Magnetic and Magnetoelectric Properties of AurivilliusThree- and Four-Layered Intergrowth Ceramics

Abstract

In this work, we have prepared intergrowth of multiferroic compounds namely Bi₄RTi₃Fe_0.7Co_0.3O₁₅-Bi₃RTi₂Fe_0.7Co_0.3O_12−δ (BRTFCO₁₅-BRTFCO₁₂) (rare earth (R) = Dy, Sm, La) by solid-state reaction method. From the X-ray diffraction Rietveld refinement, the structure of the intergrowths was found to be orthorhombic in which satisfactory fittings establish the existence of three-layered (space group: b 2 c b) and four-layered compounds (space group: A2₁am). Analysis of magnetic measurements confirmed a larger magnetization for theSm-modified intergrowth compound (BSTFCO₁₅-BSTFCO₁₂) compared to Dy- and La-doped ones. The emergence of higher magnetic properties can be due to distortion in the unit cell when some Bi³⁺ ions are replaced with the Sm³⁺, bonding of Fe³⁺-O-Co³⁺ as well as a possible mixture of Fe_xCo_y-type nanoparticles that are formed generally in the synthesis of intergrowths. The changes in the magnetic state of the Aurivillius intergrowths have been reflected in the magnetoelectric (ME) coupling: higher ME coefficient (~30 mV/Cm-Oe) at lower magnetic fields and is constant up to 3 kOe. The results were corroborated by Raman spectroscopy and variation of temperature with magnetization data. The results revealed that the RE-modified intergrowth route is an effective preparative method for higher-layer Aurivillius multiferroic ceramics.

1. Introduction

Aurivillius oxides with Bismuth layer structure and the chemical formula: (A_n−1B_nO_3n+1)²⁻ (Bi₂O₂)²⁺ are known to possess excellent magnetoelectric properties and stable ferroelectric polarization [1,2]. Here A- stands for mono-, di-, or trivalent cation or a mixture of them at a 12-coordinate site. The term B- represents tri-, tetra-, or pentavalent cations at the 6-coordinate site. The term ‘n’ indicates the number of layer compounds or perovskite units [1]. The structure of the compounds consists of alternative perovskite blocks of (A_m−1B_mO_3m+1)²⁻ and are interleaved with the bismuth oxide layers (Bi₂O₂)²⁺ along with the pseudo-tetragonal c-axis. The said above compounds are found to be interesting from the scientific and technological point of view like in piezoelectric devices, memory devices, pyro sensors, and magnetoelectric sensors [2,3,4,5]. Bi₄Ti₃O₁₂ compound is a three-layered compound having a high transition temperature (Tc~675 °C) and high electrical leakage currents. To reduce its limitation, rare earth elements were substituted for Bi³⁺ions for fine-tuning the physical properties. In addition, Sr²⁺ modified bismuth-based four-layered compounds were found to be phase-stable and the Sr-O bond suppresses the bismuth evaporation during the sintering stage. The intergrowth compounds of Sr modified 3-layered and five-layered were found to be interesting owing to the fact of their high dielectric properties [6,7]. Alternatively, many Aurivillius phase compounds have shown multiferroic properties where B-site is modified with Ti/Fe/Co ions [8,9,10,11,12]. In other words, by doping rare-earth ions for A-site and magnetic ions in B- sites were found to have room temperature (RT) multiferroic applications and therefore these materials are referred to as future bismuth layered structure ferroelectromagnetic (BLSFEM) materials [8,9,10,11,12,13,14,15,16].

In view of the aforementioned importance of BLSFEM, many researchers were concentrating, globally, to improve electrical and magnetic properties by adopting the following strategy:

• Doping A-site with rare earth (RE) ions.
• Magnetic ion (Co/Fe) doping at B (Ti)-site.
• Adopting the intergrowth method between two different layer compounds.

