Comprehensive theoretical and experimental study of electrical transport mechanism on BiFeO3 multiferroic nanoparticles

Journal of Alloys and Compounds 720 (2017) 47e53

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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Comprehensive theoretical and experimental study of electrical transport mechanism on BiFeO3 multiferroic nanoparticles ~o, J.A. Souza* F.E.N. Ramirez, E. Marinho Jr., C.R. Lea Universidade Federal do ABC, CEP 09210-170, Santo Andr e, SP, Brazil

a r t i c l e i n f o

a b s t r a c t

Article history: Received 20 January 2017 Received in revised form 12 May 2017 Accepted 19 May 2017 Available online 22 May 2017

We have performed a comprehensive study of electrical properties at high temperatures of single-phase BiFeO3 nanoparticles. Our results revealed a strong dependence of electrical resistivity behavior and activation energy on thermal cycle due to adsorption/desorption of water molecules. Ionic conductivity due to oxygen vacancies as well as Hþ ions should dominate the electronic current after the initial heating and cooling cycle explaining the change in the charge transport properties due to thermal, temperature, and atmosphere history. Electronic structure simulations indicated that the characteristics of the conductivity in the original sample is consistent with band-to-band transitions, expected in a material rich in oxygen. © 2017 Elsevier B.V. All rights reserved.

Keywords: Multiferroic nanoparticles Electrical properties Humidity adsorption Charge carrier delocalization Oxygen vacancy Formation energy

1. Introduction Multiferroics materials, which show coexistence of ferromagnetic (FM) and ferroelectric ordering, are of great interest because of many technological applications, such as information storage, spintronic, gas sensors and solar energy conversion [1e6]. Very high electrical resistivity is pursued in order to improve the multiferroic properties and, consequently, their possible technological applications. Perovskite BiFeO3 is a very interesting system because the coexistence of both ferroelectricity and antiferromagnetic orders over a wide temperature range above room temperature [7,8]. This compound is a rhombohedrally distorted perovskite system belonging to the space group R3c. The ferroelectric Curie temperature (TC) and Neel Temperature (TN) of bulk BiFeO3 compound are TC ~ 1100 K and TN ¼ 643 K, respectively [9e11]. A canted spin structure gives a spiral modulation with a periodicity of 62 nm, incommensurate with the crystal lattice [12,13]. The electrical properties of BiFeO3 nanoparticles are intricate and can occur in different regimes. It is known that stoichiometric single-phase samples are hard to obtain because the presence of oxygen vacancies and ions of Fe with mixed valence states (Fe2þ

* Corresponding author. E-mail address: [email protected] (J.A. Souza). http://dx.doi.org/10.1016/j.jallcom.2017.05.191 0925-8388/© 2017 Elsevier B.V. All rights reserved.

and Fe3þ) in the crystalline structure of BiFeO3 [14e20]. Besides the intrinsic semiconductor mechanism, the presence of these defects may lead to the appearance of other types of transport process depending on the temperature interval. The low-temperature dielectric relaxation is considered as a typical characteristic of electronic hopping between Fe2þ and Fe3þ inside the grain whereas the middle-temperature process arises from electrically heterogeneous regions such as grain boundary. At high temperature, the process is related to mobility of oxygen vacancies. Each of these mechanisms has different activation energy and is predominant in each temperature range [21,22]. BiFeO3 compound therefore exhibits gas sensor properties, with electrical resistivity being very sensitive to the atmosphere [23,24]. It is known that this property is mainly caused by physical or chemical adsorption of several atomic species nearby the sample surface [25e27]. In this context, the temperature dependence of the electrical resistivity is expected to be strongly influenced by humidity. However, works about the electric transport of BiFeO3 at high temperature are scarce in the literature. Indeed, the effect of humidity adsorption with thermal treatment on electrical resistivity is not discussed in literature. In this work, we have performed a comprehensive study of electrical properties at high temperatures of single-phase BiFeO3 nanoparticles with crystallite average size of ~60 nm. The influence of thermal cycles and humidity adsorption on electrical resistivity measurements along

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with computational simulations were studied. We have performed quantum simulations to obtain a better understanding on the behavior of point defects including hydrogen impurities in the electronic properties of BiFeO3. Even though there is a lot of study in this system, measurements of dc electrical resistivity as a function of temperature taken into account the humidity contribution are not found in literature. Our results revealed strong dependence of electrical conductivity behavior and activation energy on humidity absorption and thermal history.

