BaGeo3 as Sintering Additive for BaTiO3–MgFe2O4 Composite Ceramics

BaTiO3–MgFe2O4 composites (30 wt% MgFe2O4) with a small addition of BaGeO3 as a sintering additive were synthesized by a one-pot Pechini-like sol–gel process. Nano-crystalline composite powders with a crystallite size of about 10 nm were obtained after reaction at 700 C for 1 h. Magnetic investigations suggest that the nano-powder is in its superparamagnetic state at room temperature. The addition of BaGeO3 leads to an improved sintering behaviour. DTA measurements reveal the formation of a liquid phase at 1164(3) &#176;C. Dense ceramic bodies (relative density $ 90%) were obtained after sintering for 1 h at 1150 &#176;C. SEM investigations prove a 0–3 connectivity and show that the addition of BaGeO3 promotes the grain growth leading to particles up to 4 mm. In contrast, fine-grained composite ceramics with smaller particles up to 230 nm were obtained after a two-step sintering process. Magnetic measurements indicate a ferrimagnetic behaviour with coercivity values up to 70 Oe depending on the sintering procedure. Furthermore, addition of BaGeO3 results in an increase of the relative permittivity, whereas the dissipation factor slightly decreases.


Introduction
Multiferroic composite materials are of high interest because of their potential applications in advanced technology e.g. as sensors, memories and actuators. [1][2][3] Composites consisting of both ferro/ferrimagnetic and ferroelectric phases can exhibit a magneto-electrical (ME) coupling effect. 4 To produce BaTiO 3 -MFe 2 O 4 composites (e.g. M ¼ Co 2+ , Ni 2+ , Mg 2+ ) various synthesis routes have been reported, e.g. mixed-oxide-, sol-gel-and hydrothermal processes. [5][6][7][8] Reactions between the BaTiO 3 and MFe 2 O 4 phase oen lead to the formation of BaFe 12 O 19 and/or hexagonal BaTiO 3 . 9,10 Therefore, the sintering temperature for composite ceramics is limited and depends on the synthesis route and particle size. For some ferrite components, the BaTiO 3 -MFe 2 O 4 composites show an insufficient sintering behaviour. [11][12][13] In particular BaTiO 3 -MgFe 2 O 4 ceramic bodies exhibit low densities. Up to now this system has rarely been investigated. 14-17 A magneto-electrical effect in BaTiO 3 -MgFe 2 O 4 composites was found by Tan et al. 14 and Tadi et al. 15 As the ME coupling strongly depends on the quality of the oxide interface, dense ceramics with a close contact between the perovskite and spinel phase are advantageous for large ME coefficients. Sintering additives improve the densication, therefore the densities of the ceramic bodies increase and/or the sintering temperature can be reduced. [18][19][20] Low sintering temperatures are of special interest as they enable to obtain ne-grained ceramics.
Furthermore, ne-grained ceramics show improved mechanical properties. 21 The inuence of sintering additives on BaTiO 3 -MFe 2 O 4 ceramics has barely been investigated to date. To our best knowledge, there are only reports on boron-and bismuthcontaining uxes as sintering additives leading to BaTiO 3 -MFe 2 O 4 glass-ceramic composites. [22][23][24] Li et al. 25  The formation of a molten phase during sintering usually promotes grain growth and should therefore be avoided if small particle sizes are desired. Previous publications showed that the formation of a molten ux is not necessary to improve the densication of ceramic bodies when sintering additives on the basis of e.g. BaGeO 3 and SiO 2 are used. [19][20][21][22][23][24][25][26][27] The aim of this paper is to investigate the inuence of BaGeO 3 as sintering additive on BaTiO 3 -MgFe 2 O 4 composite ceramics with a 0-3 connectivity. The samples were synthesized using a one-pot sol-gel method leading to nano-sized powders. Phase formation during the sintering process was monitored by XRD and thermal analysis. The sintering behaviour and microstructure of resulting compacts were determined by dilatometry and SEM. Magnetic measurements were carried out both on composite powders and ceramic bodies. Furthermore, the inuence of BaGeO 3 on the dielectric behaviour was studied.
elsewhere. 17 This composition is abbreviated as BT 0.95 G 0.05 -MF-0.3 in the following. Additionally composites with the composition 0.7Â (BaTi 0. 9  The other compositions were prepared by the same procedure using corresponding stoichiometric quantities. To obtain ceramic bodies, the calcined powders were uniaxially pressed at about 85 MPa into pellets (green density 2.0 g cm À3 ) without using any pressing aid or binder. These pellets were placed on a ZrO 2 bre mat. Two different sintering procedures were used; conventional sintering (heating up with 10 K min À1 , soaking at this temperature, cooling down with 10 K min À1 ) and two-step sintering (heating up with 20 K min À1 to 1150 C (T 1 ), then fast cooled with 30 K min À1 to a lower temperature T 2 and held at T 2 for 20 h).

Characterization
X-ray powder diffraction patterns were collected at room temperature on a Bruker D8-Advance diffractometer, equipped with a onedimensional silicon strip detector (LynxEye™) and operating with Cu-K a radiation. Crystallite sizes were determined from the XRD line broadening using the Scherrer equation and the integral peak breadth (soware suite WinXPOW 28 ). Dilatometric investigations were performed in owing synthetic air (50 ml min À1 ) in a Setaram TMA 92-16.18 dilatometer. Differential thermoanalytic (DTA) measurements in owing synthetic air (20 ml min À1 ) were done using a Netzsch STA 449 system. Scanning electron microscope images were recorded with a Philips XL30 ESEM in the backscattered electron mode (BSE). For magnetic measurements a Quantum Design PPMS 9 was used. Hysteresis loops were obtained with magnetic eld cycling between À90 and +90 kOe. In addition, the temperature dependent magnetic moments were measured in the temperature range of 5-300 K under eld-cooled (FC) and zero-eld-cooled (ZFC) conditions. An Impedance Analyzer 4192A (Hewlett Packard) was used for permittivity measurements up to 13 MHz. A eutectic Ga-In alloy was applied as electrical contact on both sides of the ceramic bodies.

Powder characterization
In a previous paper 17 we reported on the preparation of (1 À x) BaTiO 3 -(x)MgFe 2 O 4 composites (x ¼ weight fraction) with 0-3 connectivity via a Pechini-like process. The BaTiO 3 -MgFe 2 O 4 composite powders were found to have an insufficient sintering behaviour resulting in ceramic bodies with moderate densities of about 70% at 1200 C. The maximum sintering temperature for BaTiO 3 -MFe 2 O 4 systems is limited because of the formation of hexagonal BaTiO 3 and/or BaFe 12 O 19 . 9,17,29 As stated in ref. 30 and 31 the densication of BaTiO 3 can be improved using BaGeO 3 as sintering additive. Guha and Kolar 32 found that a maximum of 1.8 mol% Ge 4+ (related to the occupation of the Ti 4+ site) can be incorporated into the BaTiO 3 lattice. For our investigations a much higher content of $5 mol% BaGeO 3 was used, resulting predominantly in a mixture of (Ge-doped) BaTiO 3 and BaGeO 3 . For simplicity this composition is described by the formula BaTi 0.95 Ge 0.05 O 3 in the following.