The intergrowth of the BLSFEM compound consists of one unit cell of one-half the unit cell of m-layered and the other half unit cell of (m + 1) layered compounds along the c-axis [6,17]. Earlier we reported the electrical properties of intergrowth of Sm-modified Bi₄Ti₃O₁₂ and SrBi₄Ti4O₁₅ [17]. The XRD pattern of intergrowth is well indexed with the SrBi₈Ti₇O₂₇ (eight layers) compound. The same results were reported on similar compounds [18,19]. Kikuchi et al. [20] reported that the regular structure for higher-layer compounds m ≥ 5, is built up of regular intergrowth of two layers of compounds of m and (m + 1) in the common direction of the c-axis. They coined these compounds as a mixed- layered compounds and the structure can be written as:

(Bi₂O₂)²⁺(A_n−1B_nO_3n+1)²⁻ − (Bi₂O₂)²⁺(A_m−1B_mO_3m+1)²⁻ or Bi₄A_m+n−2B_m+nO_3(m+n)+6

Here, the terms m and n represent the perovskite layers. It should remember here that the sum of ‘m’ and ‘n’ should be an even layer compound. Based on this, one can presume that SrBi₈Ti₇O₂₇ is built up with a mixed layer or intergrowth of m = 3 and m = 4 [20]. However, the plausible reason for the intergrowth structure of high-layer compounds is not known completely. On account of the complexity of the structure, a large number of layer compounds (m ≥ 5) are being prepared by adopting the intergrowth route [21,22,23]. It was also reported that the ferroelectric properties of intergrowth compounds have shown higher ferroelectric properties compared to that of reactant layer compounds namely 3- and 4- layers compounds [7]. A recent report by Zhang et al. [8] revealed that multiferroic intergrowth formed out of Bi₃NdTi₂Fe_1-xCo_xO_12-δ-Bi₄NdTi₃Fe_1-xCo_xO₁₅ (x = 0.1–0.7) have shown higher magnetic and magnetoelectric (ME) properties for x = 0.3. From transmission electron microscopy, it was reported modulation of 3 and 4-layered stacking sequences can be altered based on the cobalt content (x); alternative 4 and 3 for x = 0.5, a dominant 3-layer for x = 0.7 [8]. They found that the Co-ion substitution in the phase of Fe (for B-site) has shown higher ferromagnetic and magnetoelectric properties [8]. In view of this, many researchers [21,22,23,24,25,26,27,28,29,30,31] are preparing Co/Fe-modified (at B-site) compounds, and therefore intergrowth route is considered to be an effective way to improve ferroic properties, especially in Aurivillius materials.

In view of the significant ME coupling of the intergrowths based on the cobalt content, we chose the composition of Bi₃NdTi₂Fe_0.7Co_0.3O_12-δ-Bi₄NdTi₃Fe_0.7Co_0.3O₁₅ as a parent compound. Since rare-earth replacement at the Bismuth site in Aurivillius is known to reduce oxygen vacancies, [5,32,33] different rare earth ions, namely Dy, La, and Sm were doped at the A–site for the improved magnetoelectric coefficient. Detailed phase quantification, magnetization, and ME coupling are presented for different compositions.

2. Materials and Methods

Aurivillius phase of intergrowth compounds Bi₃RTi₂Fe_0.7Co_0.3O_12-δ-Bi₄RTi₃Fe_0.7Co_0.3O₁₅ (rare earth; R = Dy, Sm, and La) were prepared by conventional solid-state reaction method. Initially, high-purity reactant oxides of Bi₂O₃, Dy₂O₃, Sm₂O₃, La₂O₃, TiO₂, Fe₂O₃, and Co₃O₄ dried at 100 °C for 6h, were used to prepare 3-layered (Bi₃RTi₂Fe_0.7Co_0.3O_12-δ) and 4-layered (Bi₄RTi₃Fe_0.7Co_0.3O₁₅) compounds. The mixed powders were calcined at 850 °C for 4 h and subsequently grounded for 12 h in the ball mill. The resultant compounds were analyzed for phase conformation by XRD measurements. Prior to the intergrowth formation, the prepared 3- and 4- layered compounds were taken in stoichiometric ratio and milled for 12 h in the planetary ball mill. The present intergrowth compounds Bi₃RTi₂Fe_0.7Co_0.3O_12-δ-Bi₄RTi₃Fe_0.7Co_0.3O₁₅ (rare earth, R = Dy, Sm, and La) are named BDTFC, BSTFC, and BLTFC respectively. All the compounds were characterized by Powder X-ray diffraction using Cu-Kα₁ radiation. The sintered powders were made into pellets of diameter, and thickness around 10 mm, and 1 mm respectively, and employed final- sintering at 900 °C for 4 h.

