large enough positive surface charge (zeta potential) in water to attract the negatively charged GO, has been successfully applied in many graphene/oxide composites in our previous researches [17e19]. Unfortunately, STO powders cannot form such stable colloid naturally because of the extremely low pzc (point of zero charge) of STO particles. To render positive surface charge to STO particles, NH2þ was grafted onto the surface of STO particle by reacting with (3Aminopropyl) trimethoxy-silane. Upon the surface modification, STO powders became positively charged in water and a maximum zeta potential of 45 mV can be obtained at pH of 3.3, which guarantees the formation of stable colloid (Supplementary Material Fig. S1). After mixing STO and GO colloids carefully, GO/STO hybrid powders with uniform distribution of GO were attained via a classic process of hetero-aggregation, in which each GO nanosheet is fully decorated by STO nano-particles, effectively hindering the recombination and agglomeration of GO nanosheets (Supplementary Material Fig. S3). The GO/STO hybrid powder was densified by SPS to attain fully dense RGO/STO composite, as shown in Table 1. For comparison, an oxygen deficient STO ceramic sample (denoted as STO-H) was also prepared by heat treatment of SPS sintered pure STO ceramic in H2/ Ar forming gas, following the traditional process of creating oxygen vacancies. It is worth noting that only the composite containing very low graphene content was prepared in this study, because RGO has been proved to be of very low thermopower that may deteriorate the thermoelectric performance of STO composite with increasing amount [19,20]. Thus it is better to depress the content of graphene in composite as long as uniform distribution of graphene in the composite can be achieved, since the negative effect could be overweighed by positive one such as highly increased electrical conductivity and maintained Seebeck coefficient arising from graphene introduced oxygen vacancy when the fraction of graphene in composite is low. Three different STO samples, including RGO/STO composite, pure STO ceramic and STO-H ceramic, were characterized by X-ray diffraction (XRD) in the first place to confirm the obtained phase. As shown in Fig. 1, the XRD patterns of all samples are in good consistence with cubic perovskite STO (PDF-#35-0734), indicating no phase transition occurred during sintering or reduction process. The absent peak of RGO in composite can be understood by its low content which is determined as 0.64 vol% (see methods) in this particular sample, while their existence as well as recovered structure from GO were unambiguously identified by Raman spectra (Supplementary Material Fig. S4). The refinement of XRD patterns shows that all samples share not only same cubic structure with space group of Pm3m(Table 1), but also very similar lattice
171
Fig. 1. XRD patterns of pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
parameter that is close to the value reported in standard card, which reveals no significant change in structure with respect to stoichiometric STO in these materials. In other words, despite the probable presence of oxygen vacancies in these STO ceramics owing to the more or less reducing process they are subjected to, the concentration of the defect is modest, in comparison to the highly oxygen deficient STO that adopts a lower symmetrically tetragonal structure [21]. It has been known that one of the most prominent effects of RGO on the microstructure of ceramic composite is the highly restrained grain growth [18], which has critical meaning for reducing thermal conductivity in thermoelectric materials. As demonstrated by electron back scattering diffraction (EBSD), without RGO the average grain size for pure STO is 13.31 mm with grain size distribution mainly ranging from 10 to 18 mm (Fig. 2a). For STO-H, the average grain size is even larger compared with pure STO, reaching 14 mm with a broader grain size distribution, due to the long-time annealing process after sintering (Fig. 2b). On the contrary, the average grain size of the composite is merely 0.58 mm for RGO/STO composite (Fig. 2c), which is more than 20 times smaller than that of pure STO and STO-H, despite the fact that the RGO content is only 0.64 vol% in composite. More interestingly, although the grains in composite are mostly concentrated on 0.6 mm, there is indeed a mixing of grain size in the composite including gains smaller than 0.2 mm and larger than 1 mm (Fig. 2c and d), which could be beneficial in improving ZT value of composite according to the theory proposed by Zhao et al. [22]. Since the detailed structural information for RGO in composite can hardly be given by EBSD or scanning electron microscopy (Supplementary Material Fig. S5), the microstructure of RGO/STO composite was observed by TEM and HRTEM, as shown in Fig. 3. The fully dense and fine grained structure in composite was reconfirmed by TEM observation (Fig. 3a). Under higher magnification, RGO platelet can be clearly identified owning to its lower
Table 1 Crystal structure, composition and relative density for pure STO, STO-H and RGO/STO composite. RGO content [vol%]
Apparent density [gcm3]
Relative density [%]
Crystal structure
Space group
Lattice parameter [Å]
Pure STO
0
5.07
99.02
Cubic
Pm3m
3.90906 (4)
STO-H
0
5.03
98.24
Cubic
Pm3m
3.90800 (4)
RGO/STO
0.64
5.05
99.02
Cubic
Pm3m
3.91017 (5)
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Fig. 2. EBSD analysis of (a) pure STO, (b) pure STO-H and (c) RGO/STO composite; (d) the grain size distributions for pure STO, pure STO-H and RGO/STO composite, respectively. (A colour version of this figure can be viewed online.)
3.2. Electrical properties of RGO/STO composites Stoichiometric STO is an insulator owing to the large band gap
(3.2eV) [23]. However, at high temperature of above 1273 K, neutral oxygen on regular oxygen site (written asOxo ) of STO is thermodynamically prone to leave the lattice, retaining behind an oxygen vacancy with two electrons [24,25].
Oxo 4Vo€ þ 2e þ 1 2O2 =
contrast compared to surrounding STO matrix in bright field image (Fig. 3b), which allows us to verify the uniform distribution of RGO platelets and the absence of severe agglomerations in the composite as well. Moreover, two types of RGO platelets can be found in the composite: the one existing at grain boundaries and the one included inside the grains (Fig. 3d). The former type is very common that can be readily found in any RGO/oxide composite [17,18]. This type of RGO can effectively pin the movement of grain boundary, which is the main reason for the highly restrained grain growth. However, the latter type that RGO completely enveloped in a grain is very rare and unique, which has never been observed in other RGO/oxide composites to the best of our knowledge, since normally the lateral size of RGO is much larger than STO particles (Supplementary Material Fig. S2) that makes grain boundary completely moving across a whole piece of platelet very difficult [18]. Noting the enveloped RGO platelets all have very small size which looks more like the debris of RGO (Fig. 3c), it is very likely that these small pieces of RGO inside grains are actually a result of the mild reaction between STO and RGO during sintering process. As the other consequence of this reaction, STO lost oxygen to become oxygen deficient STO ceramic, which will be demonstrated in the next section.
(1)
Reducing environment like vacuum can further assist the escape of oxygen, leading to a slightly conductive STO ceramic after SPS (pure STO), although the conductivity of which is generally low and unstable because of very limited and inhomogeneous oxygen vacancies generated in this process. As shown in Fig. 4, after repeating measurement to achieve a steady performance (Supplementary Material Fig. S6), the as-sintered pure STO shows a very low electrical conductivity of 17.9 Sm1 at room temperature, which increases slightly with increasing temperature, reaching 176 Sm1 at 760 K finally. To promote more oxygen vacancies for tuning thermoelectric properties of perovskite materials, researchers have to employ stronger reducing conditions usually involving long time annealing in hydrogen gas [13,14], though the process is both time and energy consuming. Following this traditional strategy, the electrical conductivity of STO-H shows much improved electrical conductivity of 640 Sm1 at room temperature and 2025 Sm1 at 570 K (maximum value), revealing considerably increased oxygen vacancies compared with pure STO. Moreover, it is very surprised that without any additional reducing processes, the electrical conductivity of as-sintered RGO/STO composite is even higher than
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Fig. 3. TEM and HRTEM images of RGO/STO composite. (a) and (b) are TEM images at low and high magnification, respectively. (c) and (d) are HRTEM images from the area indicated by black square and black dash line square respectively, showing RGO on the grain boundary (c) and inside the grain (d).
that of STO-H, manifesting a conductivity of 1633 Sm1 at room temperature, and reaching the maximum 4669 Sm1 at 470 K. Theoretically, the improvement of electrical conductivity may arise from the increase in carrier mobility (m) or carrier density (n), or both of them because the conductivity can be written as:
s ¼ mne
(2)
Therefore, to clarify the reason behind the largely enhanced electrical conductivity in RGO/STO composite, it is necessary to figure out which component leads to the increase most. Using the hall measurement, the carrier density and correspondingmobility at room temperature were determined for pure STO, STO-H and composite (Table 2). It is found that in contrast to the pure STO, the STO-H has both increased carrier density and mobility, owing to the increased oxygen vacancies induced by reduction. More strikingly, the carrier density in RGO/STO composite (1.59 1020cm3) is about 3.5 times higher than STO-H (4.54 1019cm3), while the mobility is only slightly larger compared to STO-H. Therefore, it can be concluded that the improvement of the electrical conductivity in RGO/STO composite mainly comes from the increased carrier density rather than mobility, which is totally different from the situation in graphene/doped ZnO composite reported recently [26].
