Journal of Alloys and Compounds 485 (2009) 372–378

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Journal of Alloys and Compounds

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Preparation and properties of the CoFe2O4–Nb–Pb(Zr,Ti)O3 multiferroiccomposites prepared in situ by gel-combustion method

C.E. Ciomaga a,∗, C. Galassi b, F. Prihor a, I. Dumitru a, L. Mitoseriu a,A.R. Iordan c, M. Airimioaei c, M.N. Palamaru c

a Faculty of Physics, Al. I. Cuza University, Iasi 700506, Romaniab ISTEC-CNR, Via Granarolo No. 64, I-48018 Faenza, Italyc Faculty of Chemistry, Al. I. Cuza University, Iasi 700506, Romania

a r t i c l e i n f o

Article history:Received 20 March 2009Received in revised form 15 May 2009Accepted 20 May 2009Available online 28 May 2009

Keywords:

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Di-phase magnetoelectric composites formed by magnetic CoFe2O4 and ferroelectric Nb–Pb(Zr,Ti)O3 wereprepared by citrate–nitrate combustion method technique by using Nb–Pb(Zr,Ti)O3 template powdersprepared by mixed oxides. Pure di-phase powder composites with a good crystallinity were obtainedafter the calcination step. The purity was maintained by sintering at temperature of 1100 ◦C/2 h, whichensured limited reactions at interfaces, while by sintering at 1250 ◦C/2 h, small amounts of secondaryphases resulted. The method allowed to produce di-phase ceramics with homogeneous microstructuresand a very good mixing of the two phases. The dielectric and magnetic investigations at room temperature

Composite materialsMicrostructureX-ray diffractionS

confirmed the formation of composite ceramics with both dielectric and magnetic properties at roomtemperature, with permittivity and magnetization resulted as sum properties from the parent Pb(Zr,Ti)O3

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canning electron microscopy and ferrite phases.

. Introduction

Materials in which coupled ferroelectric and ferromagneticrders occur simultaneously in the same range of temperaturesre known as magnetoelectric (ME) multiferroics and can beither single-phase or composites (di- or polyphasic systems)1,2]. Most of the known single-phase ME materials are still fareyond the practical needs because they show a very weak MEoupling and also low values of the ME coefficient (of the orderf 1–20 mV/cm Oe) [3]. Another alternative method to obtainultiferroics is to produce composite materials consisting of piezo-

lectric and piezomagnetic phases. It has been reported that suchomposites may exhibit a larger ME effect than single-phase mate-ials [4]. The conceptual conditions for the ME effect in compositesere pointed out by Boomgaard [5], as following: (i) the individ-al phases should be in chemical equilibrium, (ii) no mismatchetween grains, (iii) high values for the magnetostriction coeffi-ient of the piezomagnetic or magnetostrictive phase and of the

iezoelectric coefficient of the ferroelectric phase, (iv) no leak ofhe accumulated charges through the phases. The ME effect is real-zed in composites on the concept of product property [6]: a suitableombination of magnetostrictive and piezoelectric phases can yield

∗ Corresponding author.E-mail address: crisfedor@yahoo.com (C.E. Ciomaga).

925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.jallcom.2009.05.101

© 2009 Elsevier B.V. All rights reserved.

a ME property [7]. If a system is formed by more than one phase,its functional properties result as being derived both by the proper-ties of the constituents and also from their reciprocal interactionsas synergy effects. Three classes of synergy effects were identified[3,8]: (i) sum property, which represents the weighted sum of thecomponents’ contributions of the constituent phases; (ii) combi-nation property denoting an effect in which, at given compositionsor under specific circumstances, the amplitude of the property ishigher in the composite than in the end compounds; (iii) productproperty, referring to effects present in the composites but not in theindividual phases. Whereas the sum and combination propertiesare describing average or enhancement of effects already presentin the parent phases, the product property only refers to novelphenomena originated from the reciprocal interaction betweenthe individual phases. The ME composites normally present func-tional characteristics given as sum property (e.g. magnetization)and as product property (e.g. magnetoelectric effect). The MEeffect in composites has an extrinsic character, depending on thecomposite microstructure and on the coupling interaction acrossmagnetostrictive–piezoelectric interfaces. Composites’ propertiesalso depend on: (a) the number, type and volumic ratio of the com-

ponent phases; (b) the individual component properties; (c) phasesconnectivity.

