Structural and chemical composition of multi-component inorganic nanostructures have stimulated technological and scientific interest to alter the physiochemical properties while developing magnetic, electric, catalysis, and spintronic devices1,2,3,4,5,6. Spinel is one of the complex structures whose physical, magnetic and electrical properties can be altered by adding dopants and using the suitable route for the synthesis7, 8. The spinel ferrites have attracted considerable interest due to their use in microwave technology, magnetic storage, and biomedical applications etc.3, 9, 10. Spinel ferrite has general formula (AII+)[B2 III+]O4 II−, where AII+ and BIII+ are the divalent and trivalent cations occupying tetrahedral (A) and octahedral [B] sites. Face-centered cubic structure of the ferrite is a result of cations and oxygen anions formulation. Divalent cation occupies either tetrahedral or octahedral sites, when it occupies tetrahedral sites, normal spinel is formed. On the other hand, when divalent cation occupies both tetrahedral as well as octahedral sites, inverse spinal is formed11. Similarly, a mixed structure can also be formed when divalent cation is distributed in both sites. The magnetic and electrical properties of ferrites depend on cation distribution and can be altered by varying the place of cation in the interstices. Cobalt and nickel ferrites (CoFe2O4 and NiFe2O4) are intensively studied spinel ferrites due to their high application potential12, 13. CoFe2O4 is the most versatile hard ferrite with mixed cubic spinel structure having Fd3m space group. CoFe2O4 exhibits high coercivity (5400 Oe), high magneto-crystalline anisotropy and moderate saturation magnetization14,15,16. Amongst several multiferroics, bismuth ferrite (BiFeO3) has been reported as one of the versatile cubic perovskites exhibiting both ferroelectricity and G-type anti-ferromagnetism above room temperature17, 18. In multiferroics, parameters such as electric polarization and magnetism generally are responsible for magnetoelectric effect19. The anti-ferromagnetism comes from unpaired electrons in the d shell of the Fe3+ with very weak ferromagnetic ordering due to canted spin structure. The ferroelectricity arises from the displacement of Fe3+ and Bi3+ in the unit cell17. These two effects are very weak in single phase material19, 20. Single phase multifunctional materials are rarely found in nature. Due to special structural features dual phase materials demonstrate different magnetic and electric properties. They demonstrate a strong multiferroic property and consequently have better application potential than single phase21. Thereby, dual phase materials have attracted much interest in research and industrial market. Such materials can be obtained by synthesizing artificial composite of ferromagnetic and ferroelectric materials22. Secondly, magneto-electric or multiferroic are important materials in recent years, where oxygen stoichiometry plays a crucial role in composition of oxides. The nonstoichiometry in oxide determines the phase stability and structural, magnetic and electrical properties of oxide materials23. These oxide materials have considerable demand in spintronic and data storage devices24, 25. In past, various investigators have reported the synthesis of either pure ferroelectric or ferromagnetic properties. To practical potential, it is necessary to develop dual phase composite materials. Synthesis of these materials is associated with the doping level and the choice of dopant.

We reported replacement of trivalent Fe3+ by Bi3+ in our previous work26. In the present paper, we report on the synthesis of dual phase Bi3+-doped CoFe2O4 nanostructures with general formula Co1−xBixFe2O4 (CBF) where x = 0.0–0.2 by sol-gel self-combustion method. The change in the phase from cubic spinel to spinel-perovskite with the substitution of the Bi3+ in place of the Co2+ has thoroughly been investigated. Structural, morphological and magnetic properties of CBF as a function of Bi3+-doping levels are measured and reported. Presence of trivalent Bi3+ instead of divalent Co2+ in the CoFe2O4 crystal produces CBF with different stoichiometries, structures, morphologies and magnetic properties.