The rst step of the Pechini-like sol-gel process leads to highly-viscous homogeneous gels in which the metal ions are mixed on a molecular level. To obtain composite powders, these dark-brown gels were calcined at 700 C for 1 h. The XRD pattern of the resulting powder shows mainly reections of BaTiO 3 , MgFe 2 O 4 , and only small amounts of BaCO 3 (ref. 33) ( Fig. 1). Aer this rst heat treatment no peaks related to Ge-containing phases could be detected. As an example, the powder 0.7Â (BaTi 0.95 Ge 0.05 O 3 )-0.3Â (MgFe 2 O 4 ) (donated as BT 0.95 G 0.05 -MF-0.3) was examined by the BET method and was found to possess a specic surface area of 34.5 m 2 g À1 , corresponding to an equivalent particle size of 32 nm. The volume-weighted average crystallite sizes for both the MgFe 2 O 4 and the BaTiO 3 phase were calculated as 10 nm from XRD line broadening. The calculated crystallite size is roughly by the factor three smaller than the size of the primary particles obtained from the BET data. This is most likely because the decomposition of the gel leads to closely joined crystallites and in turn surface areas unavailable for nitrogen adsorption. 34 An alternative explanation is that each primary particle may consist of several crystallite domains.
Temperature-dependent magnetic investigations (H ¼ 500 Oe) show a superparamagnetic behaviour for the BT 0.95 G 0.05 -MF-0.3 powder ( Fig. 2A). The eld-cooled (FC) curve increases continuously with decreasing temperature, while the zero-eld-curve (ZFC) curve has a maximum at 20 K, the so called blocking temperature. The superparamagnetic behaviour is conrmed by eld-dependent magnetic measurements. At 300 K the coercivity value is close to zero, whereas below the blocking temperature the hysteresis loop at 5 K shows a coercivity value of 146 Oe (Fig. 2B).

Composite ceramics
DTA investigations (shown in Fig. 3) of BT 0.95 G 0.05 -MF-0.3 and BT 0.9 G 0.1 -MF-0.3 samples show the formation of a liquid phase (Graph 3a,b). The DTA signal increases with rising BaGeO 3 content. The onset temperature of the melting peak was determined as 1164 (3) C. In contrast, the BT-MF-0.3 composite without BaGeO 3 does not reveal any DTA signal up to 1200 C (Graph 3c). Fig. 4 shows the non-isothermal dilatometric measurement up to 1200 C in owing air of compacts from powder BT 0.95 G 0.05 -MF-0.3 and BT-MF-0.3 calcined at 700 C for 1 h, respectively. The shrinkage process of BT 0.95 G 0.05 -MF-0.3 (Graph 4a) starts at about 800 C and the shrinkage rate achieves a maximum at 1102 C with a value of À1.68% min À1 . Interestingly, this shrinkage maximum occurs below the liquid phase formation at 1164 C. The shrinkage rate indicates that the densication is dominated by sliding processes (viscous ow). Diffusion as the dominant process causes only shrinkage rates of about 10 À4 to 10 À1 % min À1 . [35][36][37][38] Investigations on the BaTiO 3 -BaGeO 3 system conrm that the addition of BaGeO 3 promotes sliding processes. 39 The shrinkage has not yet been nished at the highest investigated temperature of 1200 C. For comparison graph 4b shows the shrinkage behaviour of BT-MF-0.3 compacts without the addition of BaGeO 3 . A maximum of the shrinkage rate can be observed at 1185 C (À0.80% min À1 ). Thus the addition of BaGeO 3 leads to improved shrinkage behaviour at relatively low temperatures without involvement of a molten phase. Fig. 5 shows the nal bulk densities of ceramic bodies aer isothermal conventional sintering at various temperatures in static air (heating up with 10 K min À1 , soaking at this temperature for 1 h, cooling down with 10 K min À1 ). The absolute bulk densities of the sintered bodies were determined from their   weight and geometric dimensions. The relative bulk densities of all ceramic bodies were calculated with respect to the theoretical density of 5.39 g cm À3 for BT-MF-0.3. 17,31 It turned out that BT 0.95 G 0.05 -MF-0.3 ceramic bodies with relative densities of 70% can be obtained even at 1050 C. By raising the sintering temperature to 1150 C dense ceramics with 91% relative density can be achieved. Aer sintering at 1200 C the relative density further increases to 95%. According to the DTA results sintering above 1150 C is characterized by liquid-phase sintering processes. As seen in Fig. 5 the formation of a liquid-phase is not necessary to obtain dense ceramic bodies.
Longer sintering times have little inuence on the density. For example, a soaking time of 10 h at 1100 C results only in minor increase of the relative density from 76 to 79% for BT 0.95 G 0.05 -MF-0.3 composites. Furthermore, an increase of the BaGeO 3 content to 10 mol% (BT 0.9 G 0.1 -MF-0.3) does not lead to rising densities. Aer sintering at 1100 C for 1 h ceramic bodies of BT 0.9 G 0.1 -MF-0.3 with a relative density of 69% (3.71 g cm À3 ) were obtained. On the other hand, BaGeO 3 -free BT-MF-0.3 ceramics reach relative densities of only 50 and 70% aer sintering at 1100 and 1200 C for 1 h, respectively.
The SEM images of BT 0.95 G 0.05 -MF-0.3 ceramic bodies depicted in Fig. 6 show a 0-3 connectivity between the spinel and the perovskite phase. In the applied BSE mode the light grains correspond to BaTiO 3 while the dark grains are MgFe 2 O 4 . Sintering at 1000 C results in grain sizes between about 90 and 200 nm, which only slightly increase to 140-350 nm aer sintering at 1100 C (Fig. 6a). Up to this temperature there is no signicant size difference between the BaTiO 3 and MgFe 2 O 4 grains. In contrast, sintering at 1150 C for 1 h leads to a considerable grain growth with BaTiO 3 grains in the range between 0.25-0.9 mm and MgFe 2 O 4 grains between 0.25 and 3 mm as seen in Fig. 6b. Abnormal grain growth occurs during sintering at 1200 C resulting in BaTiO 3 grain sizes between 0.5 and 1.5 mm while the size of the ferrite grains ranges between 0.5 and 4 mm. The appearance of white rods with dimensions up to 23 Â 4 mm aer sintering at 1200 C indicates the formation of hexagonal BaTiO 3 (ref. 40) (Fig. 6c) in accordance with the XRD results discussed below.
Using a two-step sintering procedure the densication at lower temperatures can be improved. 30,41 In this approach, the compacts were rst heated rapidly (20 K min À1 ) to 1150 C (T 1 ), then fast cooled (30 K min À1 ) to a lower temperature T 2 and held at T 2 for 20 h (see inset in Fig. 5). Two-step sintering with T 2 ¼ 900 C leads to a relative bulk density of 78% (4.23 g cm À3 ), whereas upon rising T 2 to 1000 C the ceramic bodies achieve a relative density of 83% (4.46 g cm À3 ). In both cases the ceramic bodies consists of grain between about 90 and 230 nm. For comparison, conventional sintering at 1000 C for 20 h leads only to a relative density of 68%.
XRD powder patterns of selected BT 0.95 G 0.05 -MF-0.3 ceramics are shown in Fig. 7. The pattern of ceramics sintered  up to 1100 C shows only reections of BaTiO 3 and MgFe 2 O 4 . At 2q z 45 no clear splitting of the BaTiO 3 (002)/(200) peak could be detected indicating a reduced tetragonality, i.e. a considerable lower c/a ratio compared to coarse-grained ceramics with more than 1 mm particles. Aer sintering at 1150 C very small traces of BaFe 12 O 19 appear, 33 whereas aer 1200 C about 19 mol% of BaTiO 3 has been converted to the hexagonal modication. Sintering at 1250 C leads to an increase of hexagonal BaTiO 3 to 54 mol%. The formation of the hexagonal BaTiO 3 modication at such comparably low temperatures is noteworthy since in pure BaTiO 3 samples the transition to the hexagonal structure occurs above 1460 C. 42 The formation of hexagonal BaTiO 3 at lower temperatures in the composites is probably caused by the incorporation of small amounts of iron-and magnesium-ions into the BaTiO 3 lattice. [43][44][45][46][47] Peaks corresponding to BaGeO 3 were not found, most likely because of its small content. In contrast, BT 0.9 G 0.1 -MF-0.3 ceramic bodies with their larger BaGeO 3 content clearly show the formation of monoclinic a-BaGeO 3 (ref. 33) (Fig. 8A, Graph a).