Inverse spinel cobalt ferrite CoFe2O4 (CF) has been paid a greatdeal of attention for its applications such as high-density magneticand magneto-optic recording media. CF exhibits ferrimagnetism

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template powders. Fig. 1 shows the FT-IR spectra of the combustionsynthesized xCF–(1 − x)PZTN composites. The spectra elucidate theposition of cations in the crystal structure with oxygen ions andtheir vibrational modes, which represents the various orderingpositions of the structural properties. The spectrum of PZT com-

C.E. Ciomaga et al. / Journal of Allo

elow 793 K [9]. The octahedral Co2+ (d7) ions in CF are in the highpin state and the tetrahedral in the octahedral Fe3+ (d5) ions aren the high spin state with the spin directions antiparallel to eachther. It also has good magnetostrictive properties which are depen-ent on composition and preparation method [10–12]. Due to itsroperties, this system was often used as magnetic component inagnetoelectric composites [13–16].

Lead zirconate titanates PbZrxTi1−xO3 (PZT) are ferroelectric per-vskites ABO3 with excellent electromechanical properties, widelysed as ultrasonic resonators, ceramic filters, high-power trans-ucers and actuators [17]. Piezoelectric PZT ceramics having theomposition at the morphotropic phase boundary (x = 0.52) cane doped with ions to form “hard” or “soft” PZTs. Hard PZTs areoped with acceptor ions such as K+, Na+ (for A site) and Fe3+,l3+, Mn3+ (for B site), creating oxygen vacancies in the lattice

18,19]. Soft PZTs are doped with donor ions such as La3+ (for Aite) and Nb5+, Sb5+ (for B site) leading to the creation of A siteacancies in the lattice. They normally have a higher permittiv-ty, larger losses and much higher piezoelectric coupling. Amonghem, Nb-doped PZT (PZTN) was mostly studied [20–22] and dueo its characteristics is commonly used for applications requiringood piezoelectric properties. This system also was chosen in thisork as ferro/piezoelectric phase to form magnetoelectric ceramic

omposite ceramics.From our previous studies concerning the ceramic composite

ystem (Ni,Zn)Fe2O4–BaTiO3 [23,24], it was demonstrated that theethod of in situ synthesis of the composite powders ensures a

etter mesoscale mixing and much more homogeneous microstruc-ures than the simple mixing of the separate powder components,eading to serious reduction of the dielectric losses. An in situol–gel method followed by a conventional sintering was used toroduce ME composites by Wu et al. [25], resulting in homoge-eous microstructures with NiFe2O4 grains well dispersed in a PZTatrix and good functional properties. Thus, it is expected that a

etter control of the microstructure and better connectivity cane achieved by in situ processing of the ferroelectric–magneticomposite. In the present paper, the gel-combustion method wasdopted to produce in situ composites formed by magnetic spineloFe2O4 and ferro/piezoelectric Nb–Pb(Zr,Ti)O3 perovskite (PZTN)nd their functional properties are reported.

. Sample preparation and experimental details

Pb0.988(Zr0.52Ti0.48)0.976Nb0.024O3 (PZTN) (�th = 8.006 g/cm3)emplate powders were firstly prepared by a conventional mixedxides method by using reagent grade PbO (Aldrich 211907, purity9.9%), ZrO2 (MEL SC101), TiO2 (Degussa P25) and Nb2O5 (Aldrich08515, 99.99%). The precursors were ball-milled in water for 48 hith Zr milling media, then freeze-dried, sieved to 250 �m and cal-

ined at 800 ◦C/4 h. The resulted powder was wet-milled in ethanolor 100 h, then dried and again sieved. The mean diameter (d50) ofhe PZTN calcined powder was of ∼0.71 �m and the specific surfacerea of 2.67 m2/g. CoFe2O4 (CF) and xCF–(1 − x)PZTN compositeowders (x = 10, 20, 30 wt.% ) were prepared by citrate–nitrateombustion method, following a similar procedure as reportedn detail elsewhere [13,26]. Stoichiometric amounts of cobalt (II)itrate hexahydrate (Co(NO3)2·6H2O p.a. Merck) and iron (III)itrate nonahydrate (Fe(NO3)3·9H2O, 99.9%, Aldrich) solutionsere mixed. The mixtures PZTN and (nitrate/citric acid solution)ere stirred for 60 min and after that they were heated at 80 ◦C,

ith continuous stirring. After the evaporation of the water excess,highly viscous gel has been obtained. Subsequently, the gelas ignited at 300 ◦C to evolve the undesirable gaseous products,

esulting in the formation of a foamy powder containing theomposite precursor. To complete the ferrite formation onto the

Compounds 485 (2009) 372–378 373

PZTN templates, the powder mixture was calcined at 500 ◦C/8 hresulting in fine composite powders of xCF–(1 − x)PZTN (x = 0.10,0.20 and 0.30), and cold isostatically pressed at 300 MPa. Thesamples were sintered at 1100 ◦C/2 h and at 1250 ◦C/2 h, in Pbatmosphere maintained by a PbZrO3 source in a closed Al2O3crucible [26].

The phase formation of the CoFe2O4 ferrite onto the PZTN tem-plates was checked by using a Fourier transform infrared (FT-IR)spectra, recorded in the range 4000–400 cm−1 with 2 cm−1 reso-lution on a Jasco 660 plus FT-IR spectrophotometer using the KBrpellet technique. X-ray diffraction measurements at room temper-ature for investigating the formation of the perovskite and spinelphases in the composite powders and in the sintered pellets wereperformed with a SHIMADZU XRD 6000 diffractometer using Ni-filtered Cu K� radiation (� = 1.5418 Å), with a scan step of 0.02◦

and a counting time of 1 s/step, for 2� ranging between 20◦ and80◦. The microstructure of the composite powders and of sinteredceramics was examined by using a scanning electron microscope(SEM) coupled with energy dispersive X-ray spectroscopy (EDX)(VEGA/TESCAN instrument). The electrical measurements wereperformed on parallel-plate capacitor configuration, by applyingPd–Ag electrodes on the polished surfaces of the sintered ceramicdisks. The complex impedance in the frequency domain (20 Hzto 2 MHz) at room temperature was determined by using animpedance bridge type Agilent E4980A Precision LCR Meter. Themagnetic properties at room temperature were determined undermagnetic fields up to 14 kOe with a Vibrating Sample MagnetometerMicroMagTM VSM model 3900 (Princeton Measurements Co.).

3. Results and discussions

The infrared spectra of the composite powders were collected inorder to check the formation of the CF spinel phase onto the PZTN

Fig. 1. FT-IR spectra of the xCF–(1 − x)PZTN composites in situ synthesized by gel-combustion.

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74 C.E. Ciomaga et al. / Journal of Allo

ound is considered a superposition of the isolated PbTiO3 andbZrO3 spectra [27]. In the IR spectra of PZTN, in 400–900 cm−1

ange, the asymmetric band centered to 600 cm−1 is assigned to theetal oxide compound, identified by XRD as the PZTN perovskite

