Effects of vacancy co‑doping on the structure, magnetic and dielectric properties of LaFeO₃ perovskite nanoparticles

Abstract

Perovskite nanoparticles La_1-x□_xFe_1-y□_yO₃, (y = x = 0.0, 0.01 and 0.02) were prepared using the citrate combustion method. The samples were crystallized in a single phase orthorhombic structure. The energy dispersive spectroscopy (EDS) confirms that the chemical composition is very close to the nominal one. The particle size values indicated that the samples were prepared in nano scale. The selected area electron diffraction (SAED) pattern shows the good crystallinity of the samples. The magnetic hysteresis loops indicate the antiferromagnetic properties of the samples. The molar magnetic susceptibility (χ_M) of LaFeO₃ sample was enhanced by vacancy doping. The value of M_s of the sample La_0.98□_0.02Fe_0.98□_0.02O₃ increased by a factor 1.9 than that of sample LaFeO₃. The investigated samples have semiconductor-like behavior as clear from the increase in the conductivity with raising the temperature and the values of activation energies. The main conduction mechanisms in the investigated samples are the small polaron tunneling (SP) and the correlated barrier hopping (CBH).

1 Introduction

The perovskite materials have a general formula of ABO₃, where A is a rare earth element and is coordinated by 12 oxygen anions. While B cation is a 3d transition metal and coordinated to six oxygen anions. Metals with different oxidation states can be replaced by the metal ions at the A and B sites due to the great stability of the ABO₃ materials [1]. Thus, novel materials can be designed with various properties for different applications.

Lanthanum orthoferrite (LaFeO₃) has a perovskite structure with space group Pbnm and has attracted great interest due to its unique physical and chemical properties. LaFeO₃ has G-type antiferromagnetic (AFM) behavior with a Nèel temperature (T_N) ~ 730 K [2]. The ordering of ferroelectric or antiferroelectric dipole moments determines the electrical properties of LaFeO₃. The < FeO₆ > octahedron in LaFeO₃ drives the coexisting magnetic and electrical configurations. The perovskite LaFeO₃ was synthesized using different methods such as electrochemical [3], solid-state [4], and sol–gel combustion [5, 6] techniques. LaFeO₃ was prepared in nano scale easily using the citrate nitrate combustion method [7].

The perovskite structure ABO₃ contains an equal amount of A and B site cations (i.e., the A/B cation ratio = 1). When the A/B cation ratio is deviated significantly from unit, the perovskite lattice structure can still be stable. Introduction of the vacancy on both A and B cation sites tilts the < BO₆ > octahedron. The octahedral tilting lowers the coordination number of the A cation from 12 to 8 [8], reducing the tension in A–O bonds and increasing lattice energy. The presence of an octahedral tilting distortion affects not only the crystal structure, but also has a profound influence on the physical properties such as electrical conductivity, magnetic superexchange interactions and other electrical properties [9]. The nonstoichiometric A-site cation perovskite ABO₃ were developed as cathodes for solid-oxide fuel cells (SOFCs) and improve the oxygen reduction reaction (ORR) activity at high temperatures [10, 11].

Spinicci et al. [12] were prepared La_(1−ε)FeO_(3−1.5ε) with ε = 0.0, 0.1, 0.2, 0.3 using thermal low-temperature decomposition, a method which ensures a high surface area. La_(1−ε)FeO_(3−1.5ε) have been used in catalytic methane combustion in the range 300–450 °C. Scafetta et al. [13] have studied the effect of La and Fe deficiency on the c-axis lattice parameter and optical absorption in La_1-xFe_1-yO_3-δ films. Fe vacancies expanded out-of-plane lattice parameters while La deficiency leads to an increase in absorption throughout the spectral range probed. Zhu et al. [14] prepared the A-site cation deficient perovskites (La_0.95FeO_3-δ) by a simple and cost-effective process. La_0.95FeO_3-δ samples have highly-efficient bifunctional electrocatalysts for ORR and oxygen evolution reaction (OER) in alkaline solutions.