However, it is still not clear what leads to the highly increased carrier density, albeit apparently the reason should be related to RGO, since the only difference between pure STO and composite is the addition of RGO. At first glance, one may ascribe the enhanced electrical conductivity to the high carrier concentration of RGO. Indeed, graphene is famous for high electrical conductivity [27], but it is questionable whether RGO has such capability that can boost the carrier density of composite by more than 40 times with a loading of only 0.64 vol%. In order to clarify this problem, a RGO/ Al2O3 composite with similar graphene loading (0.6 vol%) was fabricated using the identical method as a special reference. Unlike STO, Al2O3 is a very stable insulator even under very harsh circumstances, which makes RGO the only conductive component in RGO/Al2O3 composite. At room temperature, the Al2O3/RGO composite shows an electrical conductivity of only 84.2 Sm1 (Fig. S7) and corresponding carrier density of 3.92 1018cm3 (Table S1) which is two orders lower than the carrier density in RGO/STO composite. Even if we assume only 10% RGO formed effective conducting path in RGO/Al2O3 composite, the carrier density that can be induced by ~0.6 vol% of RGO could not be higher than 3.92 1019cm3. Therefore, rather than the high electrical conductivity of RGO, there must be another reason accounting for the
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to grain boundary. In contrast, the O-K edge on the grain boundary lost details at higher energy loss in the composite, which is consistent with the observation for oxygen deficient STO [28e30]. The Ti-L edge can provide more powerful evidence since the explicit distinction between Ti3þ and Ti4þ formal valences [28,31,32]. Because every oxygen vacancy can nominally donate two electrons to nearest Ti d band and the number of oxygen is three times more than titanium, the sensitivity of Ti is six times higher than O. It is found that in RGO/STO composite, the EELS spectra of Ti-L edge for grain boundary with RGO are distinctly different from that for inner grain without RGO (no carbon signal in the area), firmly proving the high concentration of Ti3þ ions close to RGO. Besides, the LAADF imaging can also provide evidence of RGO promoted oxygen vacancies (Fig. S8), since the LAADF signal can detect the strain fields from oxygen vacancies through their distortion of the adjacent cation sites and subsequent dechannelling of the electron beam from the cation columns [28]. Therefore, we can conclude that the amount of oxygen vacancy in composite that is mainly concentrated at the vicinity of RGO is much higher compared with pure STO. Considering the small pieces of RGO inside STO grains, the reason behind the highly increased carrier density for RGO/STO composite becomes much clear: RGO in composite served as carbon source that promoted the formation of oxygen vacancy in STO composite via mildly reacting with oxygen atoms on the surface of STO grain. During this process, some RGO platelets finally lost adequate size to hinder the movement of grain boundary and were enclosed by STO grains.
3.3. Thermal conductivity and ZT value The Seebeck coefficient of pure STO, pure STO-H and RGO/STO composite with the function of temperature is plotted in Fig. 4b. Compared with pure STO, the Seebeck coefficient of RGO/STO composite is just decreased by 16.6%, despite the two orders higher electrical conductivity. For degenerate semiconductors, two significant parameters that are responsible to the Seebeck coefficient are carrier density n and effective mass m* at determined temperature, which can be expressed as the following equation [33]:
. S ¼ 8p2 k2B 3eh2 m Tðp=3nÞ2=3
Fig. 4. (a) Electrical conductivity, (b) Seebeck coefficient, and (c) power factor of pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
high carrier density in RGO/STO composite, which is very likely to be the oxygen vacancy in RGO/STO composite. To prove our speculation, Electrical Energy-loss Spectroscopy (EELS) was employed to investigate the electronic structure change in RGO/STO composite and pure STO. As shown in Fig. 5, the O-K edge fine structure in pure STO is completely same from inner grain
(3)
where kB is Boltzmann constant; h is Planck constant; e is the carrier charge; and T is the absolute value. Therefore, the increase of carrier density usually results in the decrease of Seebeck coefficient if the effective mass is unchanged. However, with the introduction of oxygen vacancies, the effective mass increased for STO (Table 2), thus largely retaining the Seebeck coefficient. As a result, the power factor of RGO/STO composite is one order higher than that of pure STO and almost 2 times higher than that of STO-H at 520 K, shown in Fig. 4c. This large value of power is even comparable with dopedSTO such as Pr-doped STO (8.03 104Wm1 K2 at 1170 K, estimated data from the figures in the paper) [34]. The thermal diffusivity (D), specific heat (Cp) and density (d) were measured to obtain the total thermal conductivity, which is showed in Fig. 6a. The electronic thermal conductivity (ke) was
Table 2 Carrier density, mobility and effective mass for pure STO, STO-H and RGO/STO composite at room temperature.
Pure STO STO-H RGO/STO a
m0 is the mass of free electron.
Carrier density (n) [cm3]
Carrier mobility (m) [cm2V1s1]
Effective mass (m*)a
3.64 1018 4.54 1019 1.59 1020
0.29 0.49 0.64
0.41m0 2.41m0 4.39m0
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175
Fig. 5. EELS spectra for pure STO and RGO/STO composite. (a) (c) Ti-L2,3 edge, the Ti-L edge shifts from a þ4 to þ3 valence with increasing oxygen vacancies, as the oxygen vacancies are electron donors. (b) (d) O-K-edge, the O-K edge on the grain boundary lost details at higher energy loss spectrum with the introduction of oxygen vacancies. (A colour version of this figure can be viewed online.)
separated from lattice thermal Wiedemann-Franz law [35]:
ke ¼ sLT
conductivity
(kl)
by
the
(4)
where L is the Lorenz number (2.44 108V2K2). It can be seen that the contribution from electronic thermal conductivity is almost negligible in pure STO (Table S2), STO-H and RGO/STO composite, while the lattice thermal conductivity plays an important role. kl for pure STO is 7.7 Wm1 K1 at 370 K and decreased with the elevated temperature, and reached 5.6 Wm1 K1 at 760 K. For STO-H, kl is even higher than pure STO, revealing the influence from larger grain size induced by long-time heat treatment in hydrogen. In contrast, the lattice thermal conductivity of RGO/STO is just 5.7 Wm1 K1 at 370 K and 3.7 Wm1 K1 at 760 K, which decreased by 26% compared with pure STO. This depressed lattice
thermal conductivity in comparison to pure STO can be ascribed to the highly restrained grain size which is 20 times smaller than that of pure STO, thanks to the filler phase of RGO. The temperature dependence of the dimensionless figure of merit (ZT) for pure STO, STO-H and RGO/STO composite is plotted in Fig. 6b. It is found the ZT value of RGO/STO composite is much higher than that of pure STO and pure STO-H, exhibiting a ZT value of 0.03 at room temperature and 0.09 at 760 K, which is 2-fold higher than STO-H at 760 K. This improvement in ZT mainly comes from three factors: i) incorporation of RGO into STO can highly enhance the electrical conductivity of the composite via generating oxygen vacancies, rather than via the high electrical conductivity of itself; ii) the increased electrical conductivity does not deteriorate the Seebeck coefficient due to the increased effective mass of electrons simultaneously; iii) the thermal conductivity
Fig. 6. Temperature dependence of (a) the total thermal conductivity, (b) the lattice thermal conductivity and (c) the ZT value for pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
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of STO composite is also decreased due to the highly suppressed grain growth induced by RGO. 4. Conclusions In summary, we demonstrate a novel, fast and effective method for preparing oxygen deficient perovskite by incorporation with graphene using STO as the model material. Through surface modification of STO and a hetero-aggregation method to obtain uniform mixing between STO powder and GO before densification, a fully dense RGO/STO composite with homogeneous dispersion of 0.64 vol% RGO in matrix has been achieved in this study. It is found that with this low filler content, the composite displays highly enhanced electrical conductivity and much depressed thermal conductivity in comparison to hydrogen reduced STO, leading to a ZT value of 0.09 at 760 K without any foreign dopants. Further investigation reveals that the highly increased electrical conductivity mainly arises from the oxygen vacancy induced by mildly reaction between RGO and STO, which is much more effective than the traditional method via long-time hydrogen reduction. Meanwhile, the thermal conductivity is decreased owning to the highly restrained average grain size induced by RGO. Therefore, incorporating with graphene could be a very effective, convenient and novel method for modulating the thermoelectric properties of perovskite ceramics. Acknowledgements The authors would like to thank the support of China Scholarship Council. This work was funded by Natural Science Foundation of China (No. 51374078，51403037, 51272042), the Opening Project of State Key Laboratory of High Performance Ceramics and Superfine Microstructure(SKL201507SIC), the Fundamental Research Funds for the Central Universities and DHU Distinguished Young Professor Program. We appreciate the generous helps from Dr. WeiWei Zhou and Xiaohao Sun in Kawasaki Lab. Many thanks are also given to Qinghuan Huo, Dr. Kosei Kobayashi and Dr. Takamichi Miyazaki in Tohoku University. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.carbon.2016.11.012. References [1] K. Koumoto, R. Funahashi, E. Guilmeau, Y. Miyazaki, A. Weidenkaff, Y.F. Wang, et al., Thermoelectric ceramics for energy harvesting, J. Am. Ceram. Soc. 96 (1) (2013) 1e23. [2] J. He, Y.F. Liu, R. Funahashi, Oxide thermoelectrics: the challenges, progress, and outlook, J. Mater. Res. 26 (15) (2011) 1762e1772. [3] I. Terasaki, Y. Sasago, K. Uchinokura, Large thermoelectric power in NaCo2O4 single crystals, Phys. Rev. B 56 (20) (1997) 12685e12687. [4] R. Funahashi, I. Matsubara, H. Ikuta, T. Takeuchi, U. Mizutani, S. Sodeoka, An oxide single crystal with high thermoelectric performance in air, Jpn. J. Appl. Phys. Part 2 Lett. 39 (11B) (2000). L1127eL9. [5] A.C. Masset, C. Michel, A. Maignan, M. Hervieu, O. Toulemonde, F. Studer, et al., Misfit-layered cobaltite with an anisotropic giant magnetoresistance: Ca3Co4O9, Phys. Rev. B 62 (1) (2000) 166e175. [6] Y. Miyazaki, K. Kudo, M. Akoshima, Y. Ono, Y. Koike, T. Kajitani, Low-temperature thermoelectric properties of the composite crystal [Ca2CoO3.34]0.614[CoO2], Jpn. J. Appl. Phys. Part 2 Lett. 39 (6A) (2000). L531eL3. [7] H. Muta, K. Kurosaki, S. Yamanaka, Thermoelectric properties of rare earth doped SrTiO3, J. Alloy Compd. 350 (1e2) (2003) 292e295. [8] S. Ohta, T. Nomura, H. Ohta, K. Koumoto, High-temperature carrier transport and thermoelectric properties of heavily La- or Nb-doped SrTiO3 single
crystals, J. Appl. Phys. 97 (3) (2005) 4. [9] S. Ohta, H. Ohta, K. Koumoto, Grain size dependence of thermoelectric performance of Nb-doped SrTiO3 polycrystals, J. Ceram. Soc. Jpn. 114 (1325) (2006) 102e105. [10] C.H. Yu, M.L. Scullin, M. Huijben, R. Ramesh, A. Majumdar, The influence of oxygen deficiency on the thermoelectric properties of strontium titanates, Appl. Phys. Lett. 92 (9) (2008) 3. [11] J. Ravichandran, W. Siemons, D.W. Oh, J.T. Kardel, A. Chari, H. Heijmerikx, et al., High-temperature thermoelectric response of double-doped SrTiO3 epitaxial films, Phys. Rev. B 82 (16) (2010) 5. [12] J. Liu, C.L. Wang, W.B. Su, H.C. Wang, P. Zheng, J.C. Li, et al., Enhancement of thermoelectric efficiency in oxygen-deficient Sr1-xLaxTiO3-d ceramics, Appl. Phys. Lett. 95 (16) (2009) 3. [13] A.V. Kovalevsky, A.A. Yaremchenko, S. Populoh, A. Weidenkaff, J.R. Frade, Effect of a-site cation deficiency on the thermoelectric performance of donorsubstituted strontium titanate, J. Phys. Chem. C 118 (9) (2014) 4596e4606. [14] L. Sagarna, S. Populoh, A. Shkabko, J. Eilertsen, A.E. Maegli, R. Hauert, et al., Influence of the oxygen content on the electronic transport properties of SrxEu1-xTiO3-d, J. Phys. Chem. C 118 (15) (2014) 7821e7831. [15] A.A. Yaremchenko, S. Populoh, S.G. Patricio, J. Macias, P. Thiel, D.P. Fagg, et al., Boosting thermoelectric performance by controlled defect chemistry engineering in Ta-Substituted strontium titanate, Chem. Mat. 27 (14) (2015) 4995e5006. [16] S. Park, J.H. An, R.D. Piner, I. Jung, D.X. Yang, A. Velamakanni, et al., Aqueous suspension and characterization of chemically modified graphene sheets, Chem. Mat. 20 (21) (2008) 6592e6594. [17] Y.C. Fan, W. Jiang, A. Kawasaki, Highly conductive few-layer Graphene/Al2O3 nanocomposites with tunable charge carrier type, Adv. Funct. Mater. 22 (18) (2012) 3882e3889. [18] Y.C. Fan, M. Estili, G. Igarashi, W. Jiang, A. Kawasaki, The effect of homogeneously dispersed few-layer graphene on microstructure and mechanical properties of Al2O3 nanocomposites, J. Eur. Ceram. Soc. 34 (2) (2014) 443e451. [19] Y.C. Fan, L.J. Kang, W.W. Zhou, W. Jiang, L.J. Wang, A. Kawasaki, Control of doping by matrix in few-layer graphene/metal oxide composites with highly enhanced electrical conductivity, Carbon 81 (2015) 83e90. [20] N.D.K. Tu, J. Choi, C.R. Park, H. Kim, Remarkable conversion between n- and ptype reduced graphene oxide on varying the thermal annealing temperature, Chem. Mat. 27 (21) (2015) 7362e7369. [21] W.H. Gong, H. Yun, Y.B. Ning, J.E. Greedan, W.R. Datars, C.V. Stager, Oxygendeficient SRTIO3-X, X ¼ 0.28, 0.17, AND 0.08-crystal-growth, crystal-structure, magnetic, and transport-properties, J. Solid State Chem. 90 (2) (1991) 320e330. [22] L.D. Zhao, B.P. Zhang, W.S. Liu, J.F. Li, Effect of mixed grain sizes on thermoelectric performance of Bi2Te3 compound, J. Appl. Phys. 105 (2) (2009) 6. [23] A. Ganguly, S. Sharma, P. Papakonstantinou, J. Hamilton, Probing the thermal deoxygenation of graphene oxide using high-resolution in situ x-ray-based spectroscopies, J. Phys. Chem. C 115 (34) (2011) 17009e17019. [24] R. Moos, K.H. Hardtl, Defect chemistry of donor-doped and undoped strontium titanate ceramics between 1000 degrees and 1400 degrees C, J. Am. Ceram. Soc. 80 (10) (1997) 2549e2562. [25] N.G. Eror, U. Balachandran, Electrical-conductivity in strontium-titanate with nonideal cationic ratio, J. Solid State Chem. 42 (3) (1982) 227e241. [26] W.H. Nam, B.B. Kim, S.G. Seo, Y.S. Lim, J.Y. Kim, W.S. Seo, et al., Structurally nanocrystalline-electrically single crystalline ZnO-reduced graphene oxide composites, Nano Lett. 14 (9) (2014) 5104e5109. [27] A.K. Geim, K.S. Novoselov, The rise of graphene, Nat. Mater. 6 (3) (2007) 183e191. [28] D.A. Muller, N. Nakagawa, A. Ohtomo, J.L. Grazul, H.Y. Hwang, Atomic-scale imaging of nanoengineered oxygen vacancy profiles in SrTiO3, Nature 430 (7000) (2004) 657e661. [29] S. Hillyard, J. Silcox, Detector geometry, thermal diffuse-scattering and strain effects in ADF stem imaging, Ultramicroscopy 58 (1) (1995) 6e17. [30] D.D. Perovic, C.J. Rossouw, A. Howie, Imaging elastic strains in high-angle annular dark-field scanning-transmission electron-microscopy, Ultramicroscopy 52 (3e4) (1993) 353e359. [31] M. Abbate, F.M.F. Degroot, J.C. Fuggle, A. Fujimori, Y. Tokura, Y. Fujishima, et al., Soft-X-ray-absorption studies of the location of extra charges induced by substitution in controlled-valence materials, Phys. Rev. B 44 (11) (1991) 5419e5422. [32] A. Ohtomo, D.A. Muller, J.L. Grazul, H.Y. Hwang, Epitaxial growth and electronic structure of LaTiOx films, Appl. Phys. Lett. 80 (21) (2002) 3922e3924. [33] P.A. Lee, T.V. Ramakrishnan, Disordered electronic systems, Rev. Mod. Phys. 57 (2) (1985) 287e337. [34] A.V. Kovalevsky, A.A. Yaremchenko, S. Populoh, A. Weidenkaff, J.R. Frade, Enhancement of thermoelectric performance in strontium titanate by praseodymium substitution, J. Appl. Phys. 113 (5) (2013) 8. [35] G.V. Chester, A. Thellung, Law of wiedemann and franz, Proc. Phys. Soc. Lond. 77 (497) (1961) 1005e1013.
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Carbon journal homepage: www.elsevier.com/locate/carbon
Graphene promoted oxygen vacancies in perovskite for enhanced thermoelectric properties Xiaopeng Feng a, Yuchi Fan b, *, Naoyuki Nomura a, Keiko Kikuchi a, Lianjun Wang c, Wan Jiang b, c, Akira Kawasaki a, ** a b c
Department of Materials Processing, Graduate School of Engineering, Tohoku University, Sendai, 9808576, Japan Institute of Functional Materials, Donghua University, Shanghai, 201620, China State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, Donghua University, Shanghai, 201620, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 September 2016 Received in revised form 2 November 2016 Accepted 7 November 2016 Available online 9 November 2016
Oxygen vacancy plays an important role in optimizing the thermoelectric properties for achieving high figure of merit in perovskite oxides. However, it remains very challenging to find a facile and effective method for creating oxygen vacancy in bulk perovskite material. Using undoped strontium titanate (STO) as a model for perovskite, it is demonstrated here that incorporating graphene into STO can promote the formation of oxygen vacancy in perovskite far more effectively than the time and energy consuming hydrogen reduction method. With only 0.64 vol% of graphene content, reduced graphene oxide (RGO)/ STO composite shows highly increased electrical conductivity and power factor compared to hydrogen reduced STO. Electron energy loss spectrum confirms the high concentration of oxygen vacancy in the area close to RGO in RGO/STO composite, which suggests a mild reaction between RGO and STO during sintering. In addition, the thermal conductivity is also depressed due to greatly restricted grain growth via adding RGO. Therefore, this study provides a fast, green and effective way of preparing oxygen deficient perovskite towards improved thermoelectric properties. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction With the ability of converting heat into electricity, thermoelectric materials have been intensively studied in the past decade due to the growing demand for energy and lack of solution for sustainable powder generation. Theoretically the efficiency of fuels can be largely improved when the waste heat is recycled via the application of thermoelectric materials with large dimensionless figure of merit, ZT ¼ sS2T/k, in which s is the electrical conductivity, S is the Seebeck coefficient, k is the thermal conductivity and T is the absolute temperature. Although numbers of high ZT materials such as Te based and Se based skutterudite and intermetallic compounds have been developed so far, most of these materials are suffered from toxicity, rarity in earth and instability [1,2]. In the regard of this situation, many researchers have revisited the possibility of using oxides as thermoelectric materials, despite the fact
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Fan), [email protected]. jp (A. Kawasaki). http://dx.doi.org/10.1016/j.carbon.2016.11.012 0008-6223/© 2016 Elsevier Ltd. All rights reserved.
that oxides are of low carrier mobility and high thermal conductivity in general. The representative p-type materials are NaCo2O4 and Ca3Co4O9 which have been found very promising especially in high temperature range, benefiting from their layered structure [3e6]. For n-type thermoelectric materials, SrTiO3 (STO) has attracted tremendous interests for its large carrier effective mass, stability at high temperature and structural tolerance with respect to doping. As a typical perovskite with composition of ABO3, STO is an intrinsic insulator without doping. Electrical conductivity can be rendered either by introducing foreign dopants to A-or B-site, or by reduction to generate oxygen vacancies. The former route is much more popular because of its obvious flexibility for tailoring properties. STO ceramics with A-site partly substituted by rare earth elements like Y, La, Sm, Gd and Dy have comparatively investigated by Muta et al., whose work suggests that the influence of A-site doping on ZT value mainly comes from the changes in thermal conductivity rather than power factor [7]. Moreover, Ohta et al. have shown that in comparison to A-site doped Sr1-xLaxTiO3, B-site doped SrTi1xNbxO3 single crystal exhibit larger power factor due to even larger density of states effective mass (m*), indicating the potential
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of Nb doping [8]. In their later work, polycrystalline ceramic of SrTi1-xNbxO3 was also fabricated. Albeit same level of power factor maintained compared to the corresponding single crystal film, thermal conductivity was not reduced effectively by grain boundaries, leading to a ZT value of 0.35 at 1000 K which is probably the highest level can be reached presently for STO ceramics at high temperature [9]. In contrast to doping with foreign atoms, oxygen deficient STO has not attracted sufficient attention so far even though this phenomenon occurs commonly when samples are fired in reductive atmosphere such as in vacuum or reducing gas. However, a few research groups have already noticed that oxygen vacancy can play an important role in optimizing the thermoelectric properties of doped STO. Yu et al. reported significantly reduced thermal conductivity in oxygen deficient La-doped STO film prepared by pulsed laser deposition (PLD) under a high vacuum of ~107 Torr at 923 K. The low thermal conductivity was mainly attributed to oxygen vacancies that suppressed heat transfer effectively [10]. Similar low thermal conductivities were obtained when PLD was conducted under a vacuum of 108 Torr at 1153 K for La-doped STO. More interestingly, it seems that the introduction of oxygen vacancies can improve the electrical conductivity by increasing carrier density without significant sacrifice of Seebeck coefficient when the A-site doping level is low (5%La), highlighting the great potential of oxygen deficient STO [11]. Apart from epitaxial films, bulk ceramics of STO with oxygen deficiency were reported as well, usually prepared by annealing in reducing gas (i.e. H2/Ar forming gas) at elevated temperature after sintering [12e15]. Nevertheless, gradient oxygen vacancy distribution in the sample cannot be avoided generally owing to the diffusion controlled process from surface to inside. Therefore more effective approach to prepare oxygen deficient STO bulk ceramics is of great importance to the enhancement of thermoelectric properties and application of STO based materials, since polycrystalline ceramics have unarguable advantage in cost in comparison to single crystals and chemically or physically deposited films. In this work, we demonstrate a novel and facile method for creating oxygen vacancy in perovskite ceramics using undoped STO as a model material. Instead of post-sintering heat treatment in reducing atmosphere, oxygen deficiency in STO was generated during spark plasma sintering (SPS) process by addition of uniformly dispersed graphene oxide (GO) in STO powders. The reduced GO (RGO) at elevated temperature served as carbon source that further reduced the STO at sintering temperature of 1473 K, leading to highly increased concentration of oxygen vacancy and residual graphene platelets as inclusion in the final STO compact. The attained RGO/oxygen deficient STO composite shows largely improved power factor and suppressed thermal conductivity compared with oxygen deficient STO prepared by long-time reduction in hydrogen, giving rise to a ZT value of 0.09 at 760 K in the absence of A- or B-site doping. 2. Experimental 2.1. Surface modification of STO The original Strontium Titanate powders (Sigma-Aldrich Co. LLC. Japan, purity of 99%, average size of 100 nm) were surface modified by (3-Aminopropyl) trimethoxy-silane (APTS, SigmaAldrich Co. LLC. Japan, purity of 97%). Firstly, STO powders were added into toluene (Wako Pure Chemical Industries, Ltd, Japan, purity of 99.99%) and dropped 4 ml APTS into the solution of STO. Then the mixture was heated to 423 K and holding for 6 h in the atmosphere of Ar with the reflux devices. The products were washed by toluene and filtered by vacuum filtration followed by