hase. The IR absorption band is associated to BO6. Guarany etl. [27] assigned the stretching vibrations of PbZr0.50Ti0.50O3, withhe �1 Ti–O stretching mode at approx. 660 and 615 cm−1 and ther–O stretching modes at approx. 770 and 547 cm−1. According tohese interpretations, for PZTN template the shoulders at approx.69 cm−1 may be associated with the Zr–O stretching mode. In theT-IR spectra of xCF–(1 − x)PZTN composites, we assigned the peakst 729 and 513 cm−1 to the Zr–O stretching mode, in accordanceith small amount of the non-stoichiometric ZrO2−x presented inRD pattern (Fig. 2). The shift of these two maxima explains therdering of the ZrO6 octahedra with increasing the temperature of00–1200 ◦C. The IR measurements in the region 400–2000 cm−1

onfirm the presence of the CoFe2O4 phase. The metal cations in

pinel ferrites are situated in two different sub-lattices namelyetrahedral (A sites) and octahedral (B sites), according to the geo-

etric configuration of the oxygen ion nearest neighbors. The band1 around 600 cm−1 is attributed to the stretching vibration of tetra-edral complexes and �2 around 400 cm−1 to that of octahedral

ig. 2. XRD pattern of the xCF–(1 − x)PZTN composites and of the PZT and CF parenthases: (a) powders with various compositions: x = 0.10, 0.20, 0.30 and (b) x = 0.20omposition as powder and ceramic sintered at 1100 ◦C/2 h and 1250 ◦C/2 h.

Compounds 485 (2009) 372–378

complexes. Due to the stretching vibration of Fe3+–O2− in tetra-hedral sites [28], our Co–ferrite shows only one absorption bandat around 590 cm−1. In Co–ferrite, the absorption maximum at670 cm−1 is due the presence of Co in tetrahedral centers [29]. Theoctahedral complexes may not be present within the experimentalrange, because the band position is nearly 400 cm−1. The IR spec-tra in Fig. 1 do not show this absorption band at around 720 cm−1,suggesting that �-Fe2O3 or/and Fe3O4 was not formed in our com-posite system. These results are in good agreement with the XRDanalysis [30].

The IR measurements in the region 400–2000 cm−1 confirmthe presence of the CoFe2O4 phase; in fact, a large band associ-ated to the Fe–O stretching mode at 590 cm−1 for all composites ispresent [28]. The band at 1640 cm−1 present in CoFe2O4 sinteredat 900 ◦C shows the existence of residues of organic compounds;in the xCF–(1 − x)PZTN after heating treatment at 1100 and 1200 ◦Cthis band disappears, indicating the absence of organic compoundsin synthesized composites. In Co–ferrite, the absorption maximumat 670 cm−1 is due the presence of Co in tetrahedral centers [29].

X-ray diffraction analysis was performed to determine the phasepurity of the xCF–(1 − x)PZTN composites, after the calcination stepand after sintering (Fig. 2(a) and (b)). From Fig. 2(a), it is observedthe most prominent peak (1 0 1) of the tetragonal phase of PZTNat 2� ∼ 30.9◦ (pdf 33-0784). For the ferrite the prominent peak islocated at 2� ∼ 35◦ (1 1 3) and the crystal structure is a rhombo-hedral spinel (pdf 79-1744). The XRD analysis of xCF–(1 − x)PZTNcomposites with x = 0.10, 0.20, 0.30 show that after calcinations at500 ◦C/8 h all the compositions present both perovskite and thespinel phases. It can be seen from the figure that the intensity offerrite peaks increases with increasing the ferrite addition. TheXRD patterns for sample with x = 0.20 sintered at 1100 ◦C/2 h andat 1250 ◦C/2 h, confirmed the successful formation of a di-phasecomposite, formed by the spinel CF and PZTN perovskite phases inthe nominal proportion, both after the calcination step and aftersintering in bulk (Fig. 2(b)). The increasing of the sintering temper-ature to 1250 ◦C aimed to lead to a better densification, resultedin the formation of some small amounts of non-stoichiometriccompounds of Zr1−xO2−y (pdf 81 1548, 81 1329, etc.) as demon-strated by the presence of maxima around 2� ∼ 24◦, ∼28◦ and ∼34◦,besides the main peaks of the parent phases (Fig. 2(b)). This residualphase resulted most probably from the use of the PbZrO3 cover-ing powder during sintering for compensating the Pb-evaporationand its amount was reduced after polishing the ceramic surfaces. Itresults that by sintering in the range of (1100–1250) ◦C, both a gooddensification and a minimization of the secondary phases can berealized. The crystallite sizes as determined by the Debye–Scherrerrelation is increasing when the sintering temperature increasesmainly for the PZTN phase, while for the CF component they remainalmost unchanged (DPZTN = 166 nm and DCF = 338 nm for sinter-ing at 1100 ◦C; DPZTN = 245 nm and DCF = 345 nm after sintering at1250 ◦C). For various compositions of the xCF–(1 − x)PZTN ceramicsintered at the same temperature, the proportion of the two phasesdetermined from the relative peak intensity ratios of the parentphases were very close to the expected nominal ones.