Arman et al. [15] were investigated the effect of cation deficiency on NdFeO₃. The self-doped vacancies on A and B sites of NdFeO₃ increased the saturation magnetization (M_s) of the sample Nd_0.97□_0.03Fe_0.97□_0.03O₃ by 1.5 times more than the parent sample NdFeO₃. The antiferromagnetic properties of NdFeO₃ perovskite were originated from magnetic interactions such as Nd–Nd, Nd–Fe, and Fe–Fe. While LaFeO₃ has Fe–Fe magnetic interaction only. The difference in ionic radii of La³⁺ ion (1.216 Å) and that of Nd³⁺ ion (1.163 Å) affects the tolerance factor and tilting of < BO₆ > octahedron.

In this present work, we studied the effect of vacancy co-doping on the physical properties such as the magnetic and dielectric properties of LaFeO₃ nanoparticles. The substitution of the vacancy on both A and B cation sites simultaneously is influenced by the structure, morphology and all physicochemical properties of the samples. The introduced A and B- site cation vacancies in LaFeO₃ enhanced Fe–Fe magnetic interaction and improved the hybridization between the La 3d and O 2p states. Furthermore, the presence of La and Fe defects increases the specific surface area. The vacancy doped samples are more applicable in many applications, such as the heavy metal removal from wastewater, cathode catalysts for lithium-oxygen (Li-O₂) batteries [16,17,18].

2 Experimental techniques

The perovskite samples La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ y, x ≤ 0.02, step 0.01) were prepared by citrate nitrate combustion method [14, 16]. The precursor metal nitrates (purity 99.9%, Sigma–Aldrich) were mixed in stoichiometric ratios with an equivalent molar ratio of citric acid in aqueous media to enhance the homogeneous mixing. To produce La and Fe vacancies, the La(NO₃)₃∙6H₂O and Fe(NO₃)₃∙9H₂O are stoichiometrically decreased by 1% and 2% to generate nonstoichiometric La and Fe defects. The LaFeO₃ perovskite has a stable structure upto 30% nonstoichiometric La and Fe defects [12, 16]. To ensure complete complexation, the pH value was adjusted to 7 using the ammonia solution as a fuel. The (citric acid/nitrate) ratio is kept constant for all samples and is equal to 1. The samples were heated on a hot plate to obtain the final product that gives a fluffy powder with a single-phase orthorhombic structure. Thermal dehydration resulted in a highly viscous liquid. The citrate nitrate combustion method is simple, low cost, and quick.

X-ray powder diffraction (XRD) was carried out to study the crystal structure of the samples using a Proker D8 advanced X-ray diffractometer with Cu-Kα radiation (λ = 1.5418 Å). The crystalline phases of the samples were identified using the International Centre for Diffraction Data (ICDD) card number 74–2203.

The samples were characterized by scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) to study the morphology of the samples using OXFORD INCA PentaFETX3-England. The shape and size of the particles were studied using a high-resolution transmission electron microscope (HRTEM) model (JEOl-2100). The magnetic properties of the samples were studied by two techniques: the first is the H-M hysteresis loop using a vibrating sample magnetometer (VSM; 9600–1 LDJ, USA) with a maximum applied field of 20 kOe at room temperature. While the second technique is measurements of dc magnetic susceptibility (χ_M) using Faraday’s method [19] at different magnetic field intensities as a function of absolute temperature.

The powdered samples were pressed to obtain pellets using a press of value 8 × 10⁵ N/m² [7]. The two surfaces of each pellet were checked for good conduction by coating the surfaces with silver paste. The electrical properties were measured on pellets of samples using the LCR meter (Hioki model 3532 Japan). The dielectric constant (ε′) and ac conductivity of the samples were measured as a function of temperature in the frequency range 100 Hz to 5 MHz. The data was triplet and taken the average value.

3 Results and discussion

Figure 1 illustrates the XRD patterns of the samples that have the formula of La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ y, x ≤ 0.02, step 0.01). The prepared samples were crystallized in a single phase orthorhombic structure with space group Pbnm. The data was indexed with ICDD card number 74–2203. The main peak (100%) of LaFeO₃ is observed for (200) plane at 2θ = 34.5°. The lattice parameters (a, b and c) were calculated based on orthorhombic unit cell from Eq. (1) and listed in Table 1. The crystallite size (L) was determined using Scherrer’s equation which is given by Eq. (2). [20]

$$\frac{1}{{d^{2} }} = \frac{{h^{2} }}{{a^{2} }} + \frac{{k^{2} }}{{b^{2} }} + \frac{{l^{2} }}{{c^{2} }}$$

(1)

$$L = \frac{0.9 \lambda }{{\beta \cos \theta }},$$

(2)

Fig. 1

XRD of the samples La_1-x□_xFe_1-y□_yO₃ as compared with ICDD card number 74–2203. The main peak (100%) is observed for (200) plane at 2θ = 34.5°

Table 1 Values of the lattice parameters, the unit cell volume (V), the theoretical density (D_x), the crystallite size from XRD (L), Williamson–Hall crystallite size (L_WH), micro strain (□), the particle size (L) and the tolerance factor for the investigated samples

where L is the average crystallite size, θ is the Bragg angle, λ is the wave length of the X-ray radiation (λ = 1.5406 Å), and β is the full width at half maximum intensity of the powder pattern peak. The values of L were listed in Table 1 and indicated that the samples are prepared in nano scale.

The crystallite size (L_WH) of the samples and the average lattice strain (ε) were calculated using the well-known Williamson–Hall equations [21]:

$$\beta \cos \theta = \, 4\varepsilon \, \sin \theta + \frac{K\lambda }{{L_{{{\text{WH}}}} }},$$

(3)

where K is the dimensionless shape factor (K = 0.9) and θ is the Bragg angle. Figure 2 illustrates the relation between βcos$\theta$ and 4sin$\theta$. The value of ε was determined from the slope of the plot in Fig. 2 and reported in Table 1. The crystallite size (L_WH) was calculated from the intercept of the fitting line with y-axis and its values confirm that the samples were crystallized in nano scale. The W–H plot shows the negative value of the strain and indicates the compressive nature of it which leads to lattice shrinkage [18].

Fig. 2

The Williamson–Hall plot for LaFeO₃ sample. The crystallite size (L_WH) was calculated from the intercept of the fitting line with y-axis

The unit cell volume and the theoretical density were calculated using the following equations, respectively.

$$V \, = \, abc$$

(4)

$$D_{x} = \frac{ZM}{{{\text{N}}_{{\text{A}}} V}},$$

(5)

where Z (Z = 4) is the number of molecules per unit cell, N_A is the Avogadro’s number, M is the molecular weight, and V is the unit cell volume. The values of D_x of the vacancy-doped samples are less than that of LaFeO₃ owing to the decrease of La³⁺ and Fe³⁺ content with investigated samples. The values of experimental density were calculated using the relation D_exp = $\frac{m}{V}$ and reported in Table 1. The values of D_exp are smaller than the theoretical density owing to oxygen deficiency.

The tolerance factor was calculated using the following equation [22]:

$$t \, = \frac{{r_{{\text{A}}} + r_{{\text{O}}} }}{{\sqrt 2 \left( {r_{{\text{B}}} + r_{{\text{O}}} } \right)}},$$

(6)

where r_A, r_B and r_O are the ionic radii of the A, B and oxygen ions respectively. The values of tolerance factors were listed in Table 1. The tolerance factor (t) is less than 1 which indicates that tilting of the < BO₆ > octahedra occurs for all the samples. The tolerance factor indicates the orthorhombic structure of the samples La_1-x□_xFe_1-y□_yO₃ and these results agree with the data obtained from XRD analysis. The values of the tolerance factor decrease with increasing the vacancy content (x and y) at the expense of La³⁺ and Fe³⁺ ions, which means that the distortion of the perovskite unit cell increases. This means that the Fe–O–Fe angle decreases with increasing the vacancy content (x and y). On the other hand, the tilting angle decreases. Besides, the values of t (0.9550 < t < 0.9514), indicated that the structure distorts.

