Crystal Structure and Related Properties of Manganese-Doped BariumTitanate Ceramics

Hans Theo Langhammer,† Thomas Mu¨ller,‡ Karl-Heinz Felgner,† and Hans-Peter Abicht‡

Departments of Physics and Chemistry, Martin-Luther-Universita¨t Halle-Wittenberg, D-06099 Halle, Germany

The influence of manganese on the crystallographic phase andthe microstructure of BaTi12xMnxO3 ceramics (0< x < 0.05)is investigated. The crystal structure at room temperaturechanges from tetragonal to hexagonal betweenx 5 0.005 and0.017, which causes a drastic change in the microstructure.The Jahn–Teller distortion caused by the MnTi

31 ions is pro-posed as the driving force for the phase transition. Annealingof the as-fired samples in both a reducing and oxidizingatmosphere restores the tetragonal phase, which is accompa-nied by a change in the microstructure based on the percent-age of tetragonal phase.

I. Introduction

MANGANESE-DOPED BaTiO3 ceramics are used for several pur-poses. The main applications are capacitors and devices

using the positive temperature coefficient of resistance (PTCR),where, as an effective acceptor dopant, manganese influences boththe electrical properties and the microstructure, because it is aneffective acceptor dopant. In capacitor materials, manganese im-proves resistivity to reducing atmospheres during sintering. Indonor-doped PTCR ceramics, manganese depresses the grain sizeand enhances the rise in electrical resistivity near the Curietemperature of BaTiO3. Much work has been done to elucidate therole of manganese regarding these effects.1–7 However, little isknown about the influence of manganese on the crystallographicstructure and the microstructure of otherwise undoped BaTiO3.

Glaister and Kay8 first reported in 1960 that certain amounts ofmanganese are sufficient to stabilize the hexagonal phase ofBaTiO3 (h-BT) at room temperature, whereas the hexagonalpolymorph of undoped material is stable in air only at temperatures.1460°C.9,10 Two possibilities are known to stabilizeh-BT atlower temperatures. The first one is firing in reducing atmo-spheres.8,11–16 In pure hydrogen, a temperature of 1330°C issufficient to stabilize the hexagonal phase.8 The formation ofoxygen vacancies and the resulting reduction of Ti41 to Ti31

according to the reaction

OO3 1 2TiTi

3 º VOzz 1 2Ti9Ti 1 1

2O21 (1)

seems to be the generally accepted reason for this stabilizationeffect, which is explained by the preferred formation of face-sharing octahedra (Ti2O9 groups), because of the reduced electro-static repulsion between Ti31 and Ti41/Ti31 ions compared withthe interaction between Ti41 ions. According to Wakamatsu,Takeuchi, and co-workers,11,12a minimum of 0.3 mol% of Ti31 is

sufficient to stabilizeh-BT at room temperature. Rec˘nik andKolar14–16also investigated the microstructure ofh-BT sintered at1360°C in a H2/Ar atmosphere and found strongly exaggeratedgrain growth, which leads to large platelike hexagonal grains.

Doping with 3d transition elements, such as manganese, iron, ornickel, is the second way of stabilizingh-BT at room tempera-ture.8,17–22Because of their small ionic radii, they substitute forTi41, which is commonly accepted in the literature. Their valencestate when they are incorporated into the lattice has been discussedin numerous experimental and theoretical papers. In the case ofmanganese,2,24–27 Glaister and Kay8 found that it is the mosteffective dopant for stabilizingh-BT. Dicksonet al.,17 and recentlyTakeuchiet al.,12 proposed a chemical bonding between the twodorbitals of titanium and manganese as a reason for stabilization ofthe face-sharing TiO6/MnO6 octahedra inh-BT. Renet al.19 statedthat the small size of the Mn41 ion, which substitutes Ti41, is adominant factor favoring the formation ofh-BT. A critical discus-sion of this assumption was published by the authors of the presentpaper.20 Similar to h-BT produced by sintering in reducingatmosphere, manganese-dopedh-BT also exhibits a pronouncedexaggerated growth of platelike grains.20,21

Comparing the published results, the whole picture ofmanganese-doped BaTiO3 is not quite clear. This is true for boththe microscopic models describing the phase transition fromcubic/tetragonal to hexagonal and also macroscopic properties,such as quantitative phase composition, depending on the dopinglevel and the parameters of the sintering process or the microstruc-ture formed. One reason for the inconsistency of the macroscopicresults seems to be the strong influence of impurities of the rawmaterials on the phase composition and microstructure. As anexample, the influence of strontium as an impurity is discussed ina subsequent paper.28

Hence, the aim of this paper is to clarify the influence ofmanganese on the macroscopic properties of BaTiO3 ceramicsmentioned above (crystallographic phase, microstructure) as wellas discuss the microscopic mechanism of the phase transition fromcubic to hexagonal and the role of manganese in the promotion ofthis process. For this purpose, we performed detailed investiga-tions of manganese-doped BaTiO3 at a doping level of between 0and 5 mol% with regard to the following properties: X-raydiffraction (XRD), electron diffraction, microstructure, and thesolubility of manganese in the BaTiO3 lattice. Both as-sinteredspecimens and samples annealed in oxidizing and reducing atmo-spheres were investigated.

II. Experimental Procedure

Ceramic powder with a nominal composition of BaTi1-xMnxO3

(0 # x # 0.05) was prepared by the conventional mixed-oxidepowder technique. After mixing (agate balls, water) and calcining(1100°C, 2 h) appropriate amounts of BaCO3 (,0.1 mol%strontium; No. 3018, Leuchtstoffwerk Breitungen GmbH, Germa-ny), TiO2, (No. 808, Merck, Darmstadt, Germany), and MnCO3

(Riedel de Haen, Germany), we fine-milled (agate balls, water)and densified the powder to disks with a diameter of 12 mm anda height of nearly 3 mm. Because of the thermodynamic instabilityof BaTiO3 in water, the Ba/Ti ratio of the green bodies amounted

B. M. Kulwicki—contributing editor

Manuscript No. 190362. Received February 20, 1998; approved July 23, 1999.Supported by the Kultusministerium des Landes Sachsen-Anhalt, Germany, and

the Deutsche Forschungsgemeinschaft under Contract No. FKZ: 1434A/0083.†Department of Physics.‡Department of Chemistry.

J. Am. Ceram. Soc.,83 [3] 605–11 (2000)

605

journal

to ;1/1.015.29 The samples were sintered in air atTs 5 1400°C for1 h and subsequently characterized with respect to their micro-structure and phase composition (see below). Then, a portion ofeach sample was annealed in air and another portion was annealedin a gas stream mixture of H2 and argon (1/1), and characterizedagain. In both cases, annealing was accomplished at 1200°C in 2 h.All heat treatments took place with heating and cooling rates of 10K/min. To avoid interfering contamination, the samples werecontained in ZrO2-covered Al2O3 dishes.

The microstructure of polished and chemically etched speci-mens was examined by optical microscopy and by scanningelectron microscopy (SEM). To determine the distribution ofmanganese in the grains and intergranular regions, wavelength-dispersive X-ray electron probe microanalysis (WDX-EPMA)(Model CAMEBAX, Cameca Instruments, Corbevoie, France)was performed. For the phase investigations at room temperatureby XRD, the sintered samples were crushed again and mixed withsilicon powder for the purpose of calibration. The phase compo-sition was determined quantitatively by analyzing the intensityratios (111)tetragonal/(103)hexagonaland (200)tetragonal/(103)hexagonal

using a diffractometer (Model D5000, Siemens Aktiengesellschaft,Karlsruhe, Germany). Samples for transmission electron micros-copy (TEM) and electron diffractometry (Model CM20Twin,Philips, Eindhoven, The Netherlands) were thinned by consecutivemechanical grinding, dimpling, and ion milling up to perforationof their central region.

III. Results

(1) X-Ray DiffractionThe results of the XRD investigations both for the as-fired and

annealed samples are shown in Fig. 1. Only the percentage oftetragonal phase as a function of the doping level of manganese isshown. If no other statement is given, the remaining percentagecorresponds to the hexagonal phase.

(A) As-Fired Samples: The region of coexistence of thetetragonal and hexagonal phases in the as-fired samples liesbetween 0.5 and 1.7 mol% manganese. Below and above thisregion, the samples are single-phase within the accuracy of theXRD method, which is not better than;2–5 mol%.

(B) Annealed Samples: The annealed samples show a signif-icant effect of restoration of the tetragonal phase. Although thisholds for both types of annealing atmospheres, the effect is clearlymore pronounced for reducing conditions. During annealing in ahighly reducing atmosphere, some manganese is expelled from thespecimens. Therefore, corrected data giving the remaining man-ganese concentration (see below) are also shown in Fig. 1. The

most remarkable finding is the nearly complete restoration of thetetragonal phase after annealing in reducing atmosphere at amanganese concentration of;1.6 mol%.

(2) Microstructure(A) As-Fired Samples: Figures 2(a)–(e) show microstruc-

tures of samples with a nominal manganese content of between 0.5and 5.0 mol%. The average grain sizes of all samples investigatedare shown in Fig. 3. Up to a manganese content of 0.5 mol%, themicrostructure looks like the well-known microstructure of un-doped liquid-phase-sintered BaTiO3 ceramics. Simultaneouslywith the appearance of the hexagonal phase at;1 mol% manga-nese, the microstructure becomes bimodal. A new grain type witha platelike shape and a grain size of#400mm develops, while thegrains of the other fraction, which exhibit the globular shape of themanganese-free ceramics, become smaller, and their percentagedecreases. At 1.6 mol% manganese, the smaller globular grainshave vanished, corresponding approximately to the disappearanceof the tetragonal phase. For nominal manganese concentrations of$2.5 mol%, once again, the microstructure becomes bimodal.With increasing manganese content, the globular grains (diameter:5–10mm) reappear, the percentage of which exceeds the percent-age of the platelike fraction at a nominal manganese content of 5mol%. Only for manganese concentrations$3 mol%, when theplatelike grains are surrounded by a sufficient amount of smallglobular grains, is their platelike shape clearly developed. At lowermanganese concentrations, these grains have grown into each otherduring the final stage of sintering and their platelike shape isdeformed.

(B) Annealed Samples: Whereas the samples annealed in airshow no remarkable change in microstructure compared with theas-sintered specimens, the samples annealed in a highly reducingatmosphere exhibit a significant reduction in grain size and achange in the grain shape, which is especially pronounced forsamples with manganese concentrations between 1.6 and 2.0mol%, according to their considerable percentage of tetragonalphase (see Fig. 1). Corresponding to Figs. 2(a)–(e), Figs. 2(f)–(j)show the microstructures of the hydrogen-annealed samples, theaverage grain sizes of which are also shown in Fig. 3.

(3) Electron DiffractionSamples were characterized by electron diffraction, because it is

not easy to distinguish between the hexagonal and cubic phase ofBaTiO3 (c-BT) by XRD measurements. An advantage of themethod is also the possibility of local measurements. Kolaret al.16

reported on small cubic grains in undopedh-BT sintered in areducing atmosphere. Because of the similarities to the microstruc-ture of the manganese-doped samples, mainly small grains wereinvestigated by electron diffraction. None of the numerous grainsinvestigated exhibited the cubic phase. Figure 4 shows the TEMimage of some small grains of a sample with 5 mol% manganesetogether with a typical hexagonal electron diffraction pattern ofone of these grains. Hence, probably all the samples investigatedwith a manganese content of$1.7 mol% seem to be completely inthe hexagonal phase.

(4) Solubility of Manganese(A) As-Fired Samples: The results of our investigation into

the solubility of manganese in BaTiO3 are shown in Fig. 5.Thesmall deviation of the integrally measured values from the nominalmanganese content seems to be due to an uncertain correction ofatomic number effect, absorption, and fluorescence excitation(ZAF) at higher manganese concentrations, because we usedmetallic manganese as a standard sample. While the nominalamount of manganese is completely incorporated up to 1.5 mol%,for higher manganese concentrations, an increasing number ofmanganese ions are not soluble in the BaTiO3 lattice of theplatelike grains but segregate at grain boundaries or form inter-granular manganese-rich phases that are clearly visible in theEPMA-measured distribution images of manganese. Identificationof these phases, the manganese content of which partially exceeds7.5 at.%, has not been successful thus far. Two different phases

Fig. 1. XRD-determined percentage of tetragonal phase (20°C) as afunction of the nominal manganese content (x) of BaTi1-xMnxO3 beforeand after annealing under both oxidizing and reducing conditions.

606 Journal of the American Ceramic Society—Langhammer et al. Vol. 83, No. 3

Fig. 2. Optical micrographs of the microstructure of BaTi1-xMnxO3 ceramics of both (a)–(e) as-fired and(f)–(j) H2-annealed samples with (a), (f)x 5 0.005; (b), (g) 0.010; (c), (h) 0.016; (d), (i) 0.030; and (e),(j) 0.050. Note different magnifications.

March 2000 Crystal Structure and Related Properties of Manganese-Doped Barium Titanate Ceramics 607

seem to occur—the first with a Ti/Ba ratio of nearly one, and asecond with a Ti/Ba ratio of nearly two. Contrary to the platelikegrains, at least 5 mol% manganese seems to be soluble in the small

globular grains of the manganese-rich samples with a bimodalmicrostructure. But, indeed, a precise determination of the man-ganese content of the small grains by EPMA in compact specimensis hardly possible because of the rather large excitation volume ofthe X-rays.

(B) Annealed Samples: Whereas the samples did not changetheir manganese content during annealing in air, some manganese

Fig. 3. Average grain size as a function of the nominal manganese content (x) of BaTi1-xMnxO3 before and after annealing under oxidizing and reducingconditions. Percentage values at the data symbols denote roughly estimated area portions of the different grain fractions. Grain sizes presented are roughlyestimated values, especially in the case of platelike grains that are obtained by random cuts in the two-dimensional polished plane.

Fig. 4. Bright-field TEM image (U 5 200 kV) of a thinned sample witha nominal composition of BaTi0.95Mn0.05O3 in the region of small globulargrains. In the upper left-hand corner, a hexagonal electron diffractionpattern of the grain in the center is shown. Electron beam coincides withthe [100] direction, and the vertical and horizontal direction of the patternimage belong to the (001) and the (010) plane, respectively.

Fig. 5. EPMA-measured manganese concentration both integrally in asquare and inside the grown grains as a function of the nominal manganesecontent of manganese-doped BaTiO3.

608 Journal of the American Ceramic Society—Langhammer et al. Vol. 83, No. 3

is expelled from the specimens during annealing in a highlyreducing atmosphere.30 The remaining manganese concentrationsare measured by EPMA. The corrected values shown in Fig. 1 arerelated to mean values averaged over an area that covers variousgrains in the central region of the specimens. The amount ofmanganese incorporated inside the grains is even somewhat lowerthan the integrally measured quantities because of manganesesegregation in triple points.

IV. Discussion

(1) As-Fired SamplesTaking into account all previous work on the factors influencing

the cubic/hexagonal phase transformation of BaTiO3 (temperature,oxygen partial pressure, concentration of some acceptor dopants)(see, e.g., Refs. 8, 12, 14), it seems to be reasonable to assume thatone common mechanism is responsible for the stability ofh-BT.Two conditions must be accounted for by that mechanism. First, asufficiently high density of oxygen vacancies is necessary toenable the microscopic steps of the transformation from the cubicto the hexagonal stacking of the Ba–O layers. The transition fromthe cubic ABCABC to the hexagonal ABACBC stacking can besummarized as a reciprocal gliding of the adjacent lattice planes Aand C, which is shown schematically in Fig. 6. Second, a minimumconcentration of BTi

31 ions (B5 titanium, manganese, iron, nickel,. . .) is necessary, which was shown in the case of B5 Ti ([Ti Ti

31]$ 0.3 mol%) by Wakamatsu and co-workers.11,12With respect toall three factors mentioned above, these two conditions can befulfilled. By means of the known defect models of BaTiO3 (see,e.g., Refs. 31, 32) it can be easily derived that, at a fixed oxygenpartial pressure, an increasing temperature causes both an increas-ing concentration of oxygen vacancies and an increasing concen-tration of TiTi

31 ions. At a sufficiently high temperature, theconcentration [TiTi

31] even exceeds the concentration of the metalvacancies. On the other hand, a decrease of the oxygen partialpressure decreases the temperature for the necessary minimumconcentration of both the oxygen vacancies and the trivalenttitanium site ions (see Eq. (1)), which is in accordance with theexperiments. According to the equation