2. Experimental and computational details BiFeO3 nanoparticle sample has been obtained by a sol-gel modified method using metal nitrates Bi(NO3)3$5H2O and Fe(NO3)3$9H2O as starting materials. The precursors were obtained by dissolving stoichiometric amounts of the metal nitrates in diluted nitric acid (HNO3). Ethylenediaminetetraacetic acid (EDTA) was used as polymerizing agent. Details of the synthesis are given elsewhere [28]. The crystal structure was studied by X-ray powder diffraction (XRD) measurements on a q-2q Bruker AXS D8 Focus diffractometer with Cu Ka radiation. Structural parameters of asprepared BiFeO3 nanoparticles were refined by using Rietveld method. For all Rietveld refinements, we have assumed the hexagonal representation of the space group R3c being the starting values ahex ¼ 5.577 Å, chex ¼ 13.86 Å, Bi (0, 0, 0.2988), Fe (0, 0, 0.197), O (0.2380, 0.3506, 1/12) [29]. In this representation, the displacement values of the Bi and Fe from its non-polar equilibrium positions are 0.049 Å and 0.197 Å, respectively. The values of the quality indices obtained from the refinement were Rwp ¼ 3.54% and Rp ¼ 2.66%. The “goodness of fit” is defined by S ¼ Rwp/Rexp, where Rexp is the expected statistical value for Rwp (2.5%). Thus, we have obtained S ¼ 1.42% which indicates a high-quality refinement. The value found for the “Chi-200 was 3.46. We have used the Scherrer equation corrected for instrumental peak broadening, which was determined with Al2O3 standard, in order to calculate the crystallite sizes (dXRD). The morphology and the average particles size of nanoparticles were studied by scanning electron microscopy (SEM) images obtained from a JMS-6701F microscope. For electrical measurements, the powder was pressed into pellets at high pressure and sintered in inert atmosphere at 500  C for 30 min. The electrical transport mechanism was studied by DC electrical resistivity measurements using a fourprobe method. Impedance spectroscopy measurements were carried out by using a Solartron Impedance Analyzer SI1260. All these electrical characterizations have been made by using a custombuilt dedicated apparatus. In all measurements, silver epoxy was used to connect the sample with the station, which is composed of alumina, and the wiring is made of platinum. Density functional theory [30,31] calculations were performed using Vienna Ab Initio Simulation Package [32] (VASP), within the projected augmented wave [33] (PAW) method. The generalized gradient approximation parameterized by Perdew-BurkeErnzerhof [34] (GGA) was used for the exchange-correlation functional. The cutoff energy for plane-wave expansion was 500 eV. Structural optimizations were performed using a conjugate gradient algorithm until the Hellmann-Feynman forces on all atoms were less than 102 eV/Å. We adopted the GGA þ U method [35] for a better description of localized Fe 3d electrons, with an effective Hubbard parameter Ueff ¼ U e J equal to 4 eV [36]. We constructed a supercell with 80 host atoms, 2a, 2b and 2c of rhombohedral R3c unit cell. In all calculations, we turn on spin polarization and we do not include the spineorbit interaction. Point defects were created by adding and/or removing atoms in the supercell. For our investigation of defect formation energies, we

have analyzed the following intrinsic and extrinsic point defects: bismuth, iron and oxygen vacancies (Bivac, Fevac and Ovac, respectively), interstitial hydrogen (Hint) and hydrogen atom into oxygen vacancy (H þ Ovac). 3. Results and discussion 3.1. Structural and morphological characterizations Fig. 1(a) shows the X-ray powder diffraction pattern along with the Rietveld refinement obtained for BiFeO3 nanoparticles sample. We can see that the sample is single-phase belonging to rhombohedral space group R3c. The unit cell parameters found from Rietveld refinement are ahex ¼ 5.5679(1) Å, chex ¼ 13.819(1) Å, V ¼ 371.02(3) Å3. The size of crystallographic domains (crystallites) calculated from Scherrer formula is 60(5) nm. Fig. 1(b), (c) and (d) shows the scanning electron microscopy (SEM) images obtained with different magnifications. We have observed that the nanoparticles are uniformly distributed forming agglomerations with porous structure, typically observed in sol-gel synthesis. Another interesting observation is the spatial organization of the nanoparticles. A careful inspection of Fig. 1(c) and (d) reveals that the nanoparticles are joined forming ring-like structures. An estimation of size distribution has revealed particles in the range of 74 ± 15 nm. Even though Scherrer formula may not necessarily yield the correct results for nanoparticles lower than 100 nm, the estimated value agrees with SEM results. 3.2. Humidity influence on electrical transport mechanism Fig. 2(a) shows the electrical resistivity (r) as a function of temperature for three subsequent thermal cycles in air atmosphere. These measurements were obtained by warming the sample to the highest temperature and cooling down back to room temperature. It is important to mention that the sample was placed into vacuum desiccator before starting the measurements. This strategy allowed us to obtain information about how the initial humidity adsorption influences the electrical resistivity. For warming curve of cycle 1, we were unable to measure the electrical resistivity at room temperature due to its huge value (it may reach 1 GU cm). Commonly this is interpreted in the literature as an evidence that the sample is highly stoichiometric with low concentration of defects or impurities [37e39] and more importantly, that there is no humidity absorbed on the surface. A detailed analysis of temperature dependence of electrical resistivity curves revealed several interesting results. Due to its high value, the electrical resistivity measurements start to be recorded only when increasing the temperature to around T ¼ 420 K (see cycle 1). Above this temperature, a semiconducting behavior is observed as the temperature is increased further. As the system is cooled down slowly from T ~700 K, the semiconducting behavior is still preserved and below 350 K a huge decreasing in the electrical resistivity is observed e a metallic-like behavior. This same metallic behavior was observed for cycles 2 and 3. We suggest that this result is due to humidity adsorption on the surface of the sample present in air atmosphere. As will be seen later, the humidity adsorption can induce formation of oxygen vacancies and hydrogen related defects, which will affect the amount of free charge carriers in the material. By further decreasing temperature below 350 K, the humidity adsorption and, consequently, the amount of free charge carries will increase causing the huge decrease in the electrical resistivity. We have repeated this measurement in an inert atmosphere and the decreasing of r is not observed confirming that moisture plays an important role in the transport mechanism around room temperature for this system. We have also observed

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Fig. 1. (a) X-ray powder diffraction pattern along with Rietveld refinement for the as-prepared BiFeO3 sample. The tick marks represent the expected Bragg reflections for the rhombohedral phase. (b), (c), and (d) SEM images showing BiFeO3 nanoparticles with different magnifications.

anomalous changes (downturn) in the slope of electrical resistivity for the three cycles. For cycle 1, this anomalous variation was observed close to 488 K whereas for cycles 2 and 3 were observed at different temperatures e 400 K and 427 K (for cycle 2) and 428 K and 446 K (for cycle 3). The nature and origin of these anomalies, which were also observed by other authors, is not clear [40e42]. One can also see a distinct change in the electrical resistivity behavior around e T* ¼ 592 K which is related to the antiferromagnetic to paramagnetic phase transition of BiFeO3 compound. We have studied the transport mechanism in the semiconducting regime and the activation energy of the system has been calculated by using Arrhenius model. In this model, the electrical resistivity is described by the equation r ¼ r0exp(Ea/KBT), where r0 is a pre-factor and Ea is the activation energy. According to Arrhenius model, the curves Ln (r) vs T1 should exhibit a linear behavior. Fig. 2(b) shows the curves Ln (r) vs T1 corresponding to the warming curves obtained in air atmosphere for three cycles. It is observed that Arrhenius mechanism is obeyed in both temperature ranges T < TN and T > TN. Interestingly, the values of activation energy obtained from linear fitting are different for each cycle. For T < TN, we have found Ea ¼ 1.53 eV, 0.84 eV, and 0.82 eV for cycles 1, 2, and 3, respectively. This result indicates that the activation energy is dependent on thermal cycles. On the other hand, these values are very close to contribution of oxygen vacancies mobility. For T > TN, we have obtained Ea ¼ 0.72(1) eV (for cycle 1) and 0.13(1) eV (for cycles 2 and 3). An abrupt decrease of Ea is observed in the high temperature paramagnetic phase. This result agrees with the behavior reported by previous works, suggesting that the oxygen vacancies become more localized in the ordered magnetic phase [43]. The delocalization processes of charge carriers and its evolution

