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Available online at www.sciencedirect.com

ScienceDirect

Journal of the European Ceramic Society 34 (2014) 943–951

Anisotropy and thermal stability of hot-forged BICUTIVOX oxygen ionconducting ceramics

Paul Fuierer a,∗, Mitchell Maier a, Jörg Exner b, Ralf Moos b

a Department of Materials & Metallurgical Engineering, New Mexico Institute of Mining & Technology, 801 Leroy Place, Socorro, NM 87801, USAb Department of Functional Materials, Faculty of Engineering Science, University of Bayreuth, 95447 Bayreuth, Germany

Received 16 July 2013; received in revised form 14 October 2013; accepted 15 October 2013Available online 14 November 2013

bstract

he oxygen ion conductivity of ultra dense, hot forged Bi4(MexV1−x)2O11 (BIMEVOX) ceramics is reported for the first time. A novel approach tolectrode preparation, including laser trenching, was developed to perform 4-probe impedance spectroscopy and determine conductivity of smallar-shaped samples with high accuracy and precision. Hot forged (HF) Bi4(Cu0.05Ti0.05V0.90)2O11 (BICUTIVOX) ceramic, with modest preferredrientation (11%), exhibits anisotropy by a factor of 2 in its conductivity. Brick layer model (BLM) analysis of complex impedance plots reveals

′

hat the relative contribution of grains and grain boundaries to the total resistance is direction dependent, in the low temperature � phase. The totalonic conductivity of BICUTIVOX, prepared either by hot-forging or conventional sintering, appears to be stable with thermal aging and cycling.nisotropy in the thermal expansion of polycrystalline BIMEVOX ceramics is also reported.

2013 Elsevier Ltd. All rights reserved.

eywords: BIMEVOX; Hot-forging; Preferred orientation; Ionic conductivity; Thermal expansion

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apiwp

. Introduction

Bