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. 2020 Nov 4;5(45):29292–29299. doi: 10.1021/acsomega.0c04076

Ho and Ti Co-Substitution Tailored Structural Phase Transition and Enhanced Magnetic Properties of BiFeO3 Thin Films

Mingjie Sun †, Liang Bai †, Wenjing Ma †, Yanqing Liu †,*, Junkai Zhang †,‡,*, Jinghai Yang †,*
PMCID: PMC7675960  PMID: 33225160

Abstract

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The polycrystalline thin films of BiFeO3 (BFO) and Bi0.90Ho0.10Fe1–xTixO (x = 0, 0.025, 0.05, 0.10, 0.15, and 0.20) were successfully synthesized by the simple sol–gel method. X-ray diffraction and Raman spectra revealed the substitution of Bi and Fe by Ho and Ti, respectively, and correspondingly a structural phase transition from the rhombohedral phase to orthorhombic phase. The field-emission scanning electron microscopy and transmission electron microscopy images indicated that the average size of the particles was decreased and the surface homogeneous agglomeration was enhanced with the increased concentration of Ti to x = 0.05. The X-ray photoelectron spectroscopy measurements illustrated that Fe3+ and O2– ions tended to increase with the Ti concentration increase, which accounted for the enhanced super-exchange interaction between Fe3+ and O2–. Because of the reduced concentration of oxygen vacancies, Ho and Ti ions with a smaller ionic radius and denser surface structure, the Ho and Ti co-substituted films with an appropriate concentration of Ti (x = 0.05) showed an optimal saturation magnetization (Ms) of 44.23 emu/cm3 and remanent magnetization (Mr) of 4.62 emu/cm3, which were approximately 1.8 times and 1.9 times than that of the pure BFO, respectively. This work opened up an effective way to modulate the structure and properties of BFO-based materials.

Introduction

Magnetoelectric multiferroics have drawn great attention because they exhibit ferroelectric and ferromagnetic properties in the single-phase simultaneously and the linear coupling between them,1−3 and they opened up prospective opportunities for plenty of potential multifunctional applications in modern technology, such as high energy density capacitors, spintronics devices, and gas sensing applications.4,5 In most of multiferroic materials, BiFeO3 (BFO) with a distorted perovskite (ABO3) structure exhibits a ferroelectricity temperature (TC ∼ 1103 K) and an antiferromagnetism temperature (TN ∼ 643 K) simultaneously beyond the room temperature, which has attracted more interest.6,7 The ferroelectric properties of BFO arise from the stereochemical activity of Bi3+ cations with 6s2 lone pair electrons, and the ferromagnetic properties stem from the Fe–O–Fe super-exchange interactions.8−11 Because of the valence fluctuation of Fe ions and oxygen vacancies, BFO has experienced some problems including high leakage current, high dielectric loss, and the secondary phase, limiting its practical applications.5,12 Moreover, the G-type antiferromagnetic spiral modulated the spin structure of BFO with a long-wavelength period of 62 nm, preventing the observation of any net magnetization and linear magnetoelectric effect.4,9

Great efforts have been carried out to resolve these shortcomings of BFO. Many investigations have manifested that the substitutions of rare-earth (RE) elements (La, Gd, Ho, Eu, etc.) at the A-site, transition element (Ni, Zn, Ti, Mn, etc.) at the B-site, and A–B-site co-substitution (Nd–Mn, Gd–Ni, Er–Co, etc.) in BFO show the valid enhancement of ferromagnetic and ferroelectric properties.3,11,13−21 The RE element Ho with the ionic radius (1.015 Å) can occupy easily the position of Bi (1.08 Å) in the BFO lattice. Ho3+ doping could suppress the evaporation of Bi, reduce the particle size, and stabilize the perovskite structure, and finally, enhance the magnetic properties.22,23 On the other hand, doping the transition-metal element Ti into B-site of BFO could compensate the charge concentration and control the fluctuation of Fe2+/Fe3+ valence. It has been reported that the Ti4+ substitution significantly reduced the leakage current.24 Moreover, the Ti4+ substitution could change the bond angle of Fe–O–Fe and distort FeO6 octahedra, leading to the increased magnetization of BFO.25 Recently, Liu et al. observed that the enhancement of magnetic properties for BFO thin films by the Ho element doping at the A-site in the perovskite lattice.14 Singh et al. obtained the net magnetization of BiPrFeTiO3 samples increased remarkably with the addition of Ti4+ ions at the B-site.15 However, to date, there are only few studies on the Ho and Ti co-substitution in BFO thin films.

