PHYSICS AND CHEMISTRY OF SOLID STATE Preparation and Characterization of Nanosized Substituted Perovskite Compounds with Orthorhombic Structure

In this work, five substituted perovskite such as (Gd 0.9 Sr 0.1 ) Mn 0.8 Co 0.2 O 3 , Tb 0.8 Sr 0.2 FeO 3 , Gd 0.6 Sr 0.4 RuO 3 , SrCe 0.95 Y 0.05 O 3 , and Mn 0.6 Co 0.4 SnO 3 were synthesized by tartrate and hydroxide precursor method. The resulting samples were characterized by inductively coupled plasma spectroscopy, energy dispersive X-ray analysis, infrared spectroscopy, thermal analysis, X-ray powder diffraction, transmission electron microscope (TEM), selected field of electron diffraction (SAED), d.c. electrical conductivity, Hall effect, dielectric measurements, and low-temperature magnetization measurements. The X-ray diffraction pattern for all compounds was indicated the formation of single-phase perovskite with orthorhombic structure except Tb 0.8 Sr 0.2 FeO 3 and Mn 0.6 Co 0.4 SnO 3 perovskite. These compounds showed a cubic and rhombohedral structure, respectively. The lattice parameter and the unit cell volume slightly decreased as ionic radii decrease in agreement with the lanthanide contraction. The average size of cation ˂ R A ˃, mismatch factor (σ 2 ), and tolerance factor (t) gives the combined effects of disorder and inhomogeneity in these compounds. The average particle size determined from TEM was in the range of 22 to 77 nm for all compounds. The different lattice planes in the SAED pattern were assigned and these were in agreement with the single crystallographic phase of these perovskites. The temperature dependence of electrical conductivity for all compounds showed a definite break in 500 K to 610 K. except the Gd 0.6 Sr 0.4 RuO 3 compound, which corresponds to semiconducting behavior. While the Gd 0.6 Sr 0.4 RuO 3 sample shows a metallic-like semiconductor. The thermoelectric power and Hall effect measurements for all compounds were n-type semiconductivity except the SrCe 0.95 Y 0.05 O 3 compound. It showed p-type semiconductivity. The frequency dependence of the dielectric constant and dielectric loss in these substituted perovskites were discussed using the Maxwell-Wagner model. Magnetic studies showed that the thermo-magnetic (ZFC and FC measurement) irreversibility for all compounds.


kites, the A-
ite is normally filled by lanthanides, alkali metals, or alkaline earth metals, while the B-site elements are typically first or second-row transition metals [10][11][12].Partial substitution, on the other hand, can occur at the A-site, the B-site, or both the A-and B-sites at the same time (A1-xA'xB1-yB'yO3±δ); The goal of substitution is to improve or fine-tune physical properties such as thermal stability, chemical stability, electrical conductivity, magnetic effects, catalytic activity, and electrodeelectrolyte compatibility in order to increase the commercial utility of perovskite devices.

The magnetic behavior of substituted manganates depends on the ratio of Mn 3+ /Mn 4+ and is regulated by the A-site cation average radius <ra> or tolerance factor, which is a geometric factor indicating the degree of distortion of the cubic perovskite tructure [13].

Maximum ferromagnetic Curie temperature in R1-xA ' xMnO3 is observed when R = La and ferromagnetism is observed for 0.1 < x < 0.5 with maximum Tc for x = 1/3 [14][15][16].When the Mn 3+ /Mn 4+ ratio in substituted manganates is 1:1, true space charge ordering of the Mn ions is observed at low temperatures (x = 0.5).Ferromagnetism and ground-state transition from ferromagnetic to antiferromagnetic transition are suppressed by the charge ordering [17].The Pr0.5Ca0.5Mn1-yCoyO3series offers a weakening of the inclination to charge order by sufficient substitutions on the manganese sites with other transition metal ions.The incidence of charge ordering state is fully suppressed for 0.03 ≤ y ≤ 0.05, and an insulator-metal transition exists in the temperature range from 80 to 60 K [18].

RFeO3 (R = rare earth) has recently become part of the perovskite oxide family with an orthorhombic cell unit.They are insulators with a good anti-ferromagnetic coupling between ions of Fe 3+ .This leads to high TN Neel temperature values.For example, for dFeO3, TN = 657 K [19].SrFeO3, on the other hand, is a TN = 130 K metallic antiferromagnet [20].It reveals a magnetic helical structure with a magnetic moment of 3.1 μB per ion Fe 4+ .Fe 4+ ions in SrFeO3 are normally thought to be in high spin 3d 4 condition with three localized electrons in the t2g band, while the fourth electron is delocalized in a broad σ* band [21][22].

Ruthenate perovskites such as SrRuO3, CaRuO3, and LaRuO3 have a low resistivity metallic property at room temperature [23], while PrRuO3 has a semiconducting property [24].SrRuO3 is ferromagnetic below Tc = 160 K followed by a small shift in Curie temperature re istivity and CaRuO3 follows a Curie Weiss law with a negative Weiss constant at high temperatures and no resistivity deviations down to 4.3 K [23,25,26].Magnetic measurements have been investigated using the solid solution Sr1-xCaxRuO3.With the replacement of Ca 2+ for Sr 2+ , these measurements revealed a decrease in the Curie temperature, suggesting a weakening of the ferromagnetic exchange interaction [27,28].

The compound ACe1-xRxO3-δ (A = Ba and Sr, R = rare earth elements) belongs to perovskite-type oxides, which in addition to good oxygen-ion conductivity, exhibit excellent high proton-conducting efficiency [29][30][31].The transport properties of these materials gen rally depend critically on the structure of crystals and chemical stoichiometry.Oxygen conduction results from the movement of oxygen ions between oxygen vacancies that are introduced by the partial substitution of R 3+ ions for Ce 4+ , while proton conduction occurs when hydrogen dissolves in the lattice through hydroxyl ions that form.A variety of researchers have studied the influence of the dopant on the crystal structure and electrical properties of A-Ce1-xRxO3-δ perovskites (A = Ba and Sr, and R = Y, Nd, Gd, Yb) [32][33][34].

Due to dielectric and piezoelectrics, Transition Metal Stannates have wide application.Cadmium stannates have been studied for potential use in solar batteries as transparent electrode materials [35].Zn, Cu, Cd, Fe, Mn stannates are semiconductors with the disordered tructure of ilmenite and can be prepared under special conditions [36].Stannate doping (ASnO3) with transition d metal has been studied for the synthesis of cobaltcontaining oxide (Sr1-x LaxSn1-xCoxO3 [37] and ironcontaining oxide Sr(Sn1-xFex)O3 [38][39].It was found that the conduction differs from ionic-electronic to electronic form for Sr (Sn1-xFex)O3-δ solid solution, with an increasing Fe to Sn ratio.These solid solutions have a cubic structure within the concentration range of 0.4 < x < 0.9.