with temperature can also be studied by impedance spectroscopy measurements. Fig. 3(a) and (b) show the imaginary part (Z00 ) of impedance as a function of frequency and temperature obtained in air atmosphere for cycles 1 and 3, respectively. All curves were normalized to value of Z00 measured at T ¼ 300 K for a better visualization. In Fig. 3(a), the numbers indicate the temperature of each measurement: (1) T ¼ 300 K, (2) T ¼ 378 K, (3) T ¼ 423 K, (4) T ¼ 473 K, (5) T ¼ 523 K, and (6) T ¼ 618 K. At T ¼ 300 K, we have observed a monotonically decreasing of Z00 with increasing frequency at room temperature. At higher temperature, Z00 starts exhibiting a well-defined maximum as a function of frequency. This maximum starts to be observed at T ¼ 378 K. The frequency corresponding to the maximum is known as relaxation frequency (fr) due to relaxation processes of charge carriers or electric dipolar moments [44]. The relaxation time t of the system can be obtained by t ¼ 1/fr. By increasing the temperature, the position of this peak is found to shift towards higher frequency side. These results suggest that at room temperature, where it was not possible to detect any relaxation process, the charge carriers are strongly trapped in agreement with the huge value of the electrical resistivity observed for cycle 1 at this temperature (1 GU cm). By increasing the temperature, a decrease in the relaxation time was observed due to initial delocalization of charge carriers caused by thermal activation. On the other hand, the temperature dependence of Z00 exhibits a different behavior for cycle 3. In this case, the well-defined maximum starts to be observed at room temperature as showed in the inset of Fig. 3(b) suggesting that the charge carriers are delocalized. The position of this peak is found to shift towards lower frequency side (trapping behavior) with increasing temperature. This metallic-like behavior is in agreement with temperature

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Fig. 2. (a) Three thermal cycles of electrical resistivity as a function of the temperature obtained in air atmosphere. (b) Ln (r) versus the inverse of temperature corresponding to the warming curves of the three cycles obtained in air atmosphere. (c) Adsorption sensibility (rmax/rmin) of water vapor molecules as a function of time.

dependence of dc electrical resistivity observed for cycle 3, where r initially increases reaching its maximum value at T ¼ 373 K. At T ¼ 373 K, relaxation processes were not detected (completely trapped). At T > 373 K, Z00 again exhibits a maximum whose position is shifted towards higher frequency side with increasing temperature, i. e., start to be delocalized again. Above this temperature, a delocalization process is thermally activated. Furthermore, it was also observed that temperature evolution of relaxation time in both cycles (1 and 3) obeys the Arrhenius equation t ¼ t0exp(Ea/kBT), where t0 is the pre-exponential factor and Ea is the activation energy. Fig. 3(c) and (d) shows the linear fitting performed on curves Ln (t) vs 1/T for cycles 1 and 3, respectively. The values of Ea