In this work, the Ho and Ti co-substitution (B0.90H0.10F1–xTxO, x = 0, 0.025, 0.05, 0.10, 0.15, and 0.20) thin films were synthesized by a simple sol–gel technique, and the changes of the structural, morphological, and magnetic properties with the increase in Ti-substitution concentration were investigated. Moreover, the underlying mechanism of enhanced magnetic properties has been investigated and discussed in detail.

Experimental Details

The polycrystalline BFO and B0.90H0.10F1–xTxO thin films were fabricated on Si(100) substrates via a simple sol–gel technique. The ultra-high purity Bi(NO3)3·5H2O, Ho(NO3)3·5H2O, Fe(NO3)3·9H2O, and Ti(OC4H9)4 with suitable proportions were dissolved in ethylene glycol and stirred well for several hours. The molar concentration of each solute in the obtained stable precursor solutions was fixed 0.2 mol/L. It is worth noting that 3 mol % excess Bi(NO3)3·5H2O was added in the experiment because of the evaporation of the Bi element. All the solutions were mixed appearing dark red in color and transparent under stirring continuously at room temperature. Then, the mixture was kept still in the air to age, as the precursor solution. The uniform precursor solution was spin-coated on the Si(100) substrate at 500 rpm for 3 s and 4000 rpm for 20 s and each composition was pre-annealed at 350 °C for 6 min, then the above steps were repeated several times in order to reach the desired thickness. These dried thin films annealed at 500 °C in air for 1 h and finally were cooled slowly to room temperature. Figure 1 shows a sketch of the process of the preparation of the above compounds.

Figure 1.

Figure 1

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Schematic diagram of the process for preparing thin film via sol–gel technology.

Characterization

The structural characteristics of the BFO and B0.90H0.10F1–xTxO thin films were analyzed by X-ray diffraction (XRD, Rigaku Corporation, Tokyo, Japan) with Cu Kα radiation. A field-emission scanning electron microscope ( Model Hitachi S-570, JEOL Ltd., Tokyo, Japan) and a transmission electron microscope ( 200 keV, JEM-2100HR, JEOL Ltd., Tokyo, Japan) were employed to obtain the surface morphology, component analysis, and microstructure of the samples. Raman spectra were carried out in backscattering geometry using a Renishaw MicroRaman spectrometer (Renishaw, London, UK). An Escala 250XI X-ray photoelectron spectrometer (Thermo Fisher Scientific, Waltham, MA, USA) was applied to study the chemical bonding and the valence states of all compounds. The magnetic hysteresis (M–H) loops were measured at room temperature using a Lake Shore 7407 vibrating sample magnetometer (VSM, Lake Shore, Columbus, OH, USA).

Results and Discussion

Figure 2a shows the XRD patterns of BFO and B0.90H0.10F1–xTxO thin films at room temperature. All diffractions of the pure BFO thin film are indexed well with the standard XRD patterns (#71-2494), showing that it possesses a rhombohedral perovskite structure with a space group of R3c. The absence of constituent oxide characteristic peaks and impurity phases assures the formation of single-phase compositions.2 The amplified images for (104) and (110) peaks near 32° are presented in Figure 2b. Two peaks exhibit an obvious separation in BFO but tend toward coalescing in the B0.90H0.10F1–xTxO thin films. Finally, the (110) and (014) peaks merge into a single peak and shift to the higher 2θ degree. According to the Bragg’s law, the lattice shrinkage of the doped samples can be expected. These indicate that the intrinsic perovskite structures transform from the rhombohedral phase (R3c) into orthorhombic phase (Pbnm) by the doping of Ho/Ho and Ti, which show that the Bi (1.08 Å) and Fe (0.645 Å) sites are replaced by smaller ionic radius elements Ho (1.015 Å) and Ti (0.605 Å).24,26 As x further increases up to 0.15 and 0.20, the diffraction peaks widen gradually, owing to the occurrence of structure transition from the rhombohedral (R3c) phase to a mixed state of cubic (Pm3̅m) and orthorhombic (Pbnm) phases. This is ascribed to the lattice distortion of BFO caused by the increase in chemical pressure of Ho and Ti.24,26,27 The structural stability of the doped samples by Goldsmith tolerance factor (t) of the perovskite structure was reckoned by using this formula

graphic file with name ao0c04076_m001.jpg 1

Figure 2.