drying at 333 K. Here modified STO powders were obtained. 2.2. Preparation of RGO/STO composite Graphene oxide (GO) was prepared by the modified Hummers method [16]. STO (2.4 g) powders were poured into 500 ml deionized water and the pH was adjusted to 3.3. Through the ultrasonic treatment for 2 h, the colloid of STO formed. The colloid of GO was dropped to the solution of STO by the method of titration with stirring. The mixture was filtered and dried at 333 K. GO/STO hybrid powders were densified by Spark Plasma Sintering (SPS) apparatus (Dr. Sinter 511-S, Sumitomo Coal Mining Co. Ltd. Japan) with the graphite dies under the pressure of 60 MPa in the vacuum atmosphere. The sintering temperature was controlled from room temperature to 1473 K at the heating rate of 100 K/min, and holding for 5 min at 1473 K. During sintering, GO was reduced to RGO. Finally, the disk of RGO/STO composite with a diameter of 10 mm and a thickness of 3 mm was obtained. Al2O3/0.6 vol% RGO composite, as well as RGO/STO composite (0, 0.64 vol%) were prepared by the same method. For comparison, the pure STO ceramic was also sintered using the same condition. Oxygen deficient STO prepared by traditional method was prepared by heat treatment of pure STO ceramic in H2/Ar flowing gas at 1573 K for 10 h. 2.3. Characterization Zeta potentials of as-prepared GO colloid and STO colloid were measured using Zetasizer Nano ZS (Malvern, UK). This instrument utilizes Laser Doppler Micro-electrophoresis technique to measure zeta potential with a red laser (633 nm). The pH values of both samples were adjusted by MPT-2Autotitrator (Malvern, UK) using HCl and NaOH solutions (0.01 M). Before the measurement, both samples were sonicated for 0.5 h to achieve homogeneous dispersion. The carbon content in the composite was measured by combustion infrared detection technique (844 series, LECO, MI, US). The transmission electron microscopy (TEM) observation, electron energy-loss spectroscopy (EELS) and low-angle annular dark field (LAADF) were performed using a JEOL JEM-ARM200F electron microscope. The electron back scattered diffraction (EBSD) was performed using JEOL JSM-7100F. The field-emission scanning electron microscopy (SEM) observation was performed by a JEOL JSM-6500F scanning electron microscope. Micro-Raman spectra were measured with a SOLAR TII Nanofinder (Tokyo Instruments Co.) with 532 nm wavelength incident laser light and a 100 objective. The X-ray diffraction (XRD) patterns were collected by Rigaku Smartlab 9kw diffractometer with Cu Ka radiation source (l ¼ 0.15418 nm). The electrical conductivity was measured by a static dc method based on the slope of the voltage versus the temperature-difference curves using ZEM-3 (Ulvac-Riko) under a low-pressure helium atmosphere. The carrier density and the mobility were measured by a Hall measurement system (Toyo ResiTest8400 series) at room temperature in the air. The thermal diffusivity was measured by a laser flash method (LFA 457, Netzsch), and the specific heat was measured using a differential scanning calorimeter (DSC 404F3, Netzsch). 3. Results and discussion 3.1. Preparation of RGO/STO composites To set a model for the effect of graphene on thermoelectric properties of STO composites, the strategy adopted here for preparing the composite should meet several requirements: i) the quality of composite prepared by this strategy should be high,
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which means the uniform dispersion and absence of agglomeration of graphene in matrix are expected. ii) Though undoped STO was selected as a model material to study here, the strategy should be universal for doped STO and other types of perovskite so that the conclusions drawn here can be easily transfer to those materials as guidance for designing novel composites with improved thermoelectric properties. Based on these considerations, GO (precursor of graphene) and STO powders were homogeneously mixed by hetero-aggregation before densification. This strategy, which is especially effective when oxide particles can form large enough positive surface charge (zeta potential) in water to attract the negatively charged GO, has been successfully applied in many graphene/oxide composites in our previous researches [17e19]. Unfortunately, STO powders cannot form such stable colloid naturally because of the extremely low pzc (point of zero charge) of STO particles. To render positive surface charge to STO particles, NH2þ was grafted onto the surface of STO particle by reacting with (3Aminopropyl) trimethoxy-silane. Upon the surface modification, STO powders became positively charged in water and a maximum zeta potential of 45 mV can be obtained at pH of 3.3, which guarantees the formation of stable colloid (Supplementary Material Fig. S1). After mixing STO and GO colloids carefully, GO/STO hybrid powders with uniform distribution of GO were attained via a classic process of hetero-aggregation, in which each GO nanosheet is fully decorated by STO nano-particles, effectively hindering the recombination and agglomeration of GO nanosheets (Supplementary Material Fig. S3). The GO/STO hybrid powder was densified by SPS to attain fully dense RGO/STO composite, as shown in Table 1. For comparison, an oxygen deficient STO ceramic sample (denoted as STO-H) was also prepared by heat treatment of SPS sintered pure STO ceramic in H2/ Ar forming gas, following the traditional process of creating oxygen vacancies. It is worth noting that only the composite containing very low graphene content was prepared in this study, because RGO has been proved to be of very low thermopower that may deteriorate the thermoelectric performance of STO composite with increasing amount [19,20]. Thus it is better to depress the content of graphene in composite as long as uniform distribution of graphene in the composite can be achieved, since the negative effect could be overweighed by positive one such as highly increased electrical conductivity and maintained Seebeck coefficient arising from graphene introduced oxygen vacancy when the fraction of graphene in composite is low. Three different STO samples, including RGO/STO composite, pure STO ceramic and STO-H ceramic, were characterized by X-ray diffraction (XRD) in the first place to confirm the obtained phase. As shown in Fig. 1, the XRD patterns of all samples are in good consistence with cubic perovskite STO (PDF-#35-0734), indicating no phase transition occurred during sintering or reduction process. The absent peak of RGO in composite can be understood by its low content which is determined as 0.64 vol% (see methods) in this particular sample, while their existence as well as recovered structure from GO were unambiguously identified by Raman spectra (Supplementary Material Fig. S4). The refinement of XRD patterns shows that all samples share not only same cubic structure with space group of Pm3m(Table 1), but also very similar lattice
171
Fig. 1. XRD patterns of pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
parameter that is close to the value reported in standard card, which reveals no significant change in structure with respect to stoichiometric STO in these materials. In other words, despite the probable presence of oxygen vacancies in these STO ceramics owing to the more or less reducing process they are subjected to, the concentration of the defect is modest, in comparison to the highly oxygen deficient STO that adopts a lower symmetrically tetragonal structure [21]. It has been known that one of the most prominent effects of RGO on the microstructure of ceramic composite is the highly restrained grain growth [18], which has critical meaning for reducing thermal conductivity in thermoelectric materials. As demonstrated by electron back scattering diffraction (EBSD), without RGO the average grain size for pure STO is 13.31 mm with grain size distribution mainly ranging from 10 to 18 mm (Fig. 2a). For STO-H, the average grain size is even larger compared with pure STO, reaching 14 mm with a broader grain size distribution, due to the long-time annealing process after sintering (Fig. 2b). On the contrary, the average grain size of the composite is merely 0.58 mm for RGO/STO composite (Fig. 2c), which is more than 20 times smaller than that of pure STO and STO-H, despite the fact that the RGO content is only 0.64 vol% in composite. More interestingly, although the grains in composite are mostly concentrated on 0.6 mm, there is indeed a mixing of grain size in the composite including gains smaller than 0.2 mm and larger than 1 mm (Fig. 2c and d), which could be beneficial in improving ZT value of composite according to the theory proposed by Zhao et al. [22]. Since the detailed structural information for RGO in composite can hardly be given by EBSD or scanning electron microscopy (Supplementary Material Fig. S5), the microstructure of RGO/STO composite was observed by TEM and HRTEM, as shown in Fig. 3. The fully dense and fine grained structure in composite was reconfirmed by TEM observation (Fig. 3a). Under higher magnification, RGO platelet can be clearly identified owning to its lower
Table 1 Crystal structure, composition and relative density for pure STO, STO-H and RGO/STO composite. RGO content [vol%]
Apparent density [gcm3]
Relative density [%]
Crystal structure
Space group
Lattice parameter [Å]
Pure STO
0
5.07
99.02
Cubic
Pm3m
3.90906 (4)
STO-H
0
5.03
98.24
Cubic
Pm3m
3.90800 (4)
RGO/STO
0.64
5.05
99.02
Cubic
Pm3m
3.91017 (5)
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Fig. 2. EBSD analysis of (a) pure STO, (b) pure STO-H and (c) RGO/STO composite; (d) the grain size distributions for pure STO, pure STO-H and RGO/STO composite, respectively. (A colour version of this figure can be viewed online.)