Experimental Section


The synthesis of CBF nanostructures was carried out by sol-gel self-combustion method with a range of x = 0.0 to 0.2. The high purity analytical reagent grade (99.99%) cobalt nitrate (Co(NO3)2 6H2O), bismuth nitrate (Bi(NO3)3 5H2O), ferric nitrate (Fe(NO3)3 9H2O) and citric acid (C6H8O7 H2O) (sd-fine, India) chemicals were used as starting materials. All reagents were weighted in molar proportions; the products of the system were synthesized by keeping constant 1:3 metal nitrate to citrate ratio. The ferric nitrate, cobalt nitrate and citric acid were initially dissolved in de-ionized water and bismuth nitrate was dissolved in concentrated HCL to get clear and agglomeration-free solution. An aqueous solution of citric acid was mixed with metal nitrate as chelating agent27 and the pH of solution was increased to 7 by an addition of ammonia solution28. The solution was kept on hot-plate with continuous stirring at 90 °C. Due to evaporation process, the solution was turned to viscous and finally, a viscous gel was obtained. On removal of complete water molecules from the mixture, the gel was automatically ignited and burnt with glowing flints. The decomposition reaction would not stop until the whole citrate complex was consumed. The auto-ignition was complete within a minute, yielding the ashes termed as precursor product with some impurities, collected as sediment which was at the bottom of conical flask. The as-prepared powders (nanostructures) of all samples were heated separately at 500 °C for 5 h and further characterized.


For the investigation of formation of the dual phase CBF composites, X-ray powder diffraction (XRD) patterns were recorded on Rigaku–denki (Japan) X-ray diffractrometer (D/MAX2500) using Cu-Kα radiation (λ = 1.5418 Å) in the 2θ range from 15 to 80° with scanning rate 10°/min. For examining the cross-out morphologies of samples and elemental composition percentage, involved in CBF, scanning electron microscope (SEM) digital photoimages and energy dispersive X-ray spectra (EDS) were used. The magnetic data for these samples were obtained with a vibrating sample magnetometer (VSM) at room temperature by Lake Shore: Model: 7404. The Mossbauer spectra were taken in transmission geometry at room temperature for which a 57Co/Rh γ-ray source was used. The velocity scale was calibrated relative to 57Fe in Rh. For the qualitative evaluation of the Mossbauer spectra recoil spectra were analyzed using WinNormosFIT software29.

Results and Discussion

Structural verification

The crystal structure and phase transition of the samples were confirmed from the XRD patterns. Figure 1 depicts the XRD spectra of CBF ferrite samples obtained for various x values i.e. 0.00, 0.05, 0.10, 0.15 and 0.2. The diffraction patterns and the relative intensity of all diffraction peaks match well to those of JCPDS card number 22–1086, supporting for the formation of CoFe2O4 30 phase type. In XRD patterns of x = 0.00 and 0.05 samples, diffraction peaks of other phases are not evidenced, which confirms the formation of single phase cubic spinel structure with space group Fd3m. This suggests that the doping of Bi3+ enters the interstices of cubic structure. For x = 0.05, no significant change in the phase of CBF structure has been detected. In accordance with JCPDS card no. 20–0169 with space group-R3c31, occurrence of an additional (101) and (110) reflection planes for x ≥ 0.10 Bi3+-doping level corroborates the existence of BiFeO3. It is to be noted that up to 0.05 doping level of Bi3+, the pure cubic spinel phase of CoFe2O4 is dominating and when Bi3+ doping level is ≥0.10, dual i.e. cubic and perovskite phase structures are evolved. The trace amount of Bi3+ (1.03 Å) can be embedded into cubic lattice, while the remaining Bi3+ could form the perovskite phase. In cubic spinel structure, (220) and (400) planes are sensitive to cation distribution on tetrahedral and octahedral sites, respectively. The cations determine magnetic moment of the ferrite32, 33. Inset of Fig. 1 shows an enlarged image of (311) peak where the peak positions are shifted to a higher angle side with substitution of Bi3+, suggesting increase of the distortion in the lattice of pure cubic structure34. This distortion is may be due to Bi3+-substitution which occupies interstices of the ferrite lattice up to x = 0.15. For x = 0.2, the increased amount of Bi3+-substitution may create a new phase of perovskite which appears simultaneously with the cubic phase of ferrite. Thus, the trace amount Bi3+ may enter into the interstices of the cubic lattice and form perovskite phase simultaneously with the spinel phase. The shifting of (311) peak position by 0.13° supports the formation of CoFe2O4 in different stoichiometry as a function of the Bi3+-substitution. The amount of Bi3+, in CoFe2O4, has been determined by the EDS analysis. The molar ratio of Co2+, Bi3+ and Fe3+ cations are given in Table 1. The Rietveld refinement of structure by using XRD data was processed by using Fullprof suite software. Figure 2 displays the Rietveld refined XRD patterns of CBF where it is clearly evident that, the refined pattern has exhibited two phases; first: the cubic spinel structure with space group Fd3m and the second: the perovskite structure with space group R3c. Figure 3 shows the variation of the spinel and perovskite phases of the CBF as a function of Bi3+ -doping level. The cubic spinel phase of CBF is decreased and the perovskite phase is increased with increase of Bi3+ -doping concentration. The unit cells of cubic and perovskite phases are shown on the sides of the graph in Fig. 3. The structure towards left side of the graph shows spinel phase of the cobalt ferrite in which Fe3+ is shown to occupy both tetrahedral as well as octahedral sites while Co2+ occupies only octahedral sites. The Bi3+ cations occupy tetrahedral sites of the perovskite phase (right side of the graph) in which Fe3+ cation occupies both tetrahedral and octahedral sites. The phase analysis and structural parameters such as lattice constant obtained from Rietveld refinement of CBF are outlined in Table 1. The variation of lattice constant “a” for spinel phase and “a = b”, “c” for perovskite phase is given in Table 1. From the Table 1, it is clear that lattice constants for both the phases have not been changed with increasing Bi3+ -substitution. In this case, we have replaced Co2+ (0.78 Å) by larger Bi3+ (1.03 Å). The substitution of larger radii cations generally increases the lattice parameter31. However, in this particular case the percentage of substitution of Bi3+ is very low. Owing to this, the lattice constant(s), for both the phases remain almost constant. The quality of the refinement was verified by corresponding figure of merit, discrepancy factor (Rwp), expected values (Rexp), and goodness fit factor (χ2).

Figure 1
figure 1

Powder XRD patterns of CBF for x = 0.0, 0.05, 0.10, 0.15 and 0.20. Inset is the (311) peak positions of CoFe2O4 for various Bi3+-doping levels i.e. various x values.

Table 1 Molar ratio, phase analysis, lattice constant and Rietveld refinement parameter i.e. Rwp, Rexp, and χ2 of CBF samples.
Figure 2
figure 2

Rietveld refinement of CBF for x = 0.2 samples (dotted black lines are from Fd3m space group spinel phase and blue lines are due to R3c space group of perovskite phase).

Figure 3
figure 3

Influence of Bi3+-doping on spinel (left) and perovskite (right) phases of CBF.

Morphological changes and elemental mapping studies

The micro-structural evolution of cobalt ferrite powders at room temperature under various Bi3+-doping levels i.e. at x = 0.0, 0.05, 0.1, 0.15, and 0.2 were studied by SEM digital images. Figure 4(a) shows the image of pristine CoFe2O4 i.e. without Bi+-doping. The formation of soft spherical crystallites where majority crystallites are aligned in the same direction by forming a dense structure is vivid from the figure. All crystallites exhibit nearly same size. Figure 4(b) shows image of CBF with 0.05 mol Bi3+-doping level where crystallites are smaller in sizes, uniformly distributed and free from the voids or pores. Figure 4(c) shows the image with doping of 0.1 mol of Bi3+, where crystallites in the form of plates and rods of bigger sizes are observed. Figure 4(d) and (e) present the surface images of CBF samples for 0.15 and 0.2 mol of Bi3+-doping levels, respectively. In previous image, regular crystallites reveal majorly rod-type surface whereas in later case, agglomerated crystallites of different shapes and sizes are evidenced. Figure 4(f) gives EDS spectrum of CBF sample obtained when x = 0.20 where the presence of O, Fe, Co, and Bi as major elements is confirmed, suggesting oxygen and Bi are successfully substituted in the crystal structure. The estimated amounts of (in atomic percentage) O, Fe, Co and Bi in the pure and CBF samples are confirmed from EDS spectra analysis which are tabulated in Table 2. The concentration of Fe has remained nearly same in all powders and the concentration of Co is decreased as the Bi3+-doping level is increased. The CBF samples obtained at x = 0.15 and 0.20 exhibit different stoichiometry.

Figure 4
figure 4

The SEM images of CBF with; (a) x = 0.0, (b) x = 0.05, (c) x = 0.10, (d) x = 0.15, (e) x = 0.20 and (g) EDS (when x = 0.20).

Table 2 The stoichiometry (%) of constituent elements present in CBF powders.