obtained from linear fitting are 1.17 eV (for cycle 1) and 0.84 eV (cycle 3). These values are in good agreement with the results obtained from electrical resistivity measurements, suggesting that the dielectric relaxations observed in both cycles are associated to short-range hopping of oxygen vacancies [21,22]. Our results indicate that the overall behavior of electrical resistivity and activation energy value are strongly dependent on amount of humidity initially adsorbed by the sample. In this context, it is important to understand the mechanisms of adsorption/desorption of water vapor molecules on the sample's surface. The adsorption of these molecules may start by chemical adsorption that occurs on “clean” surface of the sample followed by physical adsorption on the layer initially chemisorbed. Eventually, condensation of water vapor molecules within the pores may also be observed. This condensation process depends on diameter and thickness of water layers physically adsorbed. For example, when the pore diameter is 20 Å, water vapor begins to condense at 15% relative humidity (RH) [45]. Both chemical and physical adsorption can introduce protons (Hþ) in the system. These protons can act as charge carriers increasing the electrical conductivity of the sample as observed here [46]. The complete desorption of these water vapor molecules depends strongly on the microstructure, concentration of oxygen vacancies, and porosity. Several works have shown that chemically adsorbed water layers are strongly attached to the grain surface. Consequently, its total desorption is reached only at high temperatures (550 K for TiO2 and 678 K for a-Fe2O3) [47]. On the other hand, physically adsorbed water layers are very susceptible to small variations of temperature being easily removed by reducing (increasing) the relative humidity (temperature). We have also studied the physical adsorption through electrical resistivity as a function of time measured at T ¼ 300 K, 323 K and 373 K as showed in Fig. 2(c). It is important to mention that before each measurement; the sample was subjected to argon flow during 30 min. The argon flow will remove water molecules physically adsorbed to the surface of the material. At this point, the electrical resistivity reaches its maximum value (rmax) and the measurement is initialized. Subsequently, the argon flow was turned off and humidity from air (approximately 60%) atmosphere is introduced by opening the chamber. The results showed in Fig. 2(c) reveal an abrupt decrease in electrical resistivity (logarithm scale) when the chamber was opened. By increasing the temperature, the decrease of r becomes less pronounced. For example, at t ¼ 3 min, the value of r was reduced in 99% and 75% for T ¼ 300 K and 323 K, respectively. This effect is due to competition between adsorption and desorption (caused by thermal effect) of humidity. Interestingly, the electrical resistivity exhibits an approximately constant behavior for T ¼ 373 K. This behavior is caused by complete desorption of water layers physically adsorbed which is commonly observed at this temperature. These results clearly show that the value of electrical resistivity is very sensitive to the interaction between sample and atmosphere (humidity). 3.3. Ab-initio simulations of point defects In order to further discuss and conciliate the results presented above, we have performed density functional simulations on BiFeO3 system. Such simulations allow us to assess the nature of atomic defects and impurities in terms of their electronic behavior, i.e., whether they are shallow or deep donors or acceptors of free carriers. We can also estimate the majority carriers present in the material under slightly different stoichiometric regimes. The concentration of atomic defects in a material depends strongly on the stoichiometric conditions of the material in equilibrium atmosphere. In an oxygen-poor atmosphere, oxygen vacancies are more frequent in BiFeO3 compound and it is expected the presence of

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Fig. 3. Frequency and temperature dependence of the imaginary part Z00 of the impedance measured at air corresponding to (a) cycle 1 and (b) cycle 3. For (a), the temperature of each measurement is indicated by number: (1) T ¼ 300 K, (2) T ¼ 378 K, (3) T ¼ 423 K, (4) T ¼ 473 K, (5) T ¼ 523 K, and (6) T ¼ 618 K. The curves of Ln (t) vs T1 are also showed for (c) cycle 1 and (d) cycle 3. Red line is the fitting by using the Arrhenius equation. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

impurities. Such impurities will preferably be incorporated on these vacant sites. On the other hand, in an oxygen-rich atmosphere, the presence of both oxygen vacancies and impurities will decrease in the system. Different impurities will be present in different concentrations and play different roles in these situations. Those related to elements that incorporate preferably on oxygen sites will be less common in oxygen-rich situations and vice-versa. The concentration of defects in a material is given by C ¼ n0exp(Ef/kBT) where n0 is the concentration of the incorporation site considered, kB is boltzmann's constant, T is the temperature and Ef is the formation energy of these defects. These energies can be evaluated by the equation [48]:


Ef ¼ Ed  Ep  mr þ ma þ qðEVBM þ me Þ

edges of BiFeO3 as obtained by our simulations. Oxygen vacancies are electron donors for values of the fermi level below 1 eV. This is indicated by the positive slope of the curve (VO are positively charged defects, indicating they donated electrons to the BiFeO3 matrix). If the matrix's fermi level shifts towards the conduction band, these vacancies will tend to be neutrally charged, as indicated by the flat line. Such transition of charge state occurring for me near the middle of the gap indicates a trap state for electrons. As a result,

(1)

where Ed/p are the total energies of the supercells containing the defect and the pristine one, respectively, mr,a is the atomic chemical potential of atoms removed or added to the material, q is the charge state considered for the defect and me is the chemical potential corresponding to Fermi level in the material. Fig. 4 shows plots of the formation energies of atomic defects, pertinent for our discussion, in the most stable charge state. Since oxygen vacancies and hydrogen related defects form the most important variables in our system, we restrained ourselves to simulating these defects. Other intrinsic defects in BiFeO3 have been analyzed theoretically in Ref. [49] and their results for VO showed good agreement with ours. Therefore, we extrapolated their results for the formation energies of iron and bismuth vacancies and included in Fig. 4 for comparison, considering the intermediary stoichiometric regime, when the system is not deficient in any of its constitutive elements. In Fig. 4, the vertical lines indicate the position of the band-

Fig. 4. Formation energies and charge state of atomic defects in stoichiometric regime in BiFeO3. Results for the formation energy curves for VBi and VFe where extrapolated from Ref. [49].

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electronic conductivity in oxygen-poor BiFeO3 should be less efficient than in samples with lower concentrations of oxygen vacancies. In addition to the deep levels introduced by VO, iron and bismuth vacancies are both electron acceptors [49]. This can be seen by the negative slope of the formation energy curves shown in Fig. 4. In oxygen-rich conditions, the formation energy of VO gets much higher than that of the other two vacancies, resulting in small concentrations of the former relative to the latter. Therefore, the material should have excess of holes, becoming a p-type semiconductor. For O-poor situations, formation energies of VFe and VBi get higher, whereas that of VO gets lower, as also indicated in Fig. 2 in Ref. [49]. In these conditions, the material should be nearly selfcompensated, with donor VO balancing out acceptors VFe and VBi. Few free carriers should be available at room temperature, resulting that conductivity in the intrinsic material in this condition should be dominated by extrinsic impurities or non-electronic mechanisms such ionic migration. We analyzed hydrogen impurities in BiFeO3, as indicated in Fig. 4. We observe a deep level near the middle of the gap associated to a neutral-to-acceptor transition for hydrogen atoms on interstitial sites in BiFeO3. H atoms incorporated on oxygen vacancies present two transitions, positive-toneutral and neutral-to-negative, characterizing an amphoteric behavior in BiFeO3. In neither O-rich nor O-poor conditions, however, these H impurities should significantly alter the concentration of free electrons or holes in intrinsic BiFeO3. Therefore, the effect of H impurities in the conductivity of this material should be through ionic current. In our experimental analyses, we first identified off-the-chart high resistivity in BiFeO3, which gradually decreased upon exposure to humidity and temperature increase. The measured activation energy from the Arrhenius plot is 1.53 eV in cycle 1. For a semiconductor, the activation energy in a conductivity regime dominated band-to-band thermally excited carriers is half its bandgap. The band-gap of thin films of BiFeO3 has been reported in the range of 2.66 eVe2.81 eV [50]. These values are slightly below twice the value for the measured activation energy. However, surface effects and domain walls are known to play major roles in the properties of thin films and nanoparticles of BiFeO3, which may shift effective band-to-band transition energies in this material [51,52]. Therefore, we believe that conductivity in cycle 1 is dominated by electric current and ionic conductivity is minimal. This indicates that the material started in cycle 1 in an oxygen-rich stoichiometric state. The analyses of impedance spectroscopy Fig. 3(a) showed that free electrons and holes were heavily localized in the sample during cycle 1, which explains why hole and electron conductivity was so low in that situation. We can see form Fig. 4 that there are several trap levels associated to VBi and VFe, as indicated by the kinks in the formation energy curves of these defects, which are the dominating defects at O-rich stoichiometric conditions. This explains the high resistivity measured in cycle 1. As the material is heated in a water rich atmosphere, oxygen vacancies can be formed, shifting it to a more oxygen-deficient situation. BiFeO3 is known to promote photoelectrochemical water splitting. Free holes act in the oxygen evolution reaction, according to the chemical reaction presented in eq. (2) [53]: 