Powder X-ray diffraction was employed for the conformation of phase using a PANalytic X’Pert³ diffractometer. The XRD data was refined using FullProf Suite software. The surface morphology and EDAX (Energy Dispersive X-ray Spectroscopy) of the samples were studied using the ZEISS-EVO185SEM instrument. The grain sizes of the prepared samples were analyzed using ImageJ software and employed a histogram with Gaussian distribution. The structural evaluation was further studiedby means of Raman spectroscopy using a Jobin-Yvon spectrometer (Horiba) using a 532 nm excitation laser. The magnetic measurements were carried out on all samples using under varying external magnetic fields using a vibrating sample magnetometer (VSM, cryogenic, UK). ME studies were measured using, Micro-Measurement Group Strain Indicator, Model 3800, and series WK strain gauge. The ferroelectric polarization–electric field (P-E) wasstudied under different driving voltages using aP-E loop tracer (Radiant technologies). All samples were electrically poled under 10 kV/cm and magnetically at 5 kOe for 30 min each, prior to ferroelectric and magnetoelectric studies.

3. Results

3.1. Powder X-ray Diffraction and Morphological Studies

To confirm the single-phase purity, XRD was carried out in the range of 20°–80° with a scan speed of 0.1°/min. The XRD patterns of both 3- layered compounds (Bi₃RTi₂Fe_0.7Co_0.3O_12-δ, R = Dy, La, and Sm) and 4- layered compounds (Bi₄RTi₃Fe_0.7Co_0.3O₁₅, R = Dy, La, and Sm) were used in Rietveld refinement. Figure 1a–c shows the XRD refinement of BDTFC, BLTFC, and BSTFC compounds. The lattice parameters were calculated by giving both space groups of 3-layered and 4-layered compounds. The compounds were found to be orthorhombic phase groups of A2₁am and b2cb with mixed phases of both 3- and 4-layered compounds. A detailed Rietveld analysis is made on these samples and the parameters were given in Figure 1. The Bragg peaks were fitted very well with the expected error with varying the size of rare earth ions [27,28]. The variation of lattice parameters and volume fraction of 3- and 4- layered phases with rare earth size is shown in Figure 2a–c. Parameters for 3-layer and 4-layer with respect to the rare earth ion incorporation for Bi-site are shown in Figure 2. The volume fraction for the 4-layer phase is found to decrease, whereas, for the 3-layer phase it increases with an increase in ionic size.

From the plot, the overall phase contribution is found to be 4-layered (see Figure 2b), and the results were consistent with the reported HAADF images [8].

A good matching between theoretical and experimental XRD peaks was demonstrated by showing reliability factors within the reliable range (>10) [8]. Based on this observation one can conclude that the presence of pyrochlore formation is mainly due to the mixed valence combination of Fe and Co ions. It should remember here that the Ti, Fe, and Co cations were occupied at the B-site (center of each layer- perovskite) with fixed occupancy.

It is a known fact that Ti⁴⁺, Co²⁺, and Fe³⁺ cations are distributed randomly at the center position of octahedral for both 3- and 4-layered compounds. A partial substitution of rare earth (R³⁺) ion with Bi³⁺ certainly affects the structural changes. To confirm the quantification of the second phase, detailed EDAX images were shown in Figure 3.

From the images, it is evident that spots were observed. The orientation of randomly oriented plate-like morphology is also evident in the SEM images. In addition, a small spot (shown as encircle) of grain-like growth is also seen in the images. This kind of detected morphology is attributed to the accumulation of polyhedral phases of impurity phase or Co-Fe segregation at the grain boundary region [8]. The appearance of the secondary- phase affects the grain growth and leads to plate-like morphology with increasing the RE incorporation. The ratio of Fe/Co content found in the elemental analysis, shown in Figure 3d–f, is found to be less than unity. This conforms to the collection of polyhedral grain over the SEM morphology, as shown as a circled region. Based on this observation one can conclude that the presence of pyrochlore formation is mainly due to the mixed value combination of Fe/Co and its ratio.

3.2. Raman Spectroscopic Studies

To understand the local structural properties, Raman spectroscopic analysis is made in the present investigation. Figure 4a–c shows the Raman spectra of BDTFC, BSTFC, and BLTFC samples respectively. Here, green lines indicate deconvoluted curves of the Raman peaks. The spectral phonon peaks above 200 cm⁻¹ marked as ν₄ are assigned to the internal mode of TiO₆ octahedra. A small shift in the fitted peak position indicates that RE⁺³ ion participation in place of Bi⁺³ ion. The broad valley-like behavior observed between 300 cm⁻¹ and 400 cm⁻¹ indicates the diffusion of Co/Fe content. Appearing two peaks, near 550 cm⁻¹, is mainly due to the substitution of Fe/Co at the Ti-site. It is difficult to explain each phonon mode and its overlapping with weak intensity. The detailed phonon modes indicated in Figure 4 are almost well-matched with the literature [29,30,31].