Mossbauer spectra analysis

Due to fine energy resolution, Mossbauer spectroscopy can be used to detect even a minute change in the nuclear realm of the iron atoms. In Mossbauer spectroscopy, γ-rays are emitted or absorbed by the crystal without energy loss. So Mossbauer spectroscopy is convenient tool to determine the cation distribution, spin magnetic moment and hyperfine interaction in spinel ferrites35. Of all CBF synthesized samples, three (x = 0.0, 0.1 and 0.2) were characterized for Mossbauer analysis at room temperature (Fig. 5). Each spectrum exhibits Zeeman pattern shape with two sub-spectra; one corresponding to Fe ions in tetrahedral A-site and other to Fe ions in octahedral B-site. The hyperfine interactions like isomer shift (IS), quadrupole splitting (QS), magnetic hyperfine field (Hf), relative area percentage (A), and cation distribution were determined from the analysis of the spectra for three samples and are presented in Table 3. It is observed that IS at A-site is increased and B-site is decreased (by a very small amount) with increasing Bi3+ -doping concentration which can be explained through the bonding ability of Fe with Co and Bi at both sites. With increasing Bi3+ -doping level, the occupancy of Fe3+ at A-site is decreased while at B-site it is increased36. The ions on B-sites are Fe3+ and Co2+ (0.67 and 0.78 Å) with smaller ionic radius sizes and ions on A-sites are Bi3+ and Fe3+ (1.03 and 0.67 Å) with larger ionic radius sizes. Due to which there is an expected increase in the orbital’s overlapping of the ions in the A-site and a decrease in the B-sites, resulting in IS changing. The IS value is increased in the A-site because of the replacement of Fe3+ with Bi3+. The super-transferred hyperfine field components are expected to be strongly influenced by the super-exchange coupling with neighboring ions and the magnetic moments of these ions37. In spinel structure, metal ions on octahedral B-site generally have strong super-exchange interactions with six neighboring metal ions on tetrahedral A-sites, while ions on the A-sites can be coupled strongly with twelve neighbouring metal ions on B-sites through the oxygen ion. The B-site ion interacts with only one A-site ion while an A-sites ion interacts with three nearest neighbouring B-sites ions. Because of large separation distance (~3.5 Å), A-sites ions are not expected to have detectable interaction with other A-sites ions (A-A interaction). The B-sites ions may interact with neighboring B-sites ions only by direct overlap. The distance between neighboring atoms lead to weak interactions (B-B interaction). The Hf at octahedral B-sites of inverse spinel ferrite is generally 10% greater than that of A-sites and this difference is usually due to covalancy38. Table 3 shows the variation of Hf at A and B-sites (HA and HB) with increasing the Bi3+ doping concentration at room temperature. The hyperfine field at both crystallographic sites for samples is decreased when x increased from 0.0 to 0.20 which can be explained with the help of mechanism of supertransferred hyperfine field components. This mechanism is strongly influenced by the super-exchange coupling with neighbouring ions and magnetic moments of these ions39. In the present samples of CBF, positions of Fe3+ at tetrahedral A-sites are replaced by the Bi3+ and super-exchange Bi(A)–Fe(B) between the ions occurs. Huang et al.40 have reported supertransferred Hf mechanism wherein Hf value of a Fe3+ coupled anti-ferromagnetically with another through superexchange path of 180° has been increased. Thus, it is anticipated that the replacement of A-sites Fe3+ with nonmagnetic Bi3+ can reduce the hyperfine field at a neighboring B-sites Fe3+. Also replacement of B-sites Fe3+ and Co2+ with nonmagnetic Bi3+ can be responsible for hyperfine field to decrease on A-sites. The QS values of CBF are given in Table 3 where the quadrupole splitting for system has shown no variation, indicating Fe3+, Co2+ and Bi3+ symmetry have not been changed between Fe3+ and their surrounding with addition of Bi3+ in the system. The Mössbauer effect technique was used to investigate these materials. The Mössbauer spectra were used to determine the sites occupancy in the spinel, which are ordered magnetically. The major magnetic interaction is A-B between A-sites and B-sites cations, the A-A and B-B interactions being much weaker. The magnetic field of Fe3+ cation depends on the nearest neighbouring cation environment41, particularly when it occupies the B-sites. In CoFe2O4, a broadening of the hyperfine lines from the B-sites due to variations in cation distribution at A-sites is noticed. The relative numbers of Fe3+ in A and B-sites were determined from intensity ratios of the outer peaks. From this, the numbers of Co2+ and Bi3+ on A and B-sites can be determined. The simplest case can be CoFe2O4 where there is only one Co2+ cation with Fe3+. From the Mössbauer spectra of CoFe2O4, the ratio of area of the outer A and B sub-lattice peak is indication of an excess of Fe3+ on B-sites and hence fraction of Co2+ on A-sites. This distribution can be described by the (Co0.034Fe0.966) [Co0.966Fe1.03] O4 chemical formula. From Table 3, populations of Fe3+ on A-sites and B-sites were estimated. For sample for x = 0.0, 0.1 and 0.2, the number of Fe3+ on A and B-sites of CBF are 48.30% and 51.69%, 46.28% and 53.71%, 47.57% and 52.42%, respectively. For pristine sample i.e. for sample with x = 0.0, 3.4% of Co2+ occupy tetrahedral A-sites. When Bi3+ is substituted, the occupied percentage of Bi3+ at A-sites is increased to 7.5% and 4.8% for x = 0.1 and x = 0.2 samples, respectively. With an addition of Bi3+ (i.e. x = 0.1), the intensity of outermost A-sites peaks (Fig. 5) is increased by 5.4% with respect to pristine sample (i.e. x = 0.0) and at x = 0.2 intensity of A peak is decreased by 0.73% (with respect to x = 0.0). This trend is consistent with the picture that Bi3+ can enter into tetrahedral sites, Co2+ enter into the octahedral sites and Fe3+ at tetrahedral sites as well as at octahedral sites.

Figure 5
figure 5

Mossbauer spectra of CBF for x = 0.0, 0.1 and 0.2.

Table 3 The IS, QS, Hf and A values, obtained from Mossbauer analysis, of CBF for various Bi3+ i.e. x values.