Recent research [40] has shown that reducing the size to the nanometer range of perovskite particles reduces their magnetization.It is therefore an important issue to increase the magnetization of nanocrystalline perovskite materials.Substituting on the A-site or B-site or both sites is one possible way to improve the magnetic properties or to increase electrical conductivity due to increased oxygen vacancies of ABO3 perovskites.A new series of substituted perovskite materials are synthesized in this paper with the formula Gd0.9Sr0.1Mn0.8Co0.2O3,Tb0.8Sr0.2FeO3,Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3,and Mn0.6Co0.4SnO3,and the effect of the substitution, on the structure, electrical and magnetic properties of these perovskites are studied.This precursor was prepared in double-distilled water using the co-precipitation method by taking Gd(NO3)3•6H2O (2.9272 g), SrCl2•6H2O (0.5765 g), MnCl2•4H2O (2.2924 g) and CoCl2•6H2O (0.2643 g).The medium's pH was changed to a sufficiently low value (pH < 5).A magnetic stirrer was used to stir the solution vigorously.Ammonium tartrate (4.4774 g in 25 mL distilled water) was then applied to this solution slowly until a permanent precipitate occurred.To get more homogeneous, stoichiometric, fine-grained powders, acetone was added in similar quantities.Pink in color was the resulting precipitate of (Gd0.9Sr0.1)Mn0.8Co0.2(C4H4O6)2.5•14H2O.The solution was filtered for 30 minutes after being stirred.For Gd 3+ , Sr 2+ , Mn 2+ , and Co 2+ , whose absence ensured absolute co-precipitation, the filtrate was tested.To accelerate the drying, the precipitate was washed with cold distilled water and then with acetone.At room temperature, it was air-dried.


I. Experimental

By following the procedure given above, other tartrate/hydroxide coprecipitates were prepared.Below are the additional related information summarized.


Terbium -strontium -iron tartrate two hydrate Tb0.8 Sr0.2 Fe(C4H4O6)2.5•2H2O

To avoid air oxidation of Fe 2+ ion, 0.5195 g iron metal powder (electrolytic) was dissolved in 5mL concentrated hydrochloric acid in a three-neck flask under a stream of dry nitrogen atmosphere.The solution on a low flame was then evaporated to dryness.The resulting crystalline green ferrous chloride residue was dissolved in double-distilled water (50 mL).1.3354 g Tb(NO3)3•6H2O (10 mL D.W.) and 2.2157 g SrCl2•6H2O (10 mL D.W.) were combined with this solution and placed into a flask (500 mL).It was applied slowly with stirring and to this ammonium tartrate solution (4.2338 g in 50 mL distilled water) was added, till a permanent precipitate occurred.It obtained a light green-colored precipitate.The precipitate was filtered and washed at room temperature with distilled water and acetone, followed by air drying.


Gadolinium -strontium -ruthenium tartrate four hydrates Gd0.6Sr0.4Ru (C4H4O6)3•4H2O

This precursor was prepared in double-distilled water by taking Gd(NO3)3•6H2O(0.8629g), SrCl2.6H2O(2.0393 g), and R

l3•3H2O (2.00 g).(50 mL).It was applied slowly with stirring and to this ammonium ta
trate solution (4.6195 g in 50 mL distilled water) was added until a permanent precipitate occurred.The graycolored precipitate was collected.To accelerate the drying, the precipitate was filtered and washed with distilled water and then with acetone.It was air-dried at room temperature.


Strontium -cerium -yttrium tartrate three hydrates Sr Ce0.95Y0.05 (C4H4O6)2.5•3H2O

The co-precipitation method was used to prepare this precursor by taking SrCl2•6H2O (2.6273 g), Ce(NO3)3•6H2O (3.0086 g), a

Y(NO3)3.5H2O(0.2698g) into double distilled water.To this ammonium tartrate soluti
n (4.0510 g dissolved in 25 mL distilled water) was added slowly with stirring, till a permanent white precipitate occurred.The precipitate has been filtered and washed at room temperature with distilled water and dry air.


Manganese -Cobalt -tin hydroxide, Mn0.6Co0.4Sn(OH)6

In 5 mL of concentrated hydrochloric acid (HCl), 2.0214 g of SnCl2•2H2O was dissolved and boiled to form a transparent solution.Dilute the solution to 25 mL

th distilled water.MnCl2•4H2O (1.0
38 g) and CoCl2•6H2O (0.8526 g) were dissolved in 10 mL of distilled water in another beaker.Then above SnCl2 solution was thoroughly mixed with these solutions.The reagent grade ammonia solution (0. M NH3) was added drop-wise to this solution until the precipitation is complete.A magnetic stirrer was used to stir the solution vigorously.The pH of the medium was maintained at ≤ 8.The grey precipitate was subsequently digested for one hour.To obtain content free of chloride ions, the solution was filtered and washed with diluted ammonia solution.At an ambient temperature, the precipitate was then dried and ground with the aid of the agate mortar.The powder sample obtained was restored using desiccators.


Synthesis of nanosized substituted perovskite compounds

In a platinum crucible under a static air atmosphere, the above precursors were decomposed and held slowly at 700 °C for two hours and then slowly cooled down to roo

temperature (30 °C/Min).This thermal treatment was adequ
te for the full decomposition of precursors to be achieved.The obtained powder was polycrystalline.This sample was then re-grounded for another two hours and re-calcined at the same temperature.The furnace was turned off and at room temperature, the sample was removed.In desiccators, substituted perovskite samples such as (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8Sr0.2FeO3,Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3and Mn0.6Co0.4SnO3have been restored.


Sample Characterization.

Microanalytical methods have been used to carry out the basic analysis of carbon and hydrogen.An inductively coupled plasma spectrometer (ICP-AES instrument ARCOS from M/S. Spectro Germany) and an Energy Di

ersive X-ray Analysis (ED
) on SEM PHILIPS XL 30 CP were used for the metal analysis of the sample.For this analysis, duplicate samples were used.Infrared spectra of the precursor were recorded using KBR pellets in the region of 4000 -450 cm -1 on the Perkin-Elmer 1600 series FTIR spectrophotometer.On a Mettler Toledo 850 instrument, the TGA and DTA were recorded.