The scanning electron micrographs of 0.10CF–0.90PZTN pow-der after the combustion step are shown in Fig. 3((a) and (b)).Since the composites are multiphase, it is desirable to know thedistribution of the constituent phases which are observed in themicrostructures. The functional properties are expected to be criti-cally dependent on the microstructure. Fig. 3 indicates the existenceof the two distinct phases with different morphologies, one is the

ferrite phase (grey color in backscattering mode: agglomeratedblocks of CF) and the other one is the ferroelectric phase (whitecolor in backscattering mode: PZTN fine particles). Fig. 4 showsthe microstructure of xCF–(1 − x)PZTN composite with x = 0.20 sin-tered at 1250 ◦C in backscattering mode, where the ferrite phase

C.E. Ciomaga et al. / Journal of Alloys and

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ig. 3. SEM micrographs of the in situ prepared composite of 0.10CF–0.90PZTNowder after the combustion step: (a) morphology; (b) backscattering image (bar:00 �m; white: PZTN particles, grey: CF).

ith grain sizes of around 10 �m appears grey and the ferroelectrichase with much finer grain sizes (2–4 �m) appears in light color.rom this picture a very good mixing of the ferroelectric and ferrite

hases, as result of the present in situ method of preparation cane observed.

The dielectric properties of the xCF–(1 − x)PZTN ceramic com-osites (x = 0, 0.10, 0.20 and 0.30) sintered at 1250 ◦C/2 h were

ig. 4. Backscattering SEM micrograph of the sintered composite 0.20CF–0.80PZTNeramic showing the mixing of the dissimilar phases (white: PZTN, grey: CF grains).

Compounds 485 (2009) 372–378 375

checked as a function of frequency in the range of 100 Hz to2 MHz, at room temperature and they are shown in Fig. 5. All theceramic samples have high dielectric constant at room temperature(Fig. 5(a)), of around 2000 for all the investigated frequencies. Thefrequency dependence of the imaginary part of permittivity presenta continuous decrease of permittivity with increasing the frequencyfor the xCF–(1 − x)PZTN composites with x = 0.10, 0.20 and 0.30 atroom temperature (Fig. 5(b)). In the case of the PZTN ceramics, over-lapped onto the general decreasing permittivity with frequency,some jumps with various amplitudes with positive/negative valuesof the real part of permittivity (Fig. 5(a)) accompanied by maximaof the imaginary part of permittivity (Fig. 5(b)) between 104 and106 Hz were observed. The PZTN is a good piezoelectric ceramicand due to this is giving rise to multiple resonances/antiresonances(characterized by positive/negative jumps of the imaginary part ofthe impedance and by maxima of the real part) observed in therange of kHz to MHz. The multiple resonances observed in our casefor PZTN are due to the activation of different piezoelectric vibra-tion modes of the ceramic sample, which comes from the randomlyorientation of the crystallites with respect with the applied field. Asclearly observed (Fig. 5((a) and (b)), the presence of the magneticphase results in the cancellation of the piezoelectric resonances inthe range of frequencies of 100 Hz to MHz or their shift to higher fre-quencies, out of the present measurement range. Further studies athigher frequencies will demonstrate the role of the composition andmicrostructures on the piezoelectric resonance shifts. Similar prop-erties were reported in other composites formed by ferroelectricsolid solution Pb(Ni1/3Nb2/3)O3–PbTiO3 and magnetic ferroxplanaBa2Zn1.2−xCoxCu0.8Fe12O22 [31], in PZT rod array within a Terfenol-D/epoxy matrix composite with a pseudo-(1–3) connectivity [32] orin situ prepared composites of NiFe2O4 and PZT ceramics obtainedby a sol–gel method followed by a conventional sintering [25] andseem to be a characteristic of the composites containing a goodpiezoelectric as ferroelectric phase in composite. Such resonancebehavior is also supported by recent theoretical calculations of theacoustic oscillations induced by the piezoelectric effect producedby the ferroelectric phase in heterogeneous magnetoelectric multi-ferroic composites [33]. Determining the resonance range for theimpedance (permittivity) in magnetoelectric composites is veryimportant since the magnetoelectric effect proved to have maxi-mum values close to the frequency of resonance in such systems[34].