3.2. Electrical properties of RGO/STO composites Stoichiometric STO is an insulator owing to the large band gap
(3.2eV) [23]. However, at high temperature of above 1273 K, neutral oxygen on regular oxygen site (written asOxo ) of STO is thermodynamically prone to leave the lattice, retaining behind an oxygen vacancy with two electrons [24,25].
Oxo 4Vo€ þ 2e þ 1 2O2 =
contrast compared to surrounding STO matrix in bright field image (Fig. 3b), which allows us to verify the uniform distribution of RGO platelets and the absence of severe agglomerations in the composite as well. Moreover, two types of RGO platelets can be found in the composite: the one existing at grain boundaries and the one included inside the grains (Fig. 3d). The former type is very common that can be readily found in any RGO/oxide composite [17,18]. This type of RGO can effectively pin the movement of grain boundary, which is the main reason for the highly restrained grain growth. However, the latter type that RGO completely enveloped in a grain is very rare and unique, which has never been observed in other RGO/oxide composites to the best of our knowledge, since normally the lateral size of RGO is much larger than STO particles (Supplementary Material Fig. S2) that makes grain boundary completely moving across a whole piece of platelet very difficult [18]. Noting the enveloped RGO platelets all have very small size which looks more like the debris of RGO (Fig. 3c), it is very likely that these small pieces of RGO inside grains are actually a result of the mild reaction between STO and RGO during sintering process. As the other consequence of this reaction, STO lost oxygen to become oxygen deficient STO ceramic, which will be demonstrated in the next section.
(1)
Reducing environment like vacuum can further assist the escape of oxygen, leading to a slightly conductive STO ceramic after SPS (pure STO), although the conductivity of which is generally low and unstable because of very limited and inhomogeneous oxygen vacancies generated in this process. As shown in Fig. 4, after repeating measurement to achieve a steady performance (Supplementary Material Fig. S6), the as-sintered pure STO shows a very low electrical conductivity of 17.9 Sm1 at room temperature, which increases slightly with increasing temperature, reaching 176 Sm1 at 760 K finally. To promote more oxygen vacancies for tuning thermoelectric properties of perovskite materials, researchers have to employ stronger reducing conditions usually involving long time annealing in hydrogen gas [13,14], though the process is both time and energy consuming. Following this traditional strategy, the electrical conductivity of STO-H shows much improved electrical conductivity of 640 Sm1 at room temperature and 2025 Sm1 at 570 K (maximum value), revealing considerably increased oxygen vacancies compared with pure STO. Moreover, it is very surprised that without any additional reducing processes, the electrical conductivity of as-sintered RGO/STO composite is even higher than
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173
Fig. 3. TEM and HRTEM images of RGO/STO composite. (a) and (b) are TEM images at low and high magnification, respectively. (c) and (d) are HRTEM images from the area indicated by black square and black dash line square respectively, showing RGO on the grain boundary (c) and inside the grain (d).
that of STO-H, manifesting a conductivity of 1633 Sm1 at room temperature, and reaching the maximum 4669 Sm1 at 470 K. Theoretically, the improvement of electrical conductivity may arise from the increase in carrier mobility (m) or carrier density (n), or both of them because the conductivity can be written as:
s ¼ mne
(2)
Therefore, to clarify the reason behind the largely enhanced electrical conductivity in RGO/STO composite, it is necessary to figure out which component leads to the increase most. Using the hall measurement, the carrier density and correspondingmobility at room temperature were determined for pure STO, STO-H and composite (Table 2). It is found that in contrast to the pure STO, the STO-H has both increased carrier density and mobility, owing to the increased oxygen vacancies induced by reduction. More strikingly, the carrier density in RGO/STO composite (1.59 1020cm3) is about 3.5 times higher than STO-H (4.54 1019cm3), while the mobility is only slightly larger compared to STO-H. Therefore, it can be concluded that the improvement of the electrical conductivity in RGO/STO composite mainly comes from the increased carrier density rather than mobility, which is totally different from the situation in graphene/doped ZnO composite reported recently [26].
However, it is still not clear what leads to the highly increased carrier density, albeit apparently the reason should be related to RGO, since the only difference between pure STO and composite is the addition of RGO. At first glance, one may ascribe the enhanced electrical conductivity to the high carrier concentration of RGO. Indeed, graphene is famous for high electrical conductivity [27], but it is questionable whether RGO has such capability that can boost the carrier density of composite by more than 40 times with a loading of only 0.64 vol%. In order to clarify this problem, a RGO/ Al2O3 composite with similar graphene loading (0.6 vol%) was fabricated using the identical method as a special reference. Unlike STO, Al2O3 is a very stable insulator even under very harsh circumstances, which makes RGO the only conductive component in RGO/Al2O3 composite. At room temperature, the Al2O3/RGO composite shows an electrical conductivity of only 84.2 Sm1 (Fig. S7) and corresponding carrier density of 3.92 1018cm3 (Table S1) which is two orders lower than the carrier density in RGO/STO composite. Even if we assume only 10% RGO formed effective conducting path in RGO/Al2O3 composite, the carrier density that can be induced by ~0.6 vol% of RGO could not be higher than 3.92 1019cm3. Therefore, rather than the high electrical conductivity of RGO, there must be another reason accounting for the
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to grain boundary. In contrast, the O-K edge on the grain boundary lost details at higher energy loss in the composite, which is consistent with the observation for oxygen deficient STO [28e30]. The Ti-L edge can provide more powerful evidence since the explicit distinction between Ti3þ and Ti4þ formal valences [28,31,32]. Because every oxygen vacancy can nominally donate two electrons to nearest Ti d band and the number of oxygen is three times more than titanium, the sensitivity of Ti is six times higher than O. It is found that in RGO/STO composite, the EELS spectra of Ti-L edge for grain boundary with RGO are distinctly different from that for inner grain without RGO (no carbon signal in the area), firmly proving the high concentration of Ti3þ ions close to RGO. Besides, the LAADF imaging can also provide evidence of RGO promoted oxygen vacancies (Fig. S8), since the LAADF signal can detect the strain fields from oxygen vacancies through their distortion of the adjacent cation sites and subsequent dechannelling of the electron beam from the cation columns [28]. Therefore, we can conclude that the amount of oxygen vacancy in composite that is mainly concentrated at the vicinity of RGO is much higher compared with pure STO. Considering the small pieces of RGO inside STO grains, the reason behind the highly increased carrier density for RGO/STO composite becomes much clear: RGO in composite served as carbon source that promoted the formation of oxygen vacancy in STO composite via mildly reacting with oxygen atoms on the surface of STO grain. During this process, some RGO platelets finally lost adequate size to hinder the movement of grain boundary and were enclosed by STO grains.