The X-ray diffraction pattern (XRD) was obtained at room temperature using a Siemens D 500 diffractometer with CuKα radiation (λ = 1.5418Å and θ = 20 -80 o ).On a JEOL-2010 transmission electron microscope, working at an accelerating vol age of 208 kV, transmission electron microscopy (TEM) and corresponding selected area electron diffraction (SAED) pattern were carried out.TEM samples were dispersed by dispersing in 2-propanol.The amorphous carbon film supported on the copper grid was deposited with one or two drops of this solution containing synthesized nano oxides and allowed to dry in the air at room temperature.It determined the average particle size and the crystalline phase of the samples.The procedure used for the measurements of d. c. electrical conductivity and thermoelectric power measurements were identical to those reported in the literature [41].Measurements of the Hall effect of the samples were analyzed using an ECOPIA instrument (HMS 3000) with a constant magnetic field of 5 Tesla (τ).A HIOKI model 3532-50 LCR Hi Tester (frequency range from 100 Hz to 5 M Hz) at room temperature was used to study the dielectric properties of the substituted perovskite nano oxides at room temperature.

A SQUID magnetometer was used to investigate the magnetic behavior of the samples (Quantum Design MPMS-55).In the 100 Oe field, sample magnetization was tested for heating through the temperature range 5 < T(K) < 325 and after cooling the sample both in the measuring field (field cooling-FC) and in the absence of the applied field (zero-field cooling -ZFC).The magnetization was also measured at room temperature as a function of the field is applied.Data were collected in a field of 50 k Oe in decreasing and increasing fields cooling from the paramagnetic phase to the selected temperature.The measurement of magnetic hysteresis was also studied at room temperature using the vibrating sample magnetometer PAREG and G model no.4508 with an applied field of up to 20 k Oe.The infrared spectrum of these precursors showed the band assigned to the water of crystallization at 3362 cm -1 , and the intense band due to νasy.at 1598 cm -1 (C=O) and bands at 1411 cm -1 and 1238 cm -1 due to νsy.(C-O).These values indicate the presence of coordinated carboxylate groups.These precursors were assigned to a chain-like polymeric octahedral structure [42,43].Measuring magnetic susceptibility at room temperature confirms the precursor's paramagnetic existence, except for strontium-cerium-yttrium tartrate, (SeCe0.95Y0.05(C4H4O6)2.5.3H2O), which is diamagnetic in nature.


II. Results and discussion

The Manganese-cobalt-tin hydroxide infrared spectrum showing broadband at ~3348 cm -1 corresponds to the presence of -OH stretching vibration and bending vibration of water molecules.The existence of more than one hydroxyl typ

iplicity of 847 -1411 cm -1 ban
s resulting from the Sn-OH group's deformation vibration having some ionic character.For each transition metal oxyhydroxyl stannate, νSn-O vibrations have been reported at ~ 521 cm -1 , suggesting a six-coordinate setting for the metal ions [44].

TGA -DTA analysis was conducted to evaluate the thermal properties of the obtained samples.Fig. 1(a) to (e) show the simultaneous TGA and DTA curves in normal air of (Gd0.9Sr0.1)Mn0.8Co0.2(C4H4O6)2.5•14H2O, Tb0.8Sr0.2Fe(C4H4O6)2.5•2H2O,samples indicated a three step decomposition, while Gd .6Sr0.4Ru(C4H4O6)3•4H2O,SrCe0.95Y0.05(C4H4O6)2.5•3H2Oand Mn0.6Co0.4Sn(OH)6indicated a two step decomposition at a temperature range of 25 -800 °C respectively.An endothermic peak at around 100°C is attributed to desorption of physically adsorbed water and residual water molecules for these samples respectively.In TGA, the mass loss starts very slowly at 45 °C and at around 180 °C indicated the removal of water molecules.The dehydrated complex was almost stable upto 210 °C.The observed mass loss and corresponding temperature ranges are shown in Table I.The isothermally heated of these samples under normal air at 150°C respectively, the infrared spectrum showed the absence of the -OH band and a considerable reduction in the intensities of all carboxylate bands.X-ray diffraction pattern showed polycrystallinity of the samples with decrease in interplanar spacing as compared to the parent compounds.The elemental analysis were in good agreement with the respective anhydrous samples except manganese-cobalt-tin hydroxide, which show the formation of Mn0.6Co0.4SnO(O )4compound.The infrared spectra of this sample showed the presence of OH by the three modes of vibrations at 3410, 1620 and 870 cm -1 , which are assigned as stretching, in plane deformation and vibrational lattice mode respectively [45].

After dehydration, there were two strong and broad peak in the DTA curve at approximately 330 °C and 465 °C for (Gd0.9Sr0.1)Mn0.8Co0.2(C4H4O6)2.5 and Tb0.8Sr0.2Fe(C4H4O6)2.5 compounds which reveals to the formation of the final oxide.The TGA showed two-step mass losses in the temperature range of 180 -410 ° and 410 -610 °C, corresponding to the formation of the respective carbonates and final perovskite oxides (see Fig. 1(a) and (b)).The observed mass losses were in reasonable agreement with the calculated values (see Table I).

The infrared spectra of the isothermally heated parent samples at 380 °C showed the bands at 1465 cm -1 , 1380 cm -1 and 1150 cm -1 to carbonate group frequency , along with the weak bands at 575 cm -1 and 430 cm -1 due to M-O stretching frequency, and no band was found at 1590 cm -1 indicating the absence of carboxyl (-COOH) group.The X-ray diffraction pattern of these isothermally heated samples showed generally broad and amorphous nature.The elemental analyses were confirmed with respective carbonate samples.This is also qualitatively test by taking small amount of these solid and adds 2 ml distilled water in a test tube.Then a dropper full of 3M H2SO4 trickle down side of test tube.The formation of small bubbles is noted, which indicate carbonate present.The another qualitative test is to prepare a thin film of baryte water (Ba (OH)2) on small platinum wire loop.Then hold near the mouth of test tube which contain above isothermal sample and HCL.The film turns opaque due to BaCO3.This is also confirmed carbonate present.The sample thus obtained at 380°C confirmed the presence of respective carbonate.For sample heated isothermally at around 490 °C, in the infrared spectra showed two absorption bands in the region 573cm -1 and 424cm -1 .These bands may be tentatively assigned to M-O stretching and O-M-O bending vibrations of perovskite phase formation respectively [45].The spectrum further showed the absence of carbonate group frequency.The X-ray diffraction pattern showed generally sharp lines, and the pattern fit with the respective perovskite oxide [46,47].The metal analysis carried out by ICPS analysis agreed well with that of the respective oxides.