To suppress the formation of hexagonal BaTiO 3 an excess of 5 mol% Ti, with respect to the Ba-content, was used in the BT 1.0 G 0.05 -MF-0.3 composite ceramics. As seen in Fig. 8A (Graph b) aer sintering at 1200 C no reections of hexagonal BaTiO 3 are detectable. Whereas, aer sintering at 1250 C of BT 1.0 G 0.05 -MF-0.3 the fraction of hexagonal BaTiO 3 is not completely supressed but reduced from 54 mol% to 31 mol% (Fig. 8A, Graph c). Therefore, at higher sintering temperatures an excess of more than 5 mol% Ti is needed to completely avoid the formation of hexagonal BaTiO 3 . On the other hand, the titanium excess leads to an increase of BaFe 12 O 19 as seen in Fig. 8B Fig. 9B. The coercivity values show a maximum at 1100 C, whereas the M s only slightly decreases for samples sintered between 1000 C and 1100 C. Sintering at higher temperatures leads to a stronger decrease of M s . The reduction of M s is probably due to the reaction between BaTiO 3 and MgFe 2 O 4 . On the other hand cation redistribution between Mg 2+ and Fe 3+ and thus a change of the inversion parameter is also possible because a similar behaviour was observed in pure MgFe 2 O 4 samples, too. 50 Two-step sintering with T 2 ¼ 1000 C and 900 C leads to similar hysteresis loops at 300 K, which are characterized by M s ¼ 9.1 emu g À1 and H c ¼ 68 and 70 Oe, respectively (not shown).

Magnetic and dielectric measurements of ceramic bodies
Temperature-dependent magnetic measurements with an applied eld of H ¼ 60 kOe of a BT 0.95 G 0.05 -MF-0.3 ceramic  sintered at 1000 C show identical values under FC and ZFC conditions, therefore in Fig. 10 only the FC curve is presented. A slight increase at lowest temperatures was found in contrast to a pure MgFe 2 O 4 ceramic sample indicating a minor paramagnetic contribution in the composites. This small paramagnetic moment is probably caused by diffusion of Fe-ions at the grain boundaries from the MgFe 2 O 4 to the BaTiO 3 phase. [51][52][53] Room-temperature dielectric measurements between 1 kHz and 13 MHz for BT 0.95 G 0.05 -MF-0.3 ceramics sintered under various conditions are shown in Fig. 11. All composite ceramics show slightly decreasing trends in the relative permittivity (3 r ) with increasing frequency (Fig. 11A) indicating dielectric dispersion due to the Maxwell-Wagner interface polarisation. 54,55 The permittivities of conventionally sintered samples increase with rising sintering temperature because of increasing densities and grain-sizes of the ceramics. It is known that the permittivity depends on the porosity and thus on the density of ceramic bodies because of pore-charging effects. 56,57 The dissipation factor (tan d) decreases for samples sintered at #1100 C, whereas sintering at 1150 C and 1200 C leads to an increase of tan d (Fig. 11B). The two-step sintered ceramic bodies show the lowest dissipation factors and the relative permittivities vary between 238-224 (T 2 ¼ 1000 C) and 201-184 (T 2 ¼ 900 C), respectively. The decrease of the dissipation   factor and increase of 3 r of the two-step sintered ceramics, compared to the conventional sintered ones, is most likely due to the improved density and a decrease of lattice defects because of the long soaking time at low temperatures. 58,59 Dielectric measurements on ceramic bodies of BT 0.95 G 0.05 -MF-0.3 and BT 1.0 G 0.05 -MF-0.3 sintered at 1200 C (see insets in Fig. 11A and B) show that the titanium excess, needed for the suppression of hexagonal BaTiO 3 , only slightly inuences the dielectric properties resulting in slightly higher 3 r and lower tan d values. Fig. 12 demonstrates dielectric measurements of both BT 0.95 G 0.05 -MF-0.3 and BT-MF-0.3 (i.e. otherwise identical samples with and without BaGeO 3 as sintering additive) composite ceramics aer conventional sintering. Considerably higher permittivity values are obtained for BT 0.95 G 0.05 -MF-0.3 ceramics primarily due to their higher densities caused by the addition of BaGeO 3 . 