2BaO1 B2O3º 2BaBa3 1 2B9Ti 1 VO

zz 1 5OO3 (2)

the incorporation of trivalent acceptor ions B31 at titanium sitesalso decreases the minimum temperature necessary to formh-BTin an analogous manner, as in the case of decreased ambientoxygen activity. A similar relation also holds for divalent accep-tors. Besides the necessary conditions, the driving force for thephase transition must be elucidated. We propose—at least in thecase of manganese-doped BaTiO3—the influence of the Jahn–Teller distortion (see, e.g., Ref. 33) as such a driving force, whichis explained as follows. Considering the BO6 complex (B5 3dtransition elements) in terms of the crystal field approach, theelectron configuration exhibits the high-spin ground state becauseof weak crystal field splitting of the oxygen ions compared to the

spin-coupling energy.§ In this case, only the electron configura-tions d1, d2, d4, d6, d7, andd9 cause a Jahn–Teller effect (JTE)because of their partial filled sublevels. Moreover, thed4 and thed9 configurations cause a stronger JTE compared to the others,because their partial filling occurs in the higher sublevel. For thatreason, the different effectivenesses of the dopants manganese,iron, and nickel in stabilizing the hexagonal polymorph of Ba-TiO3

8 could be explained easily. Whereas the Mn31(d4) ions causea complete transformation intoh-BT at 1400°C for a rather smallconcentration of dopants (1.7 mol%), it is necessary to add 10mol% FeO or NiO to reach 50%h-BT at the same temperature.8 Ofcourse, the concentration of Fe21(d6) and Ni31(d7), exhibiting aweak JTE, is rather small at high temperatures in air as the ambientatmosphere.34 But concentrations on the order of 0.3 mol% shouldbe sufficient, such as in the case of TiTi

31. To prove this assumptionwith another electron configuration showing a strong JTE, we alsoinvestigated the effect of copper because of thed9 configuration ofCu21. Samples with a nominal composition of Ba1.02Ti0.98Cu0.02O3

and sintered at 1400°C for 1 h exhibit;31% hexagonal phase anda similar microstructure with large platelike grains similar to thecase of manganese. Probably because of the rather high effectiveionic radii of Cu21 and Cu1 at 73 and 77 pm, respectively,35 adefinite excess of barium is necessary to incorporate a significantamount but, of course, far from all of the copper ions at titaniumsites. A considerable amount of copper is a component of thesecondary phase, which is necessary for liquid-phase-assistedgrain growth. Furthermore, because of the well-known valencechange from Cu21 to Cu1 at $1000°C,36 the real concentration ofCuTi

21 is even smaller. Hence, these results corroborate the highefficiency of the incorporation of Cu21 compared to the effect ofiron and nickel. Because thed1 electron configuration of Ti31 alsocauses a JTE, this effect could be the driving force for the cubic3hexagonal phase transition in the case of undoped material as well.

Looking at the microstructural development, three differentgrain growth mechanisms occur that alternate continuously forincreasing manganese content. At low manganese concentrationsof 0 # x # 0.015, the well-known exaggerated grain growth ofundoped BaTiO3 occurs, which is progressively hindered by theincreasing manganese content. Forx $ 0.01, a second growthmechanism appears that results in large platelike grains ofh-BT.Obviously, their growth mechanism is of the same type that isdescribed by Kolaret al. for reducedh-BT as a liquid-phase-assisted one.14–16The pronounced anisotropic grain growth resultsfrom the surface energy anisotropy of the hexagonal polymorph.15

At manganese concentrations$0.025, this exaggerated graingrowth is progressively hindered again, probably because of thesolubility limit of the manganese ions. A third growth mechanismgradually dominates, leading to hexagonal globular grains with alimited size.

(2) Annealed SamplesThe partial restoration of the tetragonal phase of BaTiO3 (t-BT)

during annealing in air can be understood easily if we assume thatthe results of the as-fired samples are not equilibrium values butfrozen states. In Table I, the dependence of the percentage oft-BTon sintering temperature is shown. A 1.6 mol% manganesecontaining sample sintered at 1200°C for 1 h in air exhibits 86%t-BT compared to only 3 mol% at a sintering temperature of1400°C. Thus, during annealing at 1200°C (see Fig. 1), the sampletends to reach its equilibrium state by transformation of someh-BTinto c/t-BT.

The specimens annealed in a highly reducing atmosphere alsoexhibit a restoration oft-BT. Consequently, the same explanationshould be valid. But, two questions arise that cannot be answeredeasily. First, following the discussion above, the ambient reducingatmosphere should originally favor a stabilization of the hexagonalphase and not vice versa. Second, annealing in a reducing

§In the case of manganese and chromium, the high spin ground state was provedby ab initio calculations.25

Fig. 6. Illustration of the reciprocal shift of the adjacent Ba–O latticeplanes A and C, causing the transformation from the cubic (ABCABC) tothe hexagonal (ABACBC) stacking. Only the layers BCA and BAC,respectively, are depicted.