The oxidative decomposition of anhydrous precursors such as Gd0.6Sr0.4Ru(C4H4O6)3and SrCe0.95Y0.05(C4H4O6)2.5 were indicated by presence of strong and broad exothermic peak on the DTA curve at around 330 °C corresponding in the formation of final oxides.

Similarly the decomposition of Mn0.6Co0.4SnO(OH)4sample was indicated by an endothermic peak in the DTA curve at 210 °C .The TGA curves for these complexes showed a continuous mass loss from 200 to 600 °C (see Fig. 1 (c), (d) and (e).These mass losses were found to be in good agreement with the formation of respective perovs ite oxide as the final product (see Table I).The X-ray diffraction pattern for the parent samples decomposed isothermally at around 500°C showed mainly respective perovskite oxide [48][49][50].

The infrared spectrum of the isothermal product , there are two strong absorption band at about 580 cm -1 and 422 cm -1 (or 650 cm -1 and 552 cm -1 for Mn0.6Co0.4SnO3perovskite) which correspond to M-O stretching and O-M-O bending vibration respectively for perovskite oxides [45].However the infrared spectra of Gd0.6Sr0.4RuO3per vskite showed these M-O bands along with a strong band at 2325 cm -1 correspond to carbon dioxide present on the isothermal product.The X-ray diffraction pattern of this sample showed a slightly broad pattern, probably that of adsorbed carbon dioxide on the perovskite (i,e Gd0.6Sr0.4RuO3(CO2)).When this   II, The compositions observed are summarized.They are ± 0.5 % of the nominated value.This finding agrees with the findings of inductively coupled plasma (ICPS).The energy dispersive X-ray analysis (EDS) also demonstrates the existence of the same cationic ratio of all-metal species as originally used (Table II).In the EDS spectrum, the peaks for all the cations were present (Fig. 2).At many places, the elemental composition analysis was uniform, which is representative of a highly homogeneous substance.This is due to the uniform mixing of all the cations.When strontium ions are substituted by rareearth ions or related ions, the negative effective charges of strontium ions must be compensated for by the creation of equal quantities of positive effective charges.This can either constitute M 4+ ions (e.g.M 4+ can be Mn, Fe, Ru, etc) or vacancies of charged oxygen.EDS and ICPS techniques have calculated the concentration of these ions (see Table II).


Structural studies

Structural analysis is made by X-ray diffraction (XRD) on the substituted perovskites and the diffractograms were obtained from a Siemens D500 diffractometer using CuKα radiation at 2 °/min scanning rate in the 2θ range of 20-80°.The XRD pattern acquired for substituted perovskite such as (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8Sr0.2Fe2O3,Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3and Mn0.6Co0.4SnO3are shown in Fig. 3.There is an excellent agreement of experimental XRD pattern with the standard JCPDS pattern confirming the monophasic perovskite structure of these substituted oxides [46][47][48][49][50].The X-ray diffraction peaks are found to be broad and it is attributed to the nanocrystalline nature of these perovskites.The peaks for these substituted perovskite appear approximately at same position as for the respective unsubstituted perovskites but with different intensities.In these cases, the substituent ions seem to dissolve / arrange in the perovskite structure to fulfill the formation of single perovskite phase.The hkl values are assigned to the peaks in the XRD are shown in Fig. 3.The lattice parameter for these substituted compounds are determined from the XRD pattern using the inter-plane spacing d of the diffraction peak and hkl denotes the miller indices of the corresponding peak in the orthorhombic or cubic relation.The correct structured orthorhombic perovskite except for the Tb0.8Sr0.2Fe2O3and Mn0.6Co0.4SnO3compounds, which is found to be a cubic and rhombohedral structure respectively, was revealed by all substituted compounds.The d-spacing and relative intensity values observed experimentally are compared with those recorded in the literature [46][47][48][49][50].The lattice par

ers, unit cell volu
e, average cation sizes of Asite perovskite < RA >, and mismatch factor (σ 2 ) for each compound are listed in Table III.It is shown that for all samples, the observed lattice parameters and unit cell volume are substantially decreased compared to the reported unsubstituted perovskites.(see Table III) [46][47][48][49][50].In agreement with the lanthanide contraction, the unit cell decreases.The bond distance between elements decreases when an electron is added to the 4f orbital of rare-earth (RE) elements.Therefore the    The size of the cation on the A-site affects the crystal symmetry significantly in perovskite (ABO3), while that of the cation on the B-site does not change the symmetry, but proportionally changes the volume of the lattice (i.e. unit cell volume).Rodrigues and Attfield [51] argued that local distortions and stresses in the structure were caused by size differences between different A-site ions.In terms of the variance of the distribution of cation radii, σ 2 = ∑ {xR 2  A -< RA >2}, a quantitative explanation of this effect was given.Here at the A-site species of the perovskite, x is the fractional occupancies for different cations, RA their ionic radii, while < RA > is the average cation size.The effects of the mismatch factor (σ 2 ) have been determined for the various cations of substituted perovskite compounds and are shown in Table III.Standard ionic radii [52][53] i.e.Gd 3+ (1.107 Å), Sr 2+  (1.310 Å), Tb 3+ (1.095 Å), Ce 3+ (1.196 Å), Y 3+ (1.075 Å), Mn 3+ (0.580 Å), Co 3+ (0.545 Å), Fe 3+ (0.550 Å), Ru 3+ (0.681 Å) and Sn 4+ (0.55 Å) were used for the calculation of < RA > mean cation sizes and 8-coordination mismatch factor (σ 2 ) for orthorhombic or cubic perovskites.The average cation size of cations at the A-site of perovskite < RA > is less than 1.21 Å, as is evident from Table III, suggesting the absence of long-range cooperative Jahn-Teller distortion in the compounds [54].However, the probability of dynamic Jahn-Teller short-length distortion cannot be eliminated [55].

The additional disorder is caused by the substitution of a small amount of Sr 2+ ion and can occur between the A-site cations due to the size mismatch factor (σ 2 ) [51].We have observed that the mismatch factor (σ 2 ) varies with compounds (Table III).This is due to segregation of the A-site cations, as a result of the large size disparities, so that the compounds are no longer microscopically homogeneous, although they appear to be concluded, provided the constant A-site cation average radius < RA > = 1.21 Å, the variation of the A-site disorder as characterized by variance σ 2 causes no change of the crystal structure.In addition to these effects, the strain field associated with local oxygen ion displacement is likely to become so small that no chemical phase separation occurs in the system.Oxygen stoichiometry was calculated by iodometric titration for all substituted substances, the error in the determination being ± 0.2.It was found that all the compounds in the 2.98 to 3.00 range had the same amount of oxygen.This means that the stoichiometry of oxygen in these compounds is retained perfectly.Using the relationship
[56] Q = √2 (< R 2 A > -< RA > 2)
, the local oxygen displacement Q was determined.The Q values are observed in the range 0.000067 to 0.000416.It concludes that in all substituted compounds, a negligible amount of oxygen displacement occurs.Tolerance factors for these compounds were determined using the relationship t = [<RA> + R(O)] / √2[<RB> + R(O)], where t = tolerance factors, R(O) = 1.40 Å is the ionic size of the O 2-, (for six coordinated radii), < RA > and < RB > are the mean ionic sizes of the A and B-site ions, respectively, for the perovskite structure, respectively [52,53].Tolerance factors (t) are shown in Table III.The tolerance facto s ranged from 0.6907 to 0.9072, well within the 0.68 -1.00 range which characterizes the stable perovskite structures.