The frequency dependence of tan ı shows that losses inxCF–(1 − x)PZTN ceramic composites (x = 0.10, 0.20 and 0.30) arehigher than in PZTN and are above unity below 1 kHz (Fig. 5(c)),as in many ferroelectric–magnetic composites reported in the lit-erature [3,7,8]. The impedance spectra of the ceramic composites(Fig. 5(d)) are apparently characterized by single semicircular arcs,whose pattern changes with composition, indicating a modifica-tion of the resistance/reactance ratio when increasing the ferriteaddition x. Similar dielectric properties were reported for PZT–Niferrite composites sintered by spark plasma sintering [35] and inPZT–(Ni, Zn) ferrite ceramics prepared by powder-in-sol precursorhybrid processing route [36]. However, the present data show thatlosses strongly reduces with increasing frequency and tan ı < 0.5above 10 kHz and generally, lower losses are obtained for the insitu prepared composites by comparison with ones produced bysimply mixing the component phases.

The apparent increase of the real part of permittivity incomposites at low frequencies is accompanied with high lossescaused by the slow charge species relaxation (e.g. space charge

or Maxwell–Wagner phenomena which are normally thermallyactivated [37]). This type of relaxation mechanism is associatedwith uncompensated surface charges at the dissimilar mate-rial (perovskite–spinel) interfaces inside the composite ceramics.Changes of the resistivity and permittivity at the ceramic grain

376 C.E. Ciomaga et al. / Journal of Alloys and Compounds 485 (2009) 372–378

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ig. 5. Dielectric properties of the xCF–(1 − x)PZTN composites with x = 0, 0.10, 0.20art of permittivity vs. frequency; (c) tangent loss vs. frequency; (d) complex imped

oundaries with respect to the grain bulk are also due toossible doping with foreign ions from the other phase, thatight cause semiconductor character if they are n-type or againaxwell–Wagner mechanisms if they are p-type [38]. Doping can

lso cause changes of the unit cell parameters: in the case of fer-oelectrics this might result in different ferroelectricity than in theure single-phase and thus changing its intrinsic contribution in theomposite. All of these possible mechanisms can give contributionso the dielectric response, besides the intrinsic ones resulting fromhe dielectric/ferroelectric characteristics of the ferroelectric and

agnetic phases. The thermally activated relaxations are normallyctive at higher temperatures in single-phase systems; it seems thaty increasing the degree of inhomogeneity of the system, their acti-ation energy is dramatically lowered and thus they manifest event room temperature. As discussed in detail in Ref. [38], such extrin-ic phenomena cause the peculiar dielectric properties observed inll the ME composites reported in the literature.