3.3. Thermal conductivity and ZT value The Seebeck coefficient of pure STO, pure STO-H and RGO/STO composite with the function of temperature is plotted in Fig. 4b. Compared with pure STO, the Seebeck coefficient of RGO/STO composite is just decreased by 16.6%, despite the two orders higher electrical conductivity. For degenerate semiconductors, two significant parameters that are responsible to the Seebeck coefficient are carrier density n and effective mass m* at determined temperature, which can be expressed as the following equation [33]:
. S ¼ 8p2 k2B 3eh2 m Tðp=3nÞ2=3
Fig. 4. (a) Electrical conductivity, (b) Seebeck coefficient, and (c) power factor of pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
high carrier density in RGO/STO composite, which is very likely to be the oxygen vacancy in RGO/STO composite. To prove our speculation, Electrical Energy-loss Spectroscopy (EELS) was employed to investigate the electronic structure change in RGO/STO composite and pure STO. As shown in Fig. 5, the O-K edge fine structure in pure STO is completely same from inner grain
(3)
where kB is Boltzmann constant; h is Planck constant; e is the carrier charge; and T is the absolute value. Therefore, the increase of carrier density usually results in the decrease of Seebeck coefficient if the effective mass is unchanged. However, with the introduction of oxygen vacancies, the effective mass increased for STO (Table 2), thus largely retaining the Seebeck coefficient. As a result, the power factor of RGO/STO composite is one order higher than that of pure STO and almost 2 times higher than that of STO-H at 520 K, shown in Fig. 4c. This large value of power is even comparable with dopedSTO such as Pr-doped STO (8.03 104Wm1 K2 at 1170 K, estimated data from the figures in the paper) [34]. The thermal diffusivity (D), specific heat (Cp) and density (d) were measured to obtain the total thermal conductivity, which is showed in Fig. 6a. The electronic thermal conductivity (ke) was
Table 2 Carrier density, mobility and effective mass for pure STO, STO-H and RGO/STO composite at room temperature.
Pure STO STO-H RGO/STO a
m0 is the mass of free electron.
Carrier density (n) [cm3]
Carrier mobility (m) [cm2V1s1]
Effective mass (m*)a
3.64 1018 4.54 1019 1.59 1020
0.29 0.49 0.64
0.41m0 2.41m0 4.39m0
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Fig. 5. EELS spectra for pure STO and RGO/STO composite. (a) (c) Ti-L2,3 edge, the Ti-L edge shifts from a þ4 to þ3 valence with increasing oxygen vacancies, as the oxygen vacancies are electron donors. (b) (d) O-K-edge, the O-K edge on the grain boundary lost details at higher energy loss spectrum with the introduction of oxygen vacancies. (A colour version of this figure can be viewed online.)
separated from lattice thermal Wiedemann-Franz law [35]:
ke ¼ sLT
conductivity
(kl)
by
the
(4)
where L is the Lorenz number (2.44 108V2K2). It can be seen that the contribution from electronic thermal conductivity is almost negligible in pure STO (Table S2), STO-H and RGO/STO composite, while the lattice thermal conductivity plays an important role. kl for pure STO is 7.7 Wm1 K1 at 370 K and decreased with the elevated temperature, and reached 5.6 Wm1 K1 at 760 K. For STO-H, kl is even higher than pure STO, revealing the influence from larger grain size induced by long-time heat treatment in hydrogen. In contrast, the lattice thermal conductivity of RGO/STO is just 5.7 Wm1 K1 at 370 K and 3.7 Wm1 K1 at 760 K, which decreased by 26% compared with pure STO. This depressed lattice
thermal conductivity in comparison to pure STO can be ascribed to the highly restrained grain size which is 20 times smaller than that of pure STO, thanks to the filler phase of RGO. The temperature dependence of the dimensionless figure of merit (ZT) for pure STO, STO-H and RGO/STO composite is plotted in Fig. 6b. It is found the ZT value of RGO/STO composite is much higher than that of pure STO and pure STO-H, exhibiting a ZT value of 0.03 at room temperature and 0.09 at 760 K, which is 2-fold higher than STO-H at 760 K. This improvement in ZT mainly comes from three factors: i) incorporation of RGO into STO can highly enhance the electrical conductivity of the composite via generating oxygen vacancies, rather than via the high electrical conductivity of itself; ii) the increased electrical conductivity does not deteriorate the Seebeck coefficient due to the increased effective mass of electrons simultaneously; iii) the thermal conductivity
Fig. 6. Temperature dependence of (a) the total thermal conductivity, (b) the lattice thermal conductivity and (c) the ZT value for pure STO, pure STO-H and RGO/STO composite. (A colour version of this figure can be viewed online.)