We calculate the particulate properties such as mean crystallite size < D > X-ray, density (Dx), apparent density (D), and porosity (P) for all substituted compounds are summarized in Table IV.As can be seen from the table, the crystallite sizes < D > X-ray are in the range between 14.85 nm to 47.72 nm.The apparent density (D) observed is lower than the X-ray density (Dx).This low density indicates the formation of a microstructure with a smaller average grain size.

For all substituted perovskite samples, the morphology of the particles produced was examined by scanning electron micrographs (SEM) (Fig. 4).The grain shape of (Gd0.9Sr0.1)Mn0.8Co0.2O3 is spherical, but with less degree of agglomeration.In comparison to this, the grains of Tb0.8Sr0.2FeO3and SrCe0.95Y0.05O3have a high degree of agglomeration with spherical samples.The grains of Gd0.6Sr0.4RuO3were plate-like and no agglomerated, while the grains of Mn0.6Co0.4SnO3had needle-like less agglomerated; but some needle grains have already been destroyed.Fig. 5 displays the TEM images of substituted perovskites.As can be seen for (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8Sr0.2FeO3and SrCe0.95Y0.05O3,(Fig. 5 (a)), (b), (d), the particles are more or less spherical with a size range between 22 nm to 77 nm for these samples.Square particles were present in the TEM picture of the Gd0.6Sr0.4RuO3sample (Fig. 5(c)); in contrast to this, the particles Mn0.6Co0.4SnO3(Fig. 5 (e)) had needle shape particles.Because it corresponds to the inner shel (ordered) region of the particles, without the severely distorted peripheral part, the mean crystallite size determined by XRD is typically smaller than the average particle size determined by TEM.The transmission electron microscopy is a local technique and determines the size of the particles in a limited field of view particle size determination from dynamic light scattering data is highly dependent on the powder dispersion process in glycerol, intended to deagglome ate the nanoparticles.Mechanical mixing, however, breaks down not all agglomerates, so the measured average particle size usually exceeds the XRD mean crystallite size.

Structural information from the selected field of electron diffraction (SAED) was also collected for all samples shown in Fig. 5.It is apparent that all the lattice planes are clearly distinct, suggesting that there is no defect in the single crystalline particle.In the SAED pattern, the various lattice planes are allocated and these are in good agreement with the crystallographic phase of the substituted perovskite system.

An infrared spectrum of a molecule is considered to be unique physical property and is characteristic of the molecule.The infrared spectrum for all substituted perovskite consists of two absorption bands associated with the stretching and bending vibration in the region 400 -900 cm -1 .The stretching vibration is assigned to an absorption band in each range at the higher frequency ν1 = 542 cm -1 and the bending vibration at the lower frequency ν2 = 442 cm -1 (Table IV).This is due to the change in the potential energy in the two natural vibrations due to repulsive forces between metal ions and oxygen ions.


Electrical conductivity studies

Fig. 6 indicates the dependence of log σ against the reciprocal temperature (T -1 ) for (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8Sr0.2FeO3,Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3and Mn0.6Co0.4SnO3compounds.With increasing temperature, the conductivity increases linearly.At a certain temperature, Tc indicates a break in the line of each compound, resulting in a change in the slope.The slope shift is typically supposed to occur at the Curie temperature or due to th presence of two parallel mechanisms of conductivity mechanism with different activation energies [57].For all compounds, the activation energy Ea (eV) was determined from the slopes of the line before and after transition point Tc and is tabulated in Table V.The values of activation energies imply that the conduction is due to the process of polarons hopping mechanism.In Region A the value of is found to be lower than in Regi n B. The conduction of impurity is favored below 500 K, thus hopping of polarons at higher temperatures.The sample of Gd0.6Sr0.4RuO3(Ru 3+ has a 4d 5 configuration) showed semiconducting rather than metallic properties.The electrical conduction would then become an activated process, via transfer of an electron from t2g (П*) to eg (σ*) levels, which would presumably also exist as band states i.e σ-bonding being stronger than П bonding.Therefore the electrical behavior in these samples under investigation is due to the electron acceptor hopping in the sublattice between Mn 3+ ↔ Mn 4+ or Ru 3+ ↔ Ru 4+ or Ce 3+ ↔ Ce 4+ ions of the respectively substituted perovskite under investigation results in increased activation energy (Ea) in the region B. The A-site cation disorder has previously been shown to decrease the break temperature, Tc [58].In all substituted perovskite compounds, it seems that the A-site cation mismatch factor (σ 2 ) tends to play a more important role than the average cation sizes < RA >.

The significantly different transport behavior on various mismatch factor values (σ 2 ) enables us to argue the essential role of A-site disorder in mediating the phase separation of B-site cations, nothing that no charge order phase appears upon the A-site disordered.In other words, the addition of a different substituent can be affected by many factors like the nature of the substituent, the amount of substituent incorporated creating structural defects, and changing its microstructure.Therefore, we turn to the magnetic behavior of these compounds in order to reveal the magnetic ground state in the B-site cations of different A-site disorder degrees.

The temperature effect of all substituted perovskites on the thermoelectric power is shown in Fig. 7. Positive charge carriers for samples of (Gd0.9Sr0.1)Mn0.8Co0.2O3and SrCe0.95Y0.05O3decrease slowly with temperatures of Fig. 6.Plot of log σ against T -1 of nanosized substituted perovskite compounds.up to 500 K and then almost constant as temperatures rise.In the entire temperature range, the positive values of the thermoelectric power confirmed the p-type semi conduction.Therefore the number of Mn 3+ or Ce 3+ ions at the B-site of these compounds is greater than the number of Mn 4+ or Ce 4+ ions at temperatures up to 500 K.As electron transform from Mn 3+ → Mn 4+ or Ce 3+ → Ce 4+ .The motion of Mn 4+ or Ce 4+ ions in the lattice of Mn 3+ or Ce 3+ ions can be visualized as the motion of a hole.