The dielectric data of ceramic composites can be compared at= 1 kHz, which is a rather high for the slow charged species and alsos out of the piezoelectric relaxation range. The permittivity valuest this frequency are: 1600, 506, 349 and 210 when x = 0, 10, 20 and

0%, respectively. The reduction of permittivity when reducing theoncentration of the ferroelectric phase (which possess a higherermittivity) is expected as a consequence of the sum property6,8] and can be calculated by using effective field models. Therere many empirical, semi-empirical and highly formalized expres-

.30 at room temperature: (a) real part of permittivity vs. frequency; (b) imaginaryplot. (Note a few piezoelectric resonances of the PZTN ceramic above 2 × 104 Hz.)

sions to predict the dielectric properties in such composites, butthe most common approximations are [39]: (i) Maxwell–Garnett,corresponding to isolated spherical inclusions embedded in a con-tinuous matrix background, valid for small filling concentrationsand (ii) Bruggeman, for which the two phases are percolated andthe properties of matrix or host are the same as the effective fieldmedium itself. The latter approach seems more adequate for ourcomposite ceramics, by considering the ceramic microstructurefrom Fig. 4, where both the ferroelectric and magnetic phases arepercolated. This also explains the high losses (above unity) in ourcomposites, due to the presence of ferrite low-resistivity conduc-tion chains.

The magnetic activity was also checked at room temperature.The magnetic characteristics of the composites follow from theferrite’s properties, which are sensitive to composition, prepa-ration method and sintering conditions. If the presence of thenon-magnetic phase in composite does not change the nature ofthe magnetic interactions (i.e. intrinsic changes of the magneticproperties of the ferrite), a proportional reduction of the magne-tization M (“sum property” [6,8]) with reducing the amount ofthe ferrite phase is expected. The hysteresis loops M(H) at room

temperature (Fig. 6) show the presence of the ordered magneticstructure derived from the unbalanced antiparallel spins as in thepure CoFe2O4 ferrite system. As a consequence of the “sum prop-erty” [6,8] and to possible interface effects, a reduction of thesaturation magnetization from ∼40 emu/g for x = 0.30–∼14 emu/g

C.E. Ciomaga et al. / Journal of Alloys and

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ig. 6. Magnetic hysteresis loops at room temperature obtained for theCF–(1 − x)PZTN composite ceramics with x = 0.10, 0.20 and 0.30 (Note:emu/g = 1 m2A/kg).

or x = 0.10 by comparison with the values reported for pure CoFe2O3akes place [12].

It results that by in situ processing of ceramic composite byol–gel-combustion method we have obtained valuable dielectricnd magnetic properties at room temperature. More studies will beake in order to see the behavior of the xCF–(1 − x)PZTN ceramic

omposites at higher frequencies (1 MHz to 1 GHz).

. Conclusions

Composites formed by CoFe2O4 and Nb–Pb(Zr,Ti)O3agnetic–ferroelectric ceramic with limited reactions at interfacesere prepared by citrate–nitrate combustion technique by usingb–Pb(Zr,Ti)O3 template powders prepared by mixed oxides.ure di-phase powder composites with a good crystallinity werebtained after the calcination step. The purity was maintainedy sintering at temperature of 1100 ◦C/2 h. Increasing the sin-ering temperature to 1250 ◦C/2 h resulted not only in a betterensification, but also in the formation of some small amountsf secondary phases (mainly of ZrO2−x), as demonstrated by theT-IR spectra and by the XRD investigation. The scanning elec-ron micrographs show the existence of the two distinct phases,ne is the ferrite phase and the other is the ferroelectric phase.he dielectric and magnetic investigations at room temperatureonfirmed the formation of a composite with both dielectricnd magnetic properties at room temperature, with permittivitynd magnetization resulted as sum properties from the parentZT and ferrite phases. High losses characteristic at frequencieselow 103 Hz are most probably due to Maxwell–Wagner effect asesulted from the noncompensated charges at the ferrite–(Nb–PZT)nterfaces.

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[

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