To confirm the chemical composition of the samples, the energy dispersive spectroscopy EDS analysis was performed. Table 2 illustrates the EDS data of the investigated samples. The figure shows that the prepared samples are nearly equal to the starting chemical composition without any noticeable impurities. The weight percentage (wt%) of elements (La, Fe, and O) are calculated theoretically from the chemical formula La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ y, x ≤ 0.02, step 0.01) and that obtained from EDS elemental analysis are shown in Table 2. The results confirm that the chemical composition is very close to the nominal one. The small variation between the values of weight percentage (wt%) (La, Fe, and O) due to oxygen deficiency and converting some Fe³⁺ to Fe²⁺.

Table 2 The composition of weight percentage (wt%) for the samples (a) LaFeO₃, (b) La_0.99□_0.01Fe_0.99□_0.01O₃ and (c) La_0.98□_0.02Fe_0.98□_0.02O₃ as obtained from EDS

Figure 3 shows HRTEM images of the samples La_1-x□_xFe_1-y□_yO₃. The images illustrate the orthorhombic structure of nanocrystalline samples. The values of particle size were listed in Table 1 and indicated that the samples were prepared in nano scale. The values of particle size agree with those obtained from XRD. The particles are agglomerated as shown in Fig. 3 owing to the magnetic properties of the samples. Different brightness from one crystallite to another appears as a result of different orientations of the crystal structure. The insets in Fig. 3 show the d-spacing of the investigated samples with values 0.26, 2.5 and 0.49 nm of the samples La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ x and y ≤ 0.02, step 0.01), respectively. The selected area electron diffraction (SAED) pattern of the investigated samples is illustrated in the insets in Fig. 3. The patterns show the good crystallinity of the samples.

Fig. 3

HRTEM images for the samples a LaFeO₃, b La_0.99□_0.01Fe_0.99□_0.01O₃ and c La_0.98□_0.02Fe_0.98□_0.02O₃ with average particle size 29 nm, 38 nm and 62 nm respectively. The insets illustrate d-spacing and SAED for the samples

Figure 4 shows the relation between the molar magnetic susceptibility (χ_M) and the absolute temperature for LaFeO₃ nanoparticles. The sample has antiferromagnetic behavior where the values of χ_M decreased rapidly with temperature until the Néel temperature then decreased slowly with temperature. The magnetic interactions originated from the superexchange interaction between the Fe³⁺ ions, while there is no magnetic interaction between La³⁺ ion and Fe³⁺ ion because La³⁺ ion is not magnetic. By increasing the vacancies in the sample, the tilting of < FeO₆ > octahedron increases, leading to a weak ferromagnetic ordering along with other crystal directions [23].

Fig. 4

The relation between the molar magnetic susceptibility and the absolute temperature for the sample LaFeO₃, at different magnetic field intensities. The sample has antiferromagnetic behavior

The comparative study between the vacancy doped samples in the molar magnetic susceptibility at the magnetic field intensity 1100 Oe is illustrated in Fig. 5. χ_M increases with increasing the vacancies on the A and B sites. The reasons for the increase of χ_M are:

• (1) Some of the Fe³⁺ ions convert to Fe⁴⁺ ions, which create the three magnetic interactions (antiferromagnetic between Fe³⁺–Fe³⁺ and Fe⁴⁺–Fe⁴⁺, ferromagnetic between Fe³⁺–Fe⁴⁺).

• (2) The tilting of < FeO₆ > octahedron affects the bond length and strength between Fe–O ions.