Figure 2

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(a) XRD patterns of BFO and Bi0.90Ho0.10Fe1–xTixO thin films and (b) magnified patterns of BFO and Bi0.90Ho0.10Fe1–xTixO thin films around 32°.

The substituted samples conform to the range of the ideal tolerance factor (t) from 0.77 to 0.99 in order to stabilize the perovskite structure.28 In this work, the t values of all samples are estimated at the range of 0.8–0.9, demonstrating that the substitution concentration we adopted is feasible in theory.

On account of the sensitivity of Raman spectroscopy in the crystalline structure symmetry, we can take advantage of it to detect the information of structural and phase transitions for BFO with the introduction of Ho and the increase in the Ti concentration. The vibrational changes of BFO and B0.90H0.10F1–xTxO thin films were analyzed by using Raman spectra at room temperature, as shown in Figure 3. The theoretical results have indicated that the rhombohedral BFO (R3c) possesses 13 optical phonon modes, including 4 longitudinal optical modes (A1) and 9 transverse optical mode (E).29,30 A1 modes in the lower frequency region are attributed to the Bi–O vibrations and E modes in the higher frequency region decrease to Fe–O vibrations.24 It is observed that the Raman spectra of all substitution samples slightly shift to a higher frequency compared with the pure BFO. Meanwhile, the intensities of A1-2 and A1-3 modes around 168.01 and 216.66 cm–1 are reduced with the doping of Ho3+ ions, meaning that the stereochemical activity of Bi 6s2 lone pair electrons decreases. Besides, no more obvious change in A1 mode occurs because of the addition of Ti4+ ions, yet as the Ti4+ ion concentration gradually increased, the E-9 modes around 609.77 cm–1 appear, being enhanced slightly because of the alterations of Fe–O bonds caused by Ti4+ ions substituted. These mean that the variety of E modes can be considered as the FeO6 octahedra bending and stretching.24,30,31Figure 3b shows the simulated diagram of the change of Fe–O bond.

Figure 3.

Figure 3

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(a) Raman scattering diagrams of BFO and Bi0.90Ho0.10Fe1–xTixO thin films at the room temperature and (b) simulated diagram of the change of the Fe–O bond.

The surface morphologies of BFO and B0.90H0.10F1–xTxO films are shown in Figure 4. From the vertical views, the wide pores are observed for the BFO thin films in Figure 4a. However, the BH0.10FO thin film (Figure 4b) exposes less pores and a more compact morphology compared with the pure BFO thin film. In Figure 4c–e, the B0.90H0.10F1–xTxO thin films exhibit a homogeneous agglomeration and denser surface structure with the increase in x from 0.025 to 0.10. This manifests that the Ho3+ and Ti4+ ions are highly effective in inhibiting particle growth. As x further increases up to 0.15 and 0.20, B0.90H0.10F1–xTxO thin films exhibit less dense and larger porosity morphology, as shown in Figure 4f,g. This may be caused by the unsustainable crystallization, consistent with the observation of XRD in this work.14,18,32 The thicknesses of the BFO and B0.90H0.10F1–xTxO (x = 0, 0.025, 0.05, 0.10, 0.15, 0.20) are 258, 228, 224, 194, 209, 221, and 227 nm, respectively, which can be seen from the cross-sectional SEM images of the samples in the inset of Figure 4. As shown in Figure 5, the grain size histogram and Gaussian fitting certify that the grain size decreases with the increase of x concentration up to 0.05 and then increases with x up to 0.10–0.20. The grain size is reduced because of the inhibition of oxygen vacancy concentration, which makes the movement of oxygen ions slower, decreasing the growth rate of crystal grains.18,32 As x further increases up to 0.15 and 0.20, the excessive Ti4+ ions do not affect the grain growth. Because of this, BH0.10FT0.15O and BH0.10FT0.20O thin films exhibit a less dense and larger porosity morphology.

Figure 4.

Figure 4

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SEM images of surface morphologies for (a) BFO; (b) BH0.10FO; (c) BH0.10FT0.025O; (d) BH0.10FT0.05O; (e) BH0.10FT0.10O; (f) BH0.10FT0.15O; and (g) BH0.10FT0.20O films.