The modes observed at 260 cm⁻¹ (ν₃) are attributed to the internal angles torsional bending of (Ti/Fe/Co)O₆ bonding. The modes observed at 360 cm⁻¹ (ν₆) are due to the stretching vibration of oxygen along the c-axis (O-Ti/Fe/Co-O) stretching mode. The mode ν₇ (470 cm⁻¹) is due to changes that occurred in the bond distance Bi atom within the Bi₂O₂ layer and its apical O-atom of the perovskite block. Two nearby modes ν₈ and ν₉ observed at 540 cm⁻¹ and 560 cm⁻¹ are due to the physically opposing oxygen atom from TiO₆ octahedral [29,32,33,34]. The stretching modes of Bi/RE-O stand in A-site shows modes at ν₁₀ and ν₁₁. The last modes ν₁₂ and ν₁₃ are related to once again bonding of stretching of octahedral. The overall Raman peaks and a slight notable difference observed near ν₁-ν₃ (low region) are attributed to A-site cation and Bi/RE⁺³ ion variation in the (Bi₂O₂)²⁺ layers. The results also suggest the rare earth ion readily substitutes within the Bi⁺³ ion [26]. The peaks near ν₈ and ν₉ are attributed to the decrease in the fitted octahedral distortion in (Ti/Fe/Co) O₆. The broad mode originating at 560 cm⁻¹ explains the substitution of RE ions at the A-site. An increase in the broadness of peak or more tendency of convolution is observed with decreasing the rare earth ion substitution for the Bi (A-) site. An increase in the intensity of the ν₉ peak indicates more structural distortion and improves the competitive electric-magnetic interactions.

3.3. Ferroelectric Studies

The Hysteresis loop (Polarization (P)-Electric field (E)) of BDTFC, BSTFC, and BLTFC samples are shown in Figure 5a–c. The measurements were performed at different voltages as depicted in the plot. The shapes of the hysteresis loops were found to be sharp at the edges. It is a known fact that the coercive and saturation fields of Aurivillius materials are very high. The piezoelectric properties (d₃₃) of BLSFEM are generally low (~50 pC/N). To improve d₃₃ and ferroelectric properties different rare earth ions were substituted for the Bi³⁺ site. If the Bi³⁺/RE³⁺ ratio increases then the deformation of the octahedron of layer-perovskite alters the dipole moment as well as depolarization properties. Finally, all the ordering of these materials tunes to multiferroic applications. The unsaturated and non-symmetric nature of loops is associated with the leaky nature, on account of the high concentration of oxygen vacancies [7,17]. It is also observed that the coercive field and polarization increase with the applied voltages, as shown in Figure 5a–c. In addition, when Ti ions get replaced with Co/Fe ions, they form defect complexes such as Fe³⁺-Vo-Fe³⁺ and Co²⁺-Vo-Co²⁺. Thus, the oxygen vacancies (Vo) get trapped by Fe³⁺ and Co²⁺ ions and hence they require ahigher electric field to exhibit saturated hysteresis loops [35]. Unfortunately, it was not possible to apply ahigh electric field on account of the limitation of the instrument. Despite these limitations, all samples have shown ferroelectric loops without a signature of a loss nature.

3.4. Magnetic Studies

It is a known fact that the electric and magnetic interaction originates from the exchange interaction between neighbouring Fe³⁺-O-Co⁺² ions. Room temperature magnetic hysteresis curves are shown in Figure 6a–c. It is reported in the literature that these materials belong to a canted AFM (antiferromagnetism) nature. On account of the canted nature, these materials have small net magnetization with an at-most saturated region at higher magnetic fields [34,35,36]. The antiferromagnetic nature is attributed on account of tilt-canted spins, and can be understood by the following equation:

$$\sigma ={\sigma }_P+\chi H$$

(1)

Here σ is total magnetic moment, σ_P is weak ferromagnetic moment and χH represents antiferromagnetic component. The canted AFM nature is generally verified with the help of the following facts:

• At low temperatures, the M-H loop resembles with ferromagnetic (FM) hysteresis loop.
• At higher magnetic fields, the magnetization shows linear behavior with the magnetic field.
• At higher magnetic fields, magnetization vs. magnetic field data shows a linear increase due to nearby magnetic ion interaction, namely iron (Fe) and cobalt (Co).
• The strength of magnetic nature can be understood by the law approach to saturation. This gives additional information about anisotropic properties.