Fig. 5

A comparative study of the molar magnetic susceptibility for the samples La_1-x□_xFe_1-y□_yO₃ at magnetic field intensity 1100 Oe. The vacancies on the A and B sites enhanced the magnetization of the samples

The Curie–Weiss law is satisfied in the investigated samples. The straight lines in the paramagnetic region of LaFeO₃ sample at 1100 Oe in the temperature range of 850–910 K are shown in Fig. 6.[15] The values of the Curie constant (C), Curie–Weiss constant (θ), and the effective magnetic moment (µ_eff) are calculated from the reciprocal of magnetic susceptibility with absolute temperature and reported in Table 3 using the following equations:

$$C = \frac{1}{{{\text{slope}}}}$$

(7)

$$\mu_{{{\text{eff}}}} = 2.83\sqrt C .$$

(8)

Fig. 6

Dependence of the reciprocal molar magnetic susceptibility in paramagnetic region for LaFeO₃ at magnetic field intensity 1100 Oe

Table 3 Values of the Curie constant (C), the Curie–Weiss constant (θ), the effective magnetic moment (μ_eff) and the Neel temperatures of the samples La_1-x□_xFe_1-y□_yO₃ at magnetic field intensity of 1100 Oe

The presence of a ferromagnetic component in the samples was illustrated by the positive values of θ. The Néel temperature of the samples was determined from the plots of (dM/dT) vs T where M is the magnetization.

Figure 7 illustrates the hysteresis loops of the samples La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ x, y ≤ 0.02, step 0.01) at room temperature. The S-shape of the hysteresis loops indicates the presence of a weak ferromagnetic component in addition to the main antiferromagnetic order in the samples. The magnetic parameters such as saturation magnetization (M_s) and remnant magnetization (M_r) are listed in Table 4. The value of saturation magnetization M_s of LaFeO₃ agrees with the data published by Gabala et al. [24]. The magnetization of the LaFeO₃ sample was enhanced by vacancy doping. The value of M_s of the sample La_0.98□_0.02Fe_0.98□_0.02O₃ increased by the factor 1.9 compared with the sample LaFeO₃. The improvement of magnetic parameters for vacancy doped samples can be ascribed to several reasons as follows: (i) The tilting of < FeO₆ > induces a variation in the Fe–O–Fe bond angle and distance; (ii) The existence of Fe³⁺ (HS) (t_2g³ e_g²) and Fe⁴⁺ ions in high spin state (HS) Fe⁴⁺ (t_2g³ e_g¹) [7]. The inset in Fig. 7 illustrates the hysteresis loop of the parent sample (LaFeO₃) has a clear shift from the origin due to the exchange bias effect. The exchange bias field (H_EB) was calculated using Eq. (9):

$$H_{{{\text{EB}}}} = \, - \, {{\left( {H_{{{\text{Left}}}} + \, H_{{{\text{right}}}} } \right)} \mathord{\left/ {\vphantom {{\left( {H_{{{\text{Left}}}} + \, H_{{{\text{right}}}} } \right)} 2}} \right. \kern-\nulldelimiterspace} 2},$$

(9)

Fig. 7

VSM of the samples La_1-x□_xFe_1-y□_yO₃ at room temperature. The inset shows the shift of hysteresis loop from the origin for the parent sample (LaFeO₃)

Table 4 Values of the saturation magnetization (M_s), the coercive field (H_c), the remanence magnetization (M_r) and the exchange bias field (H_EB) for the samples

where H_Left and H_right are the intercept of the magnetization on the −ve and +ve on the H-axis. The values of exchange bias were reported in Table 4. The exchange bias is present due to the exchange coupling between the ferromagnetic-like shell and the antiferromagnetic core in the nanoparticle [25].

Figure 8 illustrates the dependence of the dielectric constant on the absolute temperature of the sample La_0.99□_0.01Fe_0.99□_0.01O₃. The trend of ε′ of the sample is the general trend of dielectric materials where ε′ increases with increasing the temperature and decreases with increasing the frequency.