March 2000 Crystal Structure and Related Properties of Manganese-Doped Barium Titanate Ceramics 609

atmosphere is much more efficient than annealing in air. The 1.6mol% manganese-containing sample even undergoes a nearlycomplete transformation fromh-BT into c/t-BT. A possible expla-nation of this phenomenon could be as follows. The crystalstructure ofh-BT is not as symmetric as the perovskite structure ofc-BT. Two types of O22 ions exist that occupy different latticesites with a slightly different bonding distance to the neighboringions. The O22 ions that occupy the face-sharing planes of the TiO6

octahedra (Ti2O9 groups) ofh-BT are more weakly bound becauseof their greater bonding length. Hence, the activation energy forthe creation of oxygen vacancies at those lattice sites is slightlylowered. Therefore, the oxygen vacancies that are produced duringannealing under reducing conditions are created preferentially atthe lattice sites of the face-sharing planes. This process destroysthe Ti2O9 groups that make up two-thirds of the TiO6 octahedra ofthe hexagonal structure of BaTiO3. Consequently, the cubic/tetragonal structure reappears. This mechanism is in competitionwith the influence of the MnTi

31 ions that should stabilize thehexagonal phase because of their Jahn–Teller distortion. Hence,the retransformed percentages of tetragonal phase decrease withincreasing manganese content. The total restoration of the tetrago-nal phase at manganese concentrations lower than;2.0 mol% canbe explained by a third factor controlling the stability of crystal-lographic phases. The effective ionic radius (ri) of MnTi

21, which islikely to be present in highly reducing conditions,34 amounts to 83pm (high spin value).35 Thus, the Goldschmidt tolerance factor,which is somewhat.1 for TiTi

41 (ri 5 60.5 pm) and MnTi31 (ri 5

64.5 pm), decreases to 1, thus stabilizing the perovskite phase. Ofcourse, these explanations are hypothetical in nature and have to beproved by further investigations.

The microstructural development of samples annealed in reduc-ing atmosphere exactly reflects the change of their phase compo-sition related to the situation in the case of the as-sintered samples.But, the drastically changed microstructure of samples with anominal manganese content of 1.0 and 1.6 mol% is difficult toexplain. Because we can assume that at 1200°C no liquid phaseoccurs, a totally diffusion-controlled material transport mechanismhas to be taken into account to describe that size reduction andtotal rearrangement of the grains. In our opinion, the high densityof oxygen vacancies during annealing has to be responsible forthese extensive solid-state grain-size-reduction and grain growthprocesses.

V. Summary

Manganese-doped BaTiO3 ceramics sintered at 1400°C in airchange their room-temperature crystallographic structure fromtetragonal to hexagonal between 0.5 and 1.7 mol% manganese.This hexagonal phase is a nonequilibrium one and can be partiallyretransformed into the cubic/tetragonal one by annealing both inair and in highly reducing atmosphere. As a driving force of thetransformation from the cubic to the hexagonal crystal structure,the influence of the Jahn–Teller distortion is proposed. Thus, it ispossible to explain the known experimental data of the phasetransformation cubic7 hexagonal, induced or stabilized by 3dtransition elements qualitatively in a satisfying manner. The

explanation of the restoration of the cubic/tetragonal phase byannealing in a reducing atmosphere is more speculative and has tobe checked by further investigations.

Three different grain growth mechanisms control the micro-structure in the whole doping range between 0 and 5 mol%manganese. One of them produces exaggerated platelike grainswith hexagonal structure that corresponds to the analogous mech-anism in undopedh-BT sintered in reducing atmosphere describedby Kolar et al.15

The surprising grain size reduction and the rearrangement of thegrains during annealing in reducing atmosphere at 1200°C is notsatisfactorily understood.

Acknowledgment:

The authors thank Dr. Christian Eisenschmidt from the Department of Physics ofthe Martin-Luther-Universita¨t Halle for performing the quantitative XRD investiga-tions and Dr. Stephan Senz from the Max-Planck-Institut fu¨r MikrostrukturphysikHalle for the careful TEM and electron diffraction investigations.