The magnetization values were found to be high when compared to the parent compound [8]. The high magnetization values were attributed to the structural distortion and domain orientations. The canted AFM or weak ferromagnetic nature is attributed to the canting spin of Fe-Co-Fe interaction or Dzyaloshinskii–Moriya (DM) interaction. Based on the magnetic nature of RE ions, a higher value of magnetization is observed for BSTFC. The results were consistent with the literature [37].

In the present investigation showing saturated behaviour at higher fields clearly indicates a spin canting and the movement of cations between given bonding (Fe/Co) plays a role to give net saturation. The saturation magnetization is generally employed by fitting the isotherm data shown in Figure 7a–c. At lower fields, (H/M) saturation magnetization curves tend to have a finite positive value. The positive intercept indicates the weak ferromagnetic nature [37]. To separate out the FM and AFM contribution in the sample, the following relation is used,

$$M\left(H\right)=\left[2\frac{M_{FM}^S}{\pi }{\tan }^{-1}\left\{\left(\frac{H\pm H_{ci}}{H_{ci}}\right)\tan \left(\frac{\pi M_{FM}^R}{2M_{FM}^S}\right)\right\}\right]+\chi H$$

(2)

The term $M_{FM}^S$ represents FM saturation magnetization, $M_{FM}^R$ represents remanent magnetization. H_ci is the coercivity and the last term χH represents the susceptibility of AFM contribution part.

From Equation (2) the first term represents FM nature and the second term indicates AFM part of strong. Here the strong spin interaction FM nature, and it becomes weak. To extract more information about FM nature and its anisotropy, the law of approach to saturation technique is used [38,39,40,41,42]:

$$M=M_s(1-\frac{a}{H}-\frac{b}{H^2})+kH$$

(3)

Here, (a/H), (b/H²) represents inhomogenuity and magneto-crystalline anisotropic parameters. The $kH$ represernts the field induced force magnetization.

Anisotropy field (H_k) can be estimated from the following formula:

$$H_k=\frac{2K_I}{{\mu }_0{M}_S}$$

(4)

where $K_I$ represents, the first anisotropic constant. Fitting of Equation (3) for sample BDTFC is depicted in Figure 7a. Whereas for samples BSTFC, BLTFC, the last term becomes zero. The fitted equations were given in Figure 7b,c respectively. A detailed aspect of this would be discussed in due course.

The inset of Figure 6 (left) shows the variation of (dM/dH) with the applied field of all samples. Inset Figure 6 (right) shows variation magnetization (M_S) with the applied magnetic field. All dM/dH curves show a single sharp peak at the origin. From this, one can speculate that all prepared samples have a single magnetic grain type and are of soft magnetic nature. It should remember that the Co⁺³ doping with Fe concentration is the same for all the samples (BDTFC, BSTFC, BLTFC), and the two slopes found at the origin explain the magnetic linkage between Fe/Co magnetic grains of layered perovskites. The presence of nanoregions which were raked as Fe, and Co-rich observed as dots in SEM pictures. The calculated ratio of intergrowth grains and magnetic nanoregions of all the samples was found to be less than 0.1.

Low-temperature magnetization by zero-field cooled (ZFC) and field cooled (FC) sequences were measured under a small external magnetic field (1 kOe) in the range of 2 to 300 K, which are shown in Figure 8a–c.The increase in M (magnetization) with decreasing temperature clearly indicates the stability and strong coupling interaction between magnetic dipoles at those temperatures (shown as inset Figure 8a–c). The divergence of ZFC and FC curves is generally attributed to the signature AFM (Antiferromagnetic) nature. In the present case, dispersion is shown in Figure 8 clearly indicating canted AFM nature.

The radius of Co²⁺ (0.745 Å) is larger when compared to Fe²⁺ (0.605 Å). The Co²⁺ substitution for Ti- ion creates oxygen vacancies [8,27]. From this one can speculate Ti³⁺-O-Ti⁴⁺ decreases the bond angle and finally weakens the magnetic exchange interaction. This leads to a decrease in the magnetic order and its interaction. Moreover, a noteworthy aspect observed in the plot of the ZFC and FC condition plot clearly confirms that magnetic weak or superparamagnetic or canted AFM nature. Similar results were also observed in the Bi₉Fe₅Ti₃O₂₇ compound [26].