60 Up to a sintering temperature of 1150 C the tan d values of BT 0.95 G 0.05 -MF-0.3 are signicantly lower than for the BT-MF-0.3 ceramics (inset in Fig. 12). The decrease of tan d, resulting from the BaGeO 3 addition, is probably related to the increase in grain boundary resistivity. 61,62 The development of the relative permittivity and the dissipation factor between À5 C and 180 C at a frequency of 1 kHz for selected BT 0.95 G 0.05 -MF-0.3 ceramics was additionally studied. Fig. 13A shows ceramics sintered up to 1100 C, whereas samples sintered at 1150 and 1200 C are represented in Fig. 13B. The relative permittivities of all samples considerably increase roughly above 50-70 C and reach a maximum at about 170 C. Similar temperature dependencies were found in various BaTiO 3 -ferrite ceramics. 63-69 A clear ferroelectric/ paraelectric phase transition of the BaTiO 3 phase was not observed. The ceramic bodies consist of BaTiO 3 grains less than 1.5 mm, which are known to lead to a reduction of the permittivity and to a attening of the phase transition peak. 70,71 Furthermore, reactions between BaTiO 3 and MgFe 2 O 4 at the grain boundaries lead to impurities and defects resulting in a lower tetragonality. Furthermore, it was shown elsewhere 72 that the ferroelectric/paraelectric phase transition peak continuously decreases and disappears with increasing ferrite content. Investigations by Sakanas et al. 55 and Gupta and Chatterjee 73 revealed that the phase transition peaks of BaTiO 3 are more visible at high frequencies because of the space charge polarisation in the composites. In fact our temperature depended dielectric measurements at a higher frequency of 1 MHz showed for a sample sintered at 1100 C an anomaly between roughly 120 and 130 C (inset in Fig. 13A) which can be assigned to a ferroelectric/paraelectric phase transition. 73

Conclusions
The effect of the sintering additive BaGeO 3 on MgFe 2 O 4 -BaTiO 3 composites was investigated. Composite samples with a composition of 0.7Â (BaTi 0.95 Ge 0.5 O 3 )-0.3Â (MgFe 2 O 4 ) were prepared by a one-pot Pechini-like synthesis. Calcination at 700 C leads to a nano-sized powder showing a superparamagnetic behaviour. DTA investigations reveal the formation of a liquid phase with an onset temperature of 1164(3) C. The addition of BaGeO 3 clearly improved the densication process of compacted powders. Conventional sintering between 1000 and 1200 C results in ceramic bodies with relative densities of 60-96%. A two-step sintering process at 1000 and  900 C with a prolonged soaking time of 20 h results in relative bulk densities of 83 and 78% with grains between about 90 and 230 nm. Small traces of BaFe 12 O 19 appear aer sintering at 1150 C and the partial formation of hexagonal BaTiO 3 is observed at 1200 C. Magnetic measurements reveal typical ferrimagnetic behaviour with coercivity values up to 70 Oe depending on the sintering procedure. Dielectric measurements show that the addition of BaGeO 3 leads to an increase of the relative permittivity, whereas the dissipation factor slightly decreases. Additionally, it is shown that the sintering regime also inuences the dielectric behaviour. Sintering at 1150 C leads to dense and coarse-grained ceramic bodies possessing high 3 r and moderate tan d values over a wide temperature range. On the other hand, ne-grained ceramic bodies with high densities can be obtained aer a two-step sintering process showing relative permittivities up to 238 at room-temperature and the lowest tan d values. The resulting composite ceramics are potential candidates for ME applications.