The thermoelectric power of Tb0.8 Sr0.2FeO3 and

0.6Co0.4SnO3compounds, on the ot
er hand, showed that they were n-type semiconductors.This change in semiconductivity in the perovskite structure depends on the existence of both the A-site cation and B-cation.In general, (i) greater A-site cation causes greater reduction stability.Which is n-type of semiconductor.These findings suggest that the substituent is an electron donor in the conduction band also creates more vacancies.Here, oxygen vacancies act as donor's centers.(ii) Secondly, rare-earth ions (such as Gd 3+ , Tb 3+ ) are Fig. 7. Plot of thermoelectric power, μ (volt K -1 ) against temperature, T (K) of nanosized substituted perovskite compounds.becoming more acidic.It competes more strongly with the B-site for binding to oxygen atoms of the perovskite.(iii) Stannate have a continuous edge-sharing of transition metal ions (i.e.Mn 2+ and Co 2+ ) and Sn 4+ octahedra is a necessary criterion for the formation of semi-metallic conductivity [59].

Similarly, n-type semiconductivity is demonstrated by the Gd0.6Sr0.4RuO3compound.This could happen only charge carriers are localized on the defect centers, which are inherently present in the solids.The Ru 4+ ion obviously traps carriers.In which, the electron hopping from site to site within the 4d band of ruthenium and secondly, the reduction by Ru-O distance and Ru-O-Ru angle, which corresponds to the electron hopping from Ru 4+ to Ru 3+ ions, respectively We have measured the Hall effect constant in the 0.54 T magnetic field at room temperature after passing current in nA through samples.

Table V reveals that (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8 Sr0.2FeO3, Gd0.6Sr0.4RuO3,Mn0.6CO0.4SnO3are n-type semiconductors due to electron hopping between A site cations, while SrCe0.95Y0.05O3suggests p-type semiconductors due to hole hopping between B site cations.The specific resistivity values were observed for all compounds around 4.199 x 10 6 to 3 705 x 10 10 ohm.cm at current passed ranging from 10 to 50 nA.The variations in the observed specific resistivities were also reflected in the mobility.The Hall coefficient at room temperature for all compounds under investigation is except the SrCe0.95Y0.05O3compound indicating electrons are the majority carrier.The SrCe0.95Y0.05O3shows p-type conductivity.The variations of the observed Hall coefficient are lso reflected in mobility.


Dielectric studies

Fig. 8 (a) and (b) show the effect of frequency (ln F/Hz) on the dielectric constant (ε') and the dielectric loss factor (tan δ) of the synthesized samples such as (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8 Sr0.2FeO3, Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3,and Mn0.6Co0.4SnO3 at the room temperature.The values of dielectric constant decreases continuously with increasing frequency (see Fig. 8(a)) and this may be due to the combined contri ution from ionic, space charge, electronic and interfacial polarizations.However, in the highfrequency region (> ln 5Hz), the variation of dielectric constant (ε') is very much low with the increasing frequency.The variation of dielectric constant with frequency (ln F/Hz) is well understood by the Maxwell-Wagner model [60].Accordingly to this model, any dielectric material is made up of two differently conducting layers, first layer being a conducting grains and other being isolated by grain boundaries that are poor conductors.The electric field in a dielectric material can stimulate the migration of charge carriers or mobile ions.However these charge carriers get accumulated at the interface region of the grain-grain boundary due to their different conducting natures.Thus, the accumulation of charge carriers can lead to polarization which in turn leads to the dielectric constant (ε').In literature, for the nanocrystalline regime, the grains and grain boundaries are large in number than the microcrystalline regime and they are found more complex in lower and higher frequencies, respectively in nano materials [61,62].At low frequencies region, the mobile charge carriers / oxygen ions in the substituted perovskite are free to align towards the applied electric field due to sufficient time available and migrate through grains and grain boundaries.These mobile charges get accumulated at the grain-grain boundary interface and therefore more dielectric constant (ε') is shown by the nanocrystalline substituted perovskites.With the increase in frequency (in Hz) of the applied field, the mobile charge carriers gradually lose their ability to follow the field and become stagnant at one point of time.This leads to decrease the value of total dielectric constant (ε') and remain constant at higher frequencies for these substituted perovskites.

The frequency (ln F / Hz) -dependent dielectric loss (tan δ) of these substituted perovskite at room temperature are presented in Fig. 8 (b).The value of loss tangent (tan δ) depends on many factors like structured homogeneity, stoichiometry and synthesis procedure.Also tan δ represents the difference in phase between the applied electric field and the current induced in the dielectric material.Similar to dielectric constant (ε'), the dielectric loss in substituted perovskites decreases at a higher rate (up to ln 4.5 Hz) with the increasingfrequency and the rate is relatively very low beyond a certain  .At low-frequency region, tan δ have high value and more or less stable at the highfrequency region.This dispersion of dielectric loss with frequency is related to Koop's phenomenological theory, similar to the Maxwell-wagner model [63].At low frequency region, which corresponds to high resistivity (due to grain boundaries) more energy required for electron exchange between M 2+ and M 3+ ions; thus the tan δ loss is high.In the high frequency range which corresponds to low resistivity (due grain) less is needed for electron transfer between M 3+ to M 2+ in the grains and hence the tan δ is small.Therefore, at higher frequencies exhibit a constant tan δ, which suggest all samples possess a loss nature.This suggests all substituted perovskites could be utilized in hi

frequency device a
plications.For this sample, the remanent magnetization and coercive force could not be properly calculated from the hysteresis loop.It should be noted that even after applying a magnetic field of 60000 Oe (or 6T); the magnetization value did not show saturation.In the sample, this shows the existence of magnetic relaxation.The explanation for not having a saturated value of magnetization may be the very fine particles (nano size) present in the sample undergoing superparamagnetic relaxation at room temperature.Another reason is that the absence of saturation in the high field magnetization (6T) indicates a

in-c
nted ferromagnetic state in this sample (i.e.antiferromagnetic exchange coupling between Gd 3+ and Mn 3+ -Mn 4+ sublattice [64].