References

1I. Burn and G. H. Maher, “High Resistivity BaTiO3 Ceramics Sintered in CO–CO2Atmospheres,”J. Mater. Sci., 10, 633 (1975).

2S. B. Desu and E. C. Subbarao, “Mn-Doped BaTiO3”; pp. 189–206 in Advancesin Ceramics, Vol. 1,Grain Boundary Phenomena in Electronic Ceramics. Edited byL. M. Levinson. American Ceramic Society, Columbus, OH, 1981.

3J. Rodel and G. Tomandl, “Degradation of Mn-Doped BaTiO3 Ceramic under aHigh d.c. Electric Field,”J. Mater. Sci., 19, 3515–23 (1984).

4F. Batllo, E. Duverger, J.-C. Jules, J.-C. Niepce, B. Jannot, and M. Maglione,“Dielectric and EPR Studies of Manganese-Doped Barium Titanate,”Ferroelectrics,109, 113–18 (1990).

5J. Illingsworth, H. M. Al-Allak, A. W. Brinkmann, and J. Woods, “The Influenceof Manganese on the Grain-Boundary Potential Barrier Characteristics of Donor-Doped Barium Titanate Ceramics,”J. Appl. Phys., 67, 2088–92 (1990).

6D. Y. Wang and K. Umeya, “Spontaneous Polarization Screening Effect andTrap-State Density at Grain Boundaries of Semiconducting Barium Titanate Ceram-ics,” J. Am. Ceram. Soc., 74, 280–86 (1991).

7Y.-C. Chen, G.-M. Lo, C.-R. Shih, L. Wu, M.-H. Chen, and K.-C. Huang,“Influence of Manganese on Lanthanum-Doped BaTiO3,” Jpn. J. Appl. Phys., 33,1412–16 (1994).

8R. M. Glaister and H. F. Kay, “An Investigation of the Cubic–HexagonalTransition in Barium Titanate,”Proc. Phys. Soc., London, 76, 763 (1960).

9R. D. Burbank and H. T. Evans Jr., “The Crystal Structure of Hexagonal BariumTitanate,”Acta Crystallogr., 1, 330–37 (1948).

10D. E. Rase and R. Roy, “Phase Equilibria in the System BaO–TiO2,” J. Am.Ceram. Soc., 38, 102–13 (1955).

11M. N. Wakamatsu, N. Takeuchi, G. C. Lai, and S. Ishida, “Effect of the FiringAtmosphere on the Cubic–Hexagonal Transition and the Chemical State of Titaniumin Barium Titanate,”Yogyo Kyokaishi(Engl. abstract),95, 1181–85 (1987).

12N. Takeuchi, S. Ishida, and M. Wakamatsu, “Mechanistic Study of Reactionsbetween Ceramics and Gases during Firing under Various Atmospheres”; pp. 51–66in Memoirs of the Faculty of Engineering and Design, Vol. 43. Kyoto Institute ofTechnology, 1995.

13O. Eibl, P. Pongratz, P. Skalicky, and H. Schmelz, “Extended Defects inHexagonal BaTiO3,” Philos. Mag. A, 60, 601–12 (1989).

14A. Recnik and D. Kolar, “Exaggerated Growth of Hexagonal Barium Titanateunder Reducing Sintering Conditions,”J. Am. Ceram. Soc., 79, 1015–18 (1996).

15D. Kolar, U. Kunaver, and A. Rec˘nik, “Exaggerated Anisotropic Grain Growth inHexagonal Barium Titanate Ceramics,”Phys. Status Solidi A, 166, 219–30 (1998).

16D. Kolar, A. Recnik, and M. Ceh, “The Origin and Growth Kinetics of Plate-LikeAbnormal Grains in Liquid Phase Sintered Barium Titanate”; pp. 331–38 inCeramicMicrostructure: Control at the Atomic Level, Plenum, New York, 1998.

17J. G. Dickson, L. Katz, and R. Ward, “Compounds with the Hexagonal BariumTitanate Structure,”J. Am. Chem. Soc., 83, 3026–29 (1961).

18L. Katz and R. Ward, “Structure Relations in Mixed Metal Oxides,”Inorg.Chem., 3, 205–11 (1964).

19F. Ren, S. Ishida, and S. Mineta, “Effect of Manganese Addition on PhaseStability of Hexagonal BaTiO3,” J. Ceram. Soc. Jpn. (Yogyo Kyokaishi), 102,106–108 (1994).