2H2 O þ h /4H þ O2 ðgÞ ðOERÞ

(2)

From there, the hydrogen ion can react with the oxygen atoms present in the lattice (Olattice), generating hydroxil radicals or water molecules and oxygen vacancies (VO), according to eq. (3). 

2H þ

O lattice /

0










OH þ VO þ H /H2 O þ VO

(3)

Actually, hydrogen flux is used to enhance the photoelect-

rochemical performance of BiFeO3 by inducing excess of oxygen vacancies [54]. Humidity-rich atmosphere and high temperature treatment should have a similar effect. When the material is cooled down back to room temperature after cycle 1, resistivity does not recover off-the-chart values. Moreover, activation energy from Arrhenius plot in cycles 2 and 3 change dramatically to about 0.84 eV, which is no longer consistent with band-to-band transitions. This indicates a change in conductivity regime, most likely associated to migration of either oxygen or hydrogen-related charged impurities. Finally, impedance spectroscopy Fig. 3(b) showed that free electrons and holes were much more delocalized at room temperature in cycle 3. This again points to the change in the stoichiometric regime to an oxygen-poor situation, where trap states related to VBi and VFe become high in formation energy and therefore, less frequent. According to the impedance spectroscopy analyses, charge trapping in cycle 3 started becoming more intense at around 373 K e the boiling temperature of water e possibly because of the introduction of hydrogen related defects as well as a reincorporation of oxygen in the system, due to greater competition of absorption/desorption of H2O in the surface of the material. Our results reveal clearly that the overall behavior of electrical conductivity of our sample is strongly affected by the interaction between sample surface and atomic species present in the atmosphere. This interaction leads to chemical and physical adsorption of water vapor molecules. We believe that metallic-like behavior of r observed at low temperatures is caused by physical desorption of these water molecules. Furthermore, the anomalous variations of r observed at high temperatures may be related to chemical desorption. On the other hand, the analysis by using Arrhenius law of both electrical resistivity and impedance suggests that the value of Ea is very sensitive to amount of humidity adsorbed by the sample. 4. Conclusions We have performed a comprehensive study of electrical properties at high temperatures of single-phase BiFeO3 nanoparticles with crystallite average size of ~60 nm. These samples were placed in a vacuum desiccator and then exposed to heating/cooling cycles in atmosphere rich in humidity, while resistivity measurements were performed. Our results showed strong hysteresis in the resistivity of these nanoparticles after the first cycle, with significant changes in the measured resistivity as well as the activation energy, indicating a change in the conductivity mechanism. Electronic structure simulations indicated that the activation energy in cycle 1 is consistent with band-to-band transitions in this wide band-gap semiconductor. The results also indicate that the small value of the relaxation time of charge carriers found for cycle 1 is consistent with oxygen-rich stoichiometric material. The thermal treatment of cycle 1 induces the formation of oxygen vacancies, changing the stoichiometric state of the original samples towards oxygen-poor regime. In this situation, ionic conductivity due to oxygen vacancies as well as Hþ ions should dominate over electronic current explaining the measured change in the activation energy of the conductivity due to thermal, temperature, and atmosphere history. Acknowledgement This material is based upon work supported by the Brazilian agencies CNPq Grants No. 455092/2014-1, 153496/2016-9, and 306431/2014-9 and Fapesp under Grants No. 2010/18364-0 and 2016/09769-3.

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