Magnetic

For (Gd0.9Sr0.1)Mn0.8Co0.2O3samples, the temperature dependence of magnetization of field-cooled (FC) and zero field-cooled (ZFC) measured at 100 Oe by a SQUID magnetometer is shown in Fig. 9 (a').This sample shows relatively soft ferromagnets, with a rapid increase in magnetization occurring at a decreasing temperature.For this sample, the Curie temperature Tc is determined from magnetization curves at 150 K. Thermal irreversibility Magnetization of FC > ZFC is observed below 156 K at 100 Oe.The FC is separated from the ZFC due to the presence of an anisotropy arising from ferromagnetic clusters.This feature is widespread in unconventional ferromagnets, such as magnetic perovskites [65,66]

n other words, th
re is a large discrepancy between ZFC and FC curve below Tc which may imply a complex magnetic state after the magnetic ordering transition.This is due to the localizing influence of A-site of perovskite cation disorder is a more prominent factor than the average cation sizes < RA > that would enhance the double-exchange mediated ferromagnetism.On Sr 2+ substituting causes the broadening and shift to lower temperature magnetization.This is related to the mismatch factor (σ 2 ) causes a decrease of Tc, which inhibits charge more strongly than it affects the ferromagnetic double exchange to the disorder in the lattice of the perovskite phase.This compound has plateaued in field-cooled (FC) curve response at the temperature 145 K to 100 K, followed by a steady rise down to 35 K and finally a decrease in magnetization below 35 K (see Fig. 9(a').The magnetization of this sample first goes from paramagnetic to ferromagnetic phase followed by the charge-ordered antiferromagnetic phase coexist.It is also een in Fig. 9 (a') that a decrease in magnetization of ZFC much below the ferromagnetic ordering temperature and a continuously increasing magnetization of FC, when measured at low magnetic fields (100 Oe), have also been observed for other ferromagnetic systems showing long-range ferromagnetic order much below the temperature at which a drop in magnetization of ZFC is observed [67].Such change in the magnetic structure can be explained based on random distribution of Gd 3+ ions i.e. disorder state results in the random magnetic potential on the Mn sublattice.So long-range ferromagnetic order in Mn-O-Mn is impossible, resulting in Mn-O-Mn cluster and the forming of the cluster-spin glass state.Because long-range ferromagnetic order is destroyed, the Tc is decreased and magnetization is weakened on Gd 3+ or Sr 2+ doping.


(b) For Tb0.8Sr0.2FeO3

To confirm the coexistence of ferromagnetic coupling and antiferromagnetic coupling, the temperature dependence of magnetization of the Tb0.8Sr0.2FeO3sample at the high applied field (10 K Oe) is shown in Fig. 9(b).It is evident from this Figure that a paramagnetic contribution appears superimposed to the ferromagnetic one.This contribution is also present in M (H) curves making the evolution of the actual value of the saturation magnetization difficult.It clearly indicates that only a fraction of Fe atoms participate in ferromagnetism, the rest contribute to the observed paramagnetic signal or may even be antiferromagnetically coupled.The curves show a rapid decay of magnetization at the low-temperature regime reflecting that the paramagnetic behavior predominates at the low-temperature regions.These results indicate that the existence of a strong antiferromagnetic interaction between Fe ions in the sample.The estimated Neel temperature is observed at TN = 87 K.The large Tb 3+ paramagnetic signal hinders the direct observation of iron ordering in the sample with higher Tb 3+ content.However, such ordering can be inferred from two features (i) TN is visible at low temperature for this compound and (ii) The out-of-phase component of magnetization a sharp rise at the same temperature TN.Such behavior also found for TN in related series, has been attributed to the large exchange coupling of Fe 3+ ions and can be attributed to a small spin canting in the antiferromagnetically orde

d state [68].However, t
e magnetizations do not go to zero (see Fig. 9(b) and the hysteresis loop is not observed at 300 K.These results demonstrate that only a minor Fe spin is ferromagnetically coupled and the predominant Fe spins remain uncoupled or even antiferromagnetic coupled.

The temperature-dependent magnetization curve of the Tb0.8Sr0.2FeO3sample in zero field cooled (ZFC) and field cooled (FC) conditions in an applied 200 Oe is shown in Fig. 9(b').A nonzero difference between FC and ZFC curve indicates ferromagnetic ordering.By taking the difference between these two quantities (denoted as ∆M = M (FC)-M (ZFC)), para-and diamagnetic contribution to the magnetization can be subtracted, leaving only a measure of the hysteretic ferromagnetic regime.Simultaneously, a non-zero difference indicates the presence of hysteresis, i.e. ferromagnetic ordering.This strongly supports the bound magnetic polaron (BMP) model for ferromagnetism in solid [69].According to this model, a donor spins of defects strongly correlating with Fe 3+ within its orbit can mediate effective interactions between them based on a Heisenberg exchange Hamiltonian [69].It can be also seen that ZFC and FC magnetization curves show a distinct irreversible behavior.The irreversible temperature (Tirt) is observed at 300 K for this sample.Tirt is the temperature below which the ZFC and FC curve split.The reduction in particle size (i.e.nanoparticles) causes in Tirt towards lower temperature.

(c) For Gd0.6Sr0.4RuO3Fig. 9(c) shows the magnetization as a function of the applied field corresponding to the Gd0.6Sr0.4RuO3sample at room temperature.The sample shows a clear narrow hysteresis loop with less coercivity, but certainly, it is still strongly ferromagnetic.It may be noted the magnetization value did not show saturation e en after applying a magnetic field of 6 T.This indicates the presence of magnetic relaxation in the sample and is believed to be governed by antiferromagnetic interaction due to high magnetic ion (Ru 3+ ) in the system.The very fine particles present (nanosized) in the sample undergoing superparamagnetic relaxation at room temperature could be the reason for not getting a saturated value of magnetization.

We have also investigated the magnetic properties of this sample (Gd0.6Sr0.4RuO3)using a SQUD magnetometer in the temperature 5 -325 K.The temperature dependence of magnetization is shown in Fig. 9(c').The figure shows a plot of zero-field-cooled (ZFC) and field-cooled (FC) magnetization measurements performed on this material by applying a field of 100 Oe.It can be seen that with temperature decreasing this compound undergo a transition from paramagnetic to a ferromagnetic state.The Curie temperature Tc (defined as the temperature corresponding to an abrupt increase of magnetization) observed for this compound at 200 K.This sample sh ws a distinct splitting between ZFC and FC measurement implying the presence of a magnetic transition temperature.This sample also exhibits a peak in its magn tization curve at 100 K.A large number of Gd 3+ paramagnetic signals hinder the direct observance of ruthenium ordering in the sample [68].Above Tc, the magnetization variation in the Gd0.6Sr0.4RuO3system obeys the Curie-Weiss law.The obtained effective paramagnetic moments reasonably agree with the sum of contribution expected from Gd and Ru ions.Both kinds of Ru ions (Ru 3+ and Ru 4+ ) are high spin state.