Figure 5.

Figure 5

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Histograms regarding particle size distribution and Gaussian fitting curve of BFO and B0.90H0.10F1–xTxO (x = 0, 0.025, 0.05, 0.10, 0.15, 0.20) thin films.

Figure 6a–c shows the TEM images of BFO, BH0.10FO, and BH0.10FT0.05O thin films. The grain size slightly decreases with the substitution of the Ho3+ and Ti4+ ions, in accordance with the results of SEM images in this work. The HRTEM images are shown in the inset of Figure 6a–c, which are taken from the circular region. The lattice spacing of typical crystalline regions of BFO, BH0.10FO, and BH0.10FT0.05O is about 0.401, 0.394, and 0.389 nm, respectively, corresponding to the (012) plane of standard BFO (#71-2494).12

Figure 6.

Figure 6

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(a–c) low magnification TEM images of BFO; BH0.10FO; BH0.10FT0.05O thin films, and the insets in (a–c) show high-resolution TEM images of all thin films.

As shown in Figure 7, in order to illustrate the enhanced magnetization, we further analyzed the XPS patterns of pure BFO and BH0.10FT0.05O samples. XPS analysis was performed to predict the oxidation states of the cation in the samples and determine the variation of magnetic properties according to the Fe ions valence state. Figure 7a–h displays the XPS spectra of the Bi 4f, Fe 2p, and O 1s areas for BFO and Bi 4f, Fe 2p, O 1s, Ho 4d, and Ti 2p for BH0.10FT0.05O thin films, respectively. Herein, the C 1s peak (284.6 eV) is taken as the reference for calibrating the binding energy values of XPS spectra.7 Bi 4f peaks are resolved in 4f7/2 and 4f5/2 and have a spacing of approximately 5.3 eV in both pure BFO and BH0.10FT0.05O (Figure 7a,b), which confirms that the existence of the Bi–O bonds and the Bi3+ valence states.18,32Figure 7c,d displays the fitted Fe 2p areas for BFO and BH0.10FT0.05O thin films. The Fe 2p doublet can be fitted with two peaks of Fe 2p3/2 and Fe 2p1/2, which are located at 710.47 and 724.15 eV, respectively, for pure BFO and located at 710.14 and 723.56 eV, respectively, for BH0.10FT0.05O. The Fe 2p3/2 and Fe 2p1/2 peaks are well fitted according to the two portions of Fe3+ ions and Fe2+ ions, illustrating that Fe3+ and Fe2+ ions coexist in the BFO and BH0.10FT0.05O thin films.7,18 Furthermore, the concentration ratio of Fe3+/Fe2+ increases from 0.45 to 0.97 with Ho and Ti co-substitution. The increase in the Fe3+ concentration is beneficial to enhance the super-exchange effect of Fe–O–Fe, thereby improving the magnetic properties.33Figure 7e,f shows the fitted O 1s spectra for the BFO and BH0.10FT0.05O thin films. The broad O 1s peaks consist of two sub-peaks, which are assigned to lattice oxygens (green lines) and oxygen vacancies (orange lines), respectively. It is believed that the ratio of oxygen vacancies to lattice oxygens is associated with the oxygen vacancy concentration, and the ratio decreases obviously from 0.87 to 0.33.34Figure 7g shows that the characteristic Ho 4d peaks consisted of Ho 4d (159.0 eV) and Ho 4d5/2 (164.3 eV), which confirms the existence of the Ho3+ valence state. This indicates that a part of Bi3+ ions have been replaced by Ho3+ ions in the BH0.10FT0.05O thin film.22Figure 7h reveals that Ti 2p peaks are resolved in Ti 2p3/2 and Ti 2p1/2, which are located at 475.86 and 465.65 eV, respectively. This result verifies that the BH0.10FT0.05O thin film contains Ti4+ because the binding energies are in accordance with the typical of TiO2.35 Ti4+ ions occupying the Fe-site of BFO could compensate the charge concentration and control Fe2+/Fe3+ valence fluctuation.15 These confirm that the Ho and Ti co-substitution prevents the formation of oxygen vacancies in the BFO structure. According to the above results, it is referred that the increased Fe3+ concentration and the lower oxygen vacancies play the active role of establishing the BFO structure and changing its magnetic properties.

Figure 7.