20H. T. Langhammer, T. Mu¨ller, A. Polity, K.-H. Felgner, and H.-P. Abicht, “Onthe Crystal and Defect Structure of Manganese-Doped Barium Titanate Ceramics,”Mater. Lett., 26, 205–10 (1996).

21J. Weiss and G. Rosenstein, “Addition of Ba2TiSi2O8 to Manganese-DopedBarium Titanate: Effect on Oxygen Diffusion and Grain-Boundary Composition,”J.Mater. Sci., 23, 3263–71 (1988).

22I. Burn, “Mn-Doped Polycrystalline BaTiO3,” J. Mater. Sci., 14, 2453–58 (1979).23H.-J. Hagemann and H. Ihrig, “Valence Change and Phase Stability of 3d-Doped

BaTiO3 Annealed in Oxygen and Hydrogen,”Phys. Rev. B: Condens. Matter, 20,3871–78 (1979).

24S. B. Desu and E. C. Subbarao, “Effect of Oxidation States of Mn on the PhaseStability of Mn-Doped BaTiO3,” Ferroelectrics, 37, 665–68 (1981).

25P. Moretti and F. M. Michel-Calendini, “Impurity Energy Levels and Stability ofCr and Mn Ions in Cubic BaTiO3,” Phys. Rev. B: Condensed Matter, 36, 3522–27(1987).

Table I. Percentage of Tetragonal Phaseof BaTi12xMnxO3 Ceramics as a Functionof the Sintering Temperature, Ts (air), for

Several Manganese Concentrations,x

Ts (°C)

x

0.016 0.020 0.050

1000 92 88 741100 90 88 691200 86 85 431300 84 74 121400 3 0 0

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26B. Milsch, “Evaluation of Lattice Site and Valence of Manganese in Polycrys-talline BaTiO3 and n-BaTiO3 by Electron Paramagnetic Resonance,”Phys. StatusSolidi A, 133, 455–64 (1992).

27J. Y. Kim, C. R. Song, and H. I. Yoo, “Mn-Doped BaTiO3: Electrical TransportProperties in Equilibrium State,”J. Electroceram., 1, 27–39 (1997).

28H. T. Langhammer, T. Mu¨ller, K.-H. Felgner, and H.-P. Abicht, “Influence ofStrontium on Manganese-doped Barium Titanate Ceramics,”Mater. Lett., 42, 21–24(2000).

29H.-P. Abicht, D. Voltzke, A. Roder, R. Schneider, and J. Woltersdorf, “TheInfluence of the Milling Liquid on the Properties of Barium Titanate Powders andCeramics,”J. Mater. Chem., 7, 487–92 (1997).

30V. Bheemineni, E. K. Chang, M. Lal, M. P. Harmer, and D. M. Smyth,“Suppression of Acceptor Solubilities in BaTiO3 Densified in Highly ReducingAtmospheres,”J. Am. Ceram. Soc., 77, 3173–76 (1994).

31J. Daniels and K. H. Ha¨rdtl, “Electrical Conductivity at High Temperatures ofDonor-Doped Barium Titanate Ceramics,”Philips Res. Rep., 31, 489–504 (1976).

32N.-H. Chan, R. K. Sharma, and D. M. Smyth, “Nonstoichiometry in UndopedBaTiO3,” J. Am. Ceram. Soc., 64, 556–62 (1981).

33S. F. A. Kettle, Physical Inorganic Chemistry; p. 166. Spectrum AcademicPublishers, Oxford, U.K., 1996.

34H.-J. Hagemann, “Acceptor Ions in BaTiO3 and SrTiO3 and Their Consequenceson the Properties of Titanate Ceramics”; p. 44 in Ph.D. Thesis. Rheinisch-Westfalische Technische Hochschule Aachen, Germany, 1980.

35R. D. Shannon, “Revised Effective Ionic Radii and Systematic Studies ofInteratomic Distances in Halides and Chalcogenides,”Acta Crystallogr., Sect. A:Cryst. Phys., Diffr. Theor. Gen. Crystallogr., 32, 751–67 (1976).

36A. F. Hollemann and E. Wiberg,Lehrbuch der Anorganischen Chemie; p. 792.W. de Gruyter, New York, 1976. M

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