3.5. Magnetoelectric Studies

Figure 9 shows the magnetic field dependence of the magnetoelectric (ME) coefficient at room temperature. An ac frequency 550 Hz was applied prior to the extraction of ME data. The value of the ME coefficient has been found 27 mV/cm-Oe, 21 mV/cm-Oe, and 14 mV/cm-Oe for BSTFC, BDTFC, and BLTFC respectively. Moreover, the ME coefficient was found to be almost constant throughout the measurement in the positive magnetic field [42]. Both magnetic and magnetoelectric properties of the studied compositions were shown enhanced results compared to the parent compound.

BSTFC compound has shown a higher ME coefficient compared to the other samples. The value is found to be higher than the parent compound (1 mV/Cm-Oe) [8]. From this one can speculate that the intrinsic nature of magnetoelectric interaction could possibly be due to magnetic and electric interaction, driven by inverse DM interaction or spin–orbit and spin exchange interaction. The cause of ME coupling, however, is yet to be explored.

As mentioned earlier, we found nanoregions in the SEM images of the samples. Such nanoregions are also reported in earlier reports corresponding to the Co_xFe_y composition resulting from the synthesis. Zhang et al. [8] also observed the occurrence of a similar phase from their SEM observations which are consistent with the current study. It can be speculated here from the ME coupling effects can be due to the formation of Co_xFe_y phases.

Given the similar conditions of materials synthesis for all the three rare-earth doped samples in this study, one can expect same the amount of Co_xFe_y phase in all these samples. In this scenario, the improved ME response observed for Sm-based intergrowth arises out of structural distortion. The tetragonal strain, orthorhombocity, and orthorhombic distortionsplay a major role in the polarization switching of Aurivillius oxides in view of octahedral distortions [33]. The replacement of Sm, La, Dy at the A-site has resulted in the orthorhombic distortion(b/a), tetragonal strain (c/a), and orthorhombocity (2(a − b)/(a + b)) of (0.9947,0.9974, 0.9943 for 4-layered phase& 0.9993, 0.9993,0.997 for 3-layered phase), (7.6242, 7.606, 7.6484 for 4-layered phase & 5.989, 5.963, 6.001 for 3-layered phase), and (0.0053, 0.0026, 0.0056 for 4-layered phase & 0.0007, 0.0007 0.0033 for 3-layered phase) respectively. Stable orthorhombocity and lower orthorhombic distortions for Sm, Dy-based intergrowths seem to be favourable for polarization switching under magnetic fields.It is also evident that the ME output is strongly dependent on the volume fraction and the existence of both 3-layered and 4-layered phase structures in the present intergrowth compounds. From this, one can conclude that the intergrowth rare earth-doped compounds have shown strong ME coupling in the Aurivillius multiferroic phases. Higher magnetization accompanied by ferroelectric polarization, rare-earth (Sm, Dy) induced structural distortions, intern leads to higher ME coupling.

4. Conclusions

After adding the stoichiometric amounts of the three-layered and four-layered compounds, we have prepared intergrowth compounds named BDTFC, BLTFC, and BSTFC respectively. The intergrowth compounds have been found to be single phase and the structure of the compound is mixed phases of 3- and 4-layered compounds. In our earlier results, and in the literature, it is also reported that the intergrowth of 3- and 4-layered compounds form an 8-layer compound. The reason for the intergrowth structure is still debatable and therefore high-layer compounds will certainly open up a new direction in the multiferroic fields.

Based on the magnetic and magnetoelectric results, one can conclude that the origin of the weak ferromagnetic nature is attributed to the mixed valence states of +2 and +3 of Fe and Co ions. The possible neighbouring exchange interaction between iron and cobalt is Fe⁺³-O-Co⁺², Co⁺³-O-Co⁺², and Co⁺³-O-Fe⁺². In addition to that, Fe⁺³-O-Fe⁺² indicates AFM super-exchange interaction. The canted spin nature, especially in Co and Fe mixed phases is generally observed in these materials, called DM interaction. Doping of RE⁺³ ions causes (Fe/Co)O₆ octahedral tilt and hence DM interaction play a dominant part. Substituting trivalent Co ion in place of Fe ion generally improves strong magnetic interaction. Intergrowth materials were found to be single domain and soft magnetic type. These multi-ferroic materials are important candidates for advanced multi-functional sensors.