Figure 7

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XPS spectra of the pure BFO thin film in the binding energy regions of (a) Bi 4f; (c) Fe 2p; and (e) O 1s; XPS spectra of the BH0.10FT0.05O thin film in the binding energy regions of (b) Bi 4f; (d) Fe 2p; (f) O 1s; (g) Ho 4d; and (h) Ti 2p.

To study the effect of Ho and Ti co-substitution on the magnetic properties, Figure 8a presents the magnetic hysteresis (M–H) loops for BFO and B0.90H0.10F1–xTxO thin films at 10 kOe. Because of the Ho and Ti co-substitution, both saturation magnetization (Ms) and remanent magnetization (Mr) of BH0.10FT0.05O are enhanced remarkably. The values of Ms and Mr for BFO, BH0.10FO, BH0.10FT0.025O, BH0.10FT0.05O, BH0.10FT0.10O, BH0.10FT0.15O, and BH0.10FT0.20O thin films are shown in Figure 8b. The values of Ms are 25.15, 30.36, 31.37, 44.23, 38.52, 36.40, and 36.27 emu/cm3. The values of Mr are 2.45, 3.03, 3.58, 4.62, 3.95, 3.48, and 3.21 emu/cm3. The zoomed-in view of the magnetic field for the Mr is shown in the inset of Figure 8a. There is a slight increase in the Mr of BH0.10FO thin film compared with that of pure BFO. However, the magnetic properties of B0.90H0.10F1–xTxO thin films are enhanced obviously with the increase in the doping concentration of Ti. When the concentration reaches x = 0.05, Ms and Mr reach the maximum values, which are approximately 1.8 times and 1.9 times, respectively, than that of the pure BFO thin film. As further increasing the concentration up to 0.10–0.20, both Mr and Ms tend to decrease slightly, agreeing well with the results of XRD measurements. The macroscopic magnetization of BFO is eliminated because of the incommensurate cycloid spin modulation with a 62 nm wavelength.36,37 The introduction of the Ho element with a smaller ions radius suppresses the inherent helical cycloid space spin structure of BFO, which is useful to release the potential magnetic characteristics of the BFO system.13,36 It is illustrated that when the Fe ions are substituted by the Ti ions, the Fe–O–Fe bond angles and Fe–O bond lengths are changed a little, as well as the FeO6 octahedra near a cationic vacancy is slightly distorted. Probably, the Fe3+ ions in the distorted FeO6 octahedra are sources of a magnetic performance, which causes an effective destruction of the space-modulated spin structure.34,37 Thus, Ho substituted on the Bi-site and Ti substituted on the Fe-site of BFO give rise to the changes in the bond length of Fe–O and distortion of FeO6 octahedra, resulting in the inhibition of the long-range cycloid spin and the improvement in the latent magnetization.38,39

Figure 8.

Figure 8

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(a) Magnetic hysteresis loops measured at room temperature with the field applied in the plane for BFO and B0.90H0.10F1–xTxO films (b) tendency of Mr and Ms for these samples.

Conclusions

In summary, pure BFO and Ho and Ti co-substitution B0.90H0.10F1–xTxO (x = 0, 0.025, 0.05, 0.10, 0.15, 0.20) thin films are successfully prepared on the Si(100) substrates via the simple sol–gel technology. The rhombohedral-to-orthorhombic structural phase transition occurs with the increase in Ho and Ti co-substitution concentration. Meanwhile, the obvious agglomerate appears at the surfaces of the Ho and Ti cosubstitution thin films, and correspondingly, the surfaces become denser. Both Ms and Mr, for Ho and Ti co-substitution thin films, are improved obviously according to the structural distortion and the super-exchange interaction enhancement because of the reduced concentration of oxygen vacancies, decreased particle size, and denser surface structure. Our findings suggest that the Ho and Ti cosubstitution is an effective method to improve the magnetic properties of BFO thin films.

Acknowledgments

This work was supported by Natural Science Foundation of China (11904128), the Thirteenth Five-Year Program for Science and Technology of Education Department of Jilin Province (JJKH20200409KJ), the Capital construction funds of Jilin Province (2020C029-5), the 2018 Provincial Talent Development Fund Funded Talent Project of Jilin Province (Research on Doping Modification Design, Preparation and Performance Improvement of Single-phase Multiferroic Materials BFO Film), and the United Laboratory of High Pressure Physics and Earthquake Science.

The authors declare no competing financial interest.

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