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of STO composite is also decreased due to the highly suppressed grain growth induced by RGO. 4. Conclusions In summary, we demonstrate a novel, fast and effective method for preparing oxygen deficient perovskite by incorporation with graphene using STO as the model material. Through surface modification of STO and a hetero-aggregation method to obtain uniform mixing between STO powder and GO before densification, a fully dense RGO/STO composite with homogeneous dispersion of 0.64 vol% RGO in matrix has been achieved in this study. It is found that with this low filler content, the composite displays highly enhanced electrical conductivity and much depressed thermal conductivity in comparison to hydrogen reduced STO, leading to a ZT value of 0.09 at 760 K without any foreign dopants. Further investigation reveals that the highly increased electrical conductivity mainly arises from the oxygen vacancy induced by mildly reaction between RGO and STO, which is much more effective than the traditional method via long-time hydrogen reduction. Meanwhile, the thermal conductivity is decreased owning to the highly restrained average grain size induced by RGO. Therefore, incorporating with graphene could be a very effective, convenient and novel method for modulating the thermoelectric properties of perovskite ceramics. Acknowledgements The authors would like to thank the support of China Scholarship Council. This work was funded by Natural Science Foundation of China (No. 51374078，51403037, 51272042), the Opening Project of State Key Laboratory of High Performance Ceramics and Superfine Microstructure(SKL201507SIC), the Fundamental Research Funds for the Central Universities and DHU Distinguished Young Professor Program. We appreciate the generous helps from Dr. WeiWei Zhou and Xiaohao Sun in Kawasaki Lab. Many thanks are also given to Qinghuan Huo, Dr. Kosei Kobayashi and Dr. Takamichi Miyazaki in Tohoku University. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.carbon.2016.11.012. References [1] K. Koumoto, R. Funahashi, E. Guilmeau, Y. Miyazaki, A. Weidenkaff, Y.F. Wang, et al., Thermoelectric ceramics for energy harvesting, J. Am. Ceram. Soc. 96 (1) (2013) 1e23. [2] J. He, Y.F. Liu, R. Funahashi, Oxide thermoelectrics: the challenges, progress, and outlook, J. Mater. Res. 26 (15) (2011) 1762e1772. [3] I. Terasaki, Y. Sasago, K. Uchinokura, Large thermoelectric power in NaCo2O4 single crystals, Phys. Rev. B 56 (20) (1997) 12685e12687. [4] R. Funahashi, I. Matsubara, H. Ikuta, T. Takeuchi, U. Mizutani, S. Sodeoka, An oxide single crystal with high thermoelectric performance in air, Jpn. J. Appl. Phys. Part 2 Lett. 39 (11B) (2000). L1127eL9. [5] A.C. Masset, C. Michel, A. Maignan, M. Hervieu, O. Toulemonde, F. Studer, et al., Misfit-layered cobaltite with an anisotropic giant magnetoresistance: Ca3Co4O9, Phys. Rev. B 62 (1) (2000) 166e175. [6] Y. Miyazaki, K. Kudo, M. Akoshima, Y. Ono, Y. Koike, T. Kajitani, Low-temperature thermoelectric properties of the composite crystal [Ca2CoO3.34]0.614[CoO2], Jpn. J. Appl. Phys. Part 2 Lett. 39 (6A) (2000). L531eL3. [7] H. Muta, K. Kurosaki, S. Yamanaka, Thermoelectric properties of rare earth doped SrTiO3, J. Alloy Compd. 350 (1e2) (2003) 292e295. [8] S. Ohta, T. Nomura, H. Ohta, K. Koumoto, High-temperature carrier transport and thermoelectric properties of heavily La- or Nb-doped SrTiO3 single
crystals, J. Appl. Phys. 97 (3) (2005) 4. [9] S. Ohta, H. Ohta, K. Koumoto, Grain size dependence of thermoelectric performance of Nb-doped SrTiO3 polycrystals, J. Ceram. Soc. Jpn. 114 (1325) (2006) 102e105. [10] C.H. Yu, M.L. Scullin, M. Huijben, R. Ramesh, A. Majumdar, The influence of oxygen deficiency on the thermoelectric properties of strontium titanates, Appl. Phys. Lett. 92 (9) (2008) 3. [11] J. Ravichandran, W. Siemons, D.W. Oh, J.T. Kardel, A. Chari, H. Heijmerikx, et al., High-temperature thermoelectric response of double-doped SrTiO3 epitaxial films, Phys. Rev. B 82 (16) (2010) 5. [12] J. Liu, C.L. Wang, W.B. Su, H.C. Wang, P. Zheng, J.C. Li, et al., Enhancement of thermoelectric efficiency in oxygen-deficient Sr1-xLaxTiO3-d ceramics, Appl. Phys. Lett. 95 (16) (2009) 3. [13] A.V. Kovalevsky, A.A. Yaremchenko, S. Populoh, A. Weidenkaff, J.R. Frade, Effect of a-site cation deficiency on the thermoelectric performance of donorsubstituted strontium titanate, J. Phys. Chem. C 118 (9) (2014) 4596e4606. [14] L. Sagarna, S. Populoh, A. Shkabko, J. Eilertsen, A.E. Maegli, R. Hauert, et al., Influence of the oxygen content on the electronic transport properties of SrxEu1-xTiO3-d, J. Phys. Chem. C 118 (15) (2014) 7821e7831. [15] A.A. Yaremchenko, S. Populoh, S.G. Patricio, J. Macias, P. Thiel, D.P. Fagg, et al., Boosting thermoelectric performance by controlled defect chemistry engineering in Ta-Substituted strontium titanate, Chem. Mat. 27 (14) (2015) 4995e5006. [16] S. Park, J.H. An, R.D. Piner, I. Jung, D.X. Yang, A. Velamakanni, et al., Aqueous suspension and characterization of chemically modified graphene sheets, Chem. Mat. 20 (21) (2008) 6592e6594. [17] Y.C. Fan, W. Jiang, A. Kawasaki, Highly conductive few-layer Graphene/Al2O3 nanocomposites with tunable charge carrier type, Adv. Funct. Mater. 22 (18) (2012) 3882e3889. [18] Y.C. Fan, M. Estili, G. Igarashi, W. Jiang, A. Kawasaki, The effect of homogeneously dispersed few-layer graphene on microstructure and mechanical properties of Al2O3 nanocomposites, J. Eur. Ceram. Soc. 34 (2) (2014) 443e451. [19] Y.C. Fan, L.J. Kang, W.W. Zhou, W. Jiang, L.J. Wang, A. Kawasaki, Control of doping by matrix in few-layer graphene/metal oxide composites with highly enhanced electrical conductivity, Carbon 81 (2015) 83e90. [20] N.D.K. Tu, J. Choi, C.R. Park, H. Kim, Remarkable conversion between n- and ptype reduced graphene oxide on varying the thermal annealing temperature, Chem. Mat. 27 (21) (2015) 7362e7369. [21] W.H. Gong, H. Yun, Y.B. Ning, J.E. Greedan, W.R. Datars, C.V. Stager, Oxygendeficient SRTIO3-X, X ¼ 0.28, 0.17, AND 0.08-crystal-growth, crystal-structure, magnetic, and transport-properties, J. Solid State Chem. 90 (2) (1991) 320e330. [22] L.D. Zhao, B.P. Zhang, W.S. Liu, J.F. Li, Effect of mixed grain sizes on thermoelectric performance of Bi2Te3 compound, J. Appl. Phys. 105 (2) (2009) 6. [23] A. Ganguly, S. Sharma, P. Papakonstantinou, J. Hamilton, Probing the thermal deoxygenation of graphene oxide using high-resolution in situ x-ray-based spectroscopies, J. Phys. Chem. C 115 (34) (2011) 17009e17019. [24] R. Moos, K.H. Hardtl, Defect chemistry of donor-doped and undoped strontium titanate ceramics between 1000 degrees and 1400 degrees C, J. Am. Ceram. Soc. 80 (10) (1997) 2549e2562. [25] N.G. Eror, U. Balachandran, Electrical-conductivity in strontium-titanate with nonideal cationic ratio, J. Solid State Chem. 42 (3) (1982) 227e241. [26] W.H. Nam, B.B. Kim, S.G. Seo, Y.S. Lim, J.Y. Kim, W.S. Seo, et al., Structurally nanocrystalline-electrically single crystalline ZnO-reduced graphene oxide composites, Nano Lett. 14 (9) (2014) 5104e5109. [27] A.K. Geim, K.S. Novoselov, The rise of graphene, Nat. Mater. 6 (3) (2007) 183e191. [28] D.A. Muller, N. Nakagawa, A. Ohtomo, J.L. Grazul, H.Y. Hwang, Atomic-scale imaging of nanoengineered oxygen vacancy profiles in SrTiO3, Nature 430 (7000) (2004) 657e661. [29] S. Hillyard, J. Silcox, Detector geometry, thermal diffuse-scattering and strain effects in ADF stem imaging, Ultramicroscopy 58 (1) (1995) 6e17. [30] D.D. Perovic, C.J. Rossouw, A. Howie, Imaging elastic strains in high-angle annular dark-field scanning-transmission electron-microscopy, Ultramicroscopy 52 (3e4) (1993) 353e359. [31] M. Abbate, F.M.F. Degroot, J.C. Fuggle, A. Fujimori, Y. Tokura, Y. Fujishima, et al., Soft-X-ray-absorption studies of the location of extra charges induced by substitution in controlled-valence materials, Phys. Rev. B 44 (11) (1991) 5419e5422. [32] A. Ohtomo, D.A. Muller, J.L. Grazul, H.Y. Hwang, Epitaxial growth and electronic structure of LaTiOx films, Appl. Phys. Lett. 80 (21) (2002) 3922e3924. [33] P.A. Lee, T.V. Ramakrishnan, Disordered electronic systems, Rev. Mod. Phys. 57 (2) (1985) 287e337. [34] A.V. Kovalevsky, A.A. Yaremchenko, S. Populoh, A. Weidenkaff, J.R. Frade, Enhancement of thermoelectric performance in strontium titanate by praseodymium substitution, J. Appl. Phys. 113 (5) (2013) 8. [35] G.V. Chester, A. Thellung, Law of wiedemann and franz, Proc. Phys. Soc. Lond. 77 (497) (1961) 1005e1013.