Also in this sample (Gd0.6Sr0.4RuO3),when Sr 2+ is substituted for Gd 3+ , its larger size and lower charge enable it to compete more strongly for oxygen electrons.At the same time, the Ru 3+ resulting from the substitution are less strongly bonded than the Ru 4+ , because of its larger size and smaller charge.These effects are additive (i e.superexchange interaction) both leading to a decreasing Ru-O interaction and presumably a narrower π*-bond.This effect is quantified by the influence of tolerance factor (t) on the superexchange interaction.This tolerance factor (t) is listed in Table III.For a perfect size match, t = 1, the Ru-O-Ru bond angle (θ) would be 180°.For t < 1 rather than a simple contraction of the bond distance, the octahedral tilt and resulting in θ < 180°.Thus the apparent decrease in tolerance factor (t) with decreasing θ for Ru -O -Ru bond angle at Bsite of perovskite.At Gd0.6Sr0.4RuO3, the bandwidth has been sufficiently decreased so that antiferromagnetic behavior is observed for this sample.

(d) For SrCe0.95Y0.05O3The temperature dependence of magnetization at a high applied field (10 K Oe) for the SrCe0.95Y0.05O3sample is shown in Fig. 9


(d).

There is an increase in magnetization because of oxygen vacancies in the sample.It is well known that cerium in SrCe0.95Y0.5O3can be Ce 3+ or Ce 4+ valence states.So the substitution of a rare-earth ion such as Y 3+ should create oxygen vacancies [70] and finally enhance the ferromagnetism.Thus, the appearance of such behavior is due to increasing the cation size at the A-sit of perovskite, which inhibits charge ordering more strongly than it affects the ferromagnetic double exchange to the disorder in the lattice of the perovskite phase.

Ferromagnetic features are also found in the SrCe0.95Y0.05O3sample by studying the temperature dependence of magnetization after a ZFC-FC process under an applied field of 100 Oe (Fig. 9 (d')).ZFC curve is obtained by initially cooling the sample from 325 to 5 K in the absence of magnetic field and heating up in 100 Oe field while recording magnetization.For the FC measurements, the sample is cooled in the same field, and magnetization is recorded while heating up (FC).It is observed that the Curie transition temperature (Tc) of this sample is above 350 K.As in the case of the sample, the magnetization against the temperature curve show irreversibility between the ZFC and FC curves due to the blocking process of superparamagnetic particles.It is worth noting that at low temperature (~ 30 K), the steep rise in magnetization seems to be intrinsic to diluted magne

c oxide, and also indica
es that the presence of a paramagnetic phase at low temperature.In addition, since the particle is in the nano regime, some of the particles could behave superparamagnetically and the two contributions (paramagnetism and superparamagnetism) may complicate the magnetic behavior at low temperatures.


(e) For Mn0.6Co0.4SnO3

The temperature dependences of magnetization of Mn0.6Co0.4SnO3taken during heating in a magnetic field 200 Oe is shown in Fig. 9 (e).The magnetization of this sample exhibits a sharp peak at 160 K, an abrupt decrease of the moment below this temperature then remains constant up to 75 K and a step like an upturn at a lower temperature (~ 20 K).The anomalies of physical properties at Tc = 160 K in this sample are due to the transformation from the high-temperature paramagnetic state into the low-temperature antiferromagnetic state.There is the presence of a short-range magnetic interaction above the transition temperature Tc.It has been argued that the presence of Mn 3+ (or Co 3+ ) and Mn 4+ (or Co 4+ ) ions and with spin-orbit coupling could be the source of the anisotropic nature of the ferromagnetic exchange interactions.While Sn 4+ is diamagnetic in this compound [36,37].The predominance of ferromagnetic exchange interaction in the high temperature would lead to the formation of the ferromagnetic state if the antiferromagnetic state had not occurred at Tc.The over-cooled clusters are presumably present in the compound at low temperatures.This is because the virtual Curie temperature in this compound practically coincides with the temperature of the magnetic transition.Fig. 9 (e') shows the temperature dependence of magnetization measured in an applied magnetic field of 200 Oe in both zero field cooled (ZFC) and field cooled (FC) modes for Mn0.6Co0.4SnO3sample.This compound shows large thermomagnetic irreversibility at applied (200 Oe).In

FC and FC magnetization
branches were observed below the temperature Tc = 150 K of gradual rise.The weak ferromagnetic ordering produced branching of ZFC and FC magnetization below Tc (Fig. 9 (e')).The weak ferromagnetism is due to the canted antiferromagnetic interaction between Mn or Co ions via oxygen i.e. superexchange interaction.

Moreover, on cooling, antiferromagnetic interaction occurred at Tc = 150 K for this compound underwent long-range antiferromagnetic ordering with weak ferromagnetism (corresponding to the gradual rises).This is the first observation in this compound.The difference in the onset of the antiferromagnetic ordering of this compound originates from the magnetic interaction between Mn or Co ions via oxygen, i.e., the Mn-O-Mn (or Co-O-Co) bond angle.The long-range antiferromagnetic ordering appeared at a higher temperature in this compound.


Conclusions

1.The X-ray diffraction pattern of single-phase perovskite with orthorhombic structure For (Gd0.9Sr0.1)Mn0.8Co0.2O3,Tb0.8Sr0.2FeO3,Gd0.6Sr0.4RuO3,SrCe0.95Y0.05O3and Mn0.6Co0.4SnO3compounds, except Tb0.8Sr0.2FeO3Mn0.6Co0.4SnO3compounds, which showed cubic and rhombohedral structure respectively.The lattice parameter and unit cell volume slightly decrease as ionic radii decrease in agreement with the lanthanide contraction.The average size of cations < RA > of A-site of perovskite, mismatch factor (σ 2 ), and tolerance factor (t) give combine effects of disorder and inhomogeneity in these perovskite compounds.

2. The temperature dependence of electrical conductivity for all substituted compounds show a definite break in 500 to 610 K except the Gd0.6Sr0.4RuO3compound, which corresponds to semiconducting behavior.While the Gd0.6Sr0.4RuO3sample shows a metallic-like semiconductor.The activation energy (Ea) in the first region corresponds to the charge carriers are localized on the defect center, while Ea in the second region corresponds to an intrinsic property of the compounds.Thermoelectric power and Hall effect measurement shows all compounds are n-type semiconductivity except SrCe0.95Y0.05O3compound, which show P-type semiconductivity.

3. Magnetic studies show the thermo-magnetic irreversibility for all substituted compounds.These compounds underwent long-range antiferromagnetic ordering with weak ferromagnetism in temperature dependence (5 -325 K) of magnetization measurements.Enhanced ferromagnetism in most samples can be related to increased surface defects due to