Magnetic properties

Magnetic hysteresis loops were recorded at room temperature. Magnetic hysteresis loops of all samples annealed at 500 °C are shown in Fig. 6. The saturation magnetization (Ms), coercivity (Hc), remanent magnetization (Mr) and remanent ratio (R) for all composition of samples are listed in Table 4. It is clear from Fig. 6 that, coercivity and remanent magnetization are increased with Bi3+-doping level. The variation of saturation magnetization with Bi3+-substitution is shown in Table 4, where increasing (up to x ≤ 0.15) and decreasing (x = 0.2) trends are evidenced. Due to addition of non-magnetic Bi3+ ions Fe3+ ions from A-sites transferred to B-sites due to which the magnetic moment of A-sites decreases. The net magnetization, being the difference between B and A-sites magnetizations, is increased due to small increase of Fe3+ on B-sites. The magnetic moment is supposed to increase with Bi3+ content; this could be explained on the basis of magnetic moment of constituent ions. On addition of non-magnetic Bi3+, concentration of Fe3+ in the A-sites is decreased, as a result magnetic moment of the sites reduces. On B-sites, concentration of Fe3+ is increased42, 43. Hence on introduction of non-magnetic Bi3+ the net magnetic moment up to x = 0.15 is increased. The magnetic moments are dropped for higher values which could be explained on the basis of spin canting. Spin canting is the effect in which non-magnetic substitution on one sub-lattice could lead to a non-collinear or canted spin arrangement on other sub-lattice44. As Bi3+ (non-magnetic) -content is increased after certain level (x = 0.15), the exchange interactions weaken and the spin magnetic moment of B sub-lattice will no longer be parallel to the spin magnetic moment of A sub-lattice. The decrease in the B sub-lattice moment can be interpreted as a spin departure from co-linearity which causes the effect known as canting45. Geller46 gave the canting approach in which individual moments on one sub-lattice are canted at different angles. Now out of the two sub-lattices i.e. B and B́, only B́ may have affected by the canting effect. It is presumed that the B́ sub-lattice is formed by the cations of B-sites those are in the neighbourhood of A-sites which contains Bi3+. With increasing concentration of Bi3+, the canting effect is increased and the spin magnetic moments of B-sites are canted from the direction of net magnetization. The coercivity is increased from 1457 to 2277 Oe with increasing Bi3+ which may lead to the fact that Hc can be enhanced by enlarging the magnetocrystalline anisotropy. For x = 0.2, number of Co2+ is decreased due to increase of Bi3+. The Bi3+, accommodated at the rhombohedral perovskite lattice sites unable to enter cubic lattice, produces structural distortion in cubic structure, resulting in decrease of coercivity. The remanent ratio, R = Mr/Ms is characteristic parameter of the material. High remanent ratio is desirable for magnetic recording and memory devices47, 48. It is an indication of the ease with which the directions of magnetization reorient to nearest easy axis magnetization direction after the magnetic field is removed. Lower value of the remanent ratio is an indication of the isotropic nature of the material. The values of R in the present case are varied from 0.54 to 0.57, showing no significant change in the value with increasing substitution of Bi3+.

Figure 6
figure 6

Magnetic hysteresis loops measured for CBF at room temperature for different ‘x’ values.

Table 4 The Ms, Hc, Mr and R values, magnetic hysteresis measurements, of CBF for different ‘x’ values.


We have demonstrated controlled synthesis of Bi3+-doped cobalt ferrite having dual phase (spinel and perovskite) structures, where spinel phase is diminished and perovskite phase is evolved with increase of Bi3+-content. The cubic spinel phase is evidenced up to 0.15 Bi3+-doping level and for x = 0.2, the perovskite phase is dominating showing impact on structural and magnetic properties of the crystal. The doping of Bi3+ has made remarkable and interesting changes in cation distribution, where Bi3+ occupy tetrahedral sites thereby replacing Fe3+ cations to octahedral sites. This is confirmed from Mossbauer spectra analysis. Saturation magnetization, corecivity and remanence magnetization are increased with increasing doping level of Bi3+ and are maximum at x = 0.15. For further increase in doping level to 0.2 of Bi3+ discussed magnetic properties are decreased, revealing dominancy of perovskite phase.