Fig. 8

Dependence of dielectric constant (ε′) on absolute temperature T (K) of the sample La_0.99□_0.01Fe_0.99□_0.01O₃ as a function of frequencies. ε′ values increases with increasing the temperature and decreases with increasing the frequency

Figure 9 shows the dependence of lnσ on the reciprocal of absolute temperature (1000/T) as a function of frequencies for the sample La_0.99□_0.01Fe_0.99□_0.01O₃. The plot data obeys the well-known Arrhenius equation [26]:

$$\sigma \, = \, \sigma_{o} \exp \, \left( { - E/kT} \right),$$

(10)

Fig. 9

The dependence of lnσ on the reciprocal of the absolute temperature as a function of frequencies of the sample La_0.99□_0.01Fe_0.99□_0.01O₃

where E is the activation energy, k is the Boltzmann’s constant and T is the absolute temperature. The presence of two different straight lines indicates the presence of two conduction mechanisms in the investigated samples. The activation energies were calculated and reported in Table 5 using the Arrhenius equation at frequency 5 MHz. The activation energy in the high temperature region (E_II) is greater than that of low temperature (E_I). The presence of two activation energies is due to the presence of two conduction mechanisms, which are the small polaron (SP) tunneling and the correlated barrier hopping (CBH) as shown in the next section.

Table 5 Values of the activation energy E_II (at high temperature region) and E_I (at low temperature) of the samples La_1-x□_xFe_1-y□_yO₃ at frequency 5 MHz

The samples have semiconducting like behavior as clear from the increase in the conductivity with an increase of temperature. The values of activation energies confirm the semiconducting trend of the samples. The cation deficient in perovskite LaFeO₃ played an important role in the enhancement of the conductivity. Introducing of vacancies in LaFeO₃ leads to formation oxygen deficiency and converting some of Fe³⁺ ions to Fe⁴⁺ ions (t_2g³ e_g¹ configuration). This will bring a localized electronic structure and lead to an inherent semiconductor nature [16].

Figure 10 illustrates the dependence of lnσ on lnω at different temperatures. The relation between lnσ and lnω obeys the universal power law. [27]

$$\sigma_{{{\text{ac}}}} \left( \omega \right) = A \omega^{s} ,$$

(11)

Fig. 10

The relation between the lnσ as a function of lnω at different temperatures of the sample La_0.99□_0.01Fe_0.99□_0.01O₃. The data obeys the universal power law

where σ_ac is the ac conductivity, ω = 2πf is the angular frequency, A is a temperature dependent constant and S is the frequency dependent exponent. Figure 11 shows the relation between the frequency exponent factor (S) and T. In the 1st temperature region (300–450 K), the increase of S with T which indicates the small polaron (SP) tunneling in this region. This originated from converting some of the Fe³⁺ to Fe⁴⁺ ions, also, the electron conduction in perovskite-type ferrites takes place through Fe⁴⁺–O–Fe³⁺ [28]. While in the 2nd temperature region (450–800 K), S decreases with T which illustrates that the correlated barrier hopping (CBH) is the main mechanism. The electrons accept enough energy to hop between the O 2p and the Fe 3d orbitals. The hopping of electrons is affected by the Fe–O distance and the Fe–O–Fe angle. The values of the activation energy (E_Π) increased in the off-stoichiometric samples owing to the presence of oxygen vacancies [13].

Fig. 11

Relation between the frequency exponent factor (S) and absolute temperature (T) for the sample La_0.99□_0.01Fe_0.99□_0.01O₃. The main conduction mechanisms are the small polaron tunneling (SP) and the correlated barrier hopping (CBH)

4 Conclusion

Orthoferrite of formula La_1-x□_xFe_1-y□_yO₃, (0.0 ≤ y, x ≤ 0.02, step 0.01) were well synthesized using the citrate combustion technique. XRD revealed that the investigated samples crystallized in a single phase orthorhombic structure with space group Pbnm. The particle size of LaFeO₃ increased gradually with increasing the vacancy content. The magnetic hysteresis loops indicate the antiferromagnetic properties of the samples. χ_M of La_0.98□_0.02Fe_0.98□_0.02O₃ increased by a factor 1.7 times more than of LaFeO₃ sample. The investigated samples have semiconducting-like behavior as clear from the increase in the conductivity with the temperature and the values of activation energies. The small polaron tunneling (SP) and the correlated barrier hopping (CBH) are the main conduction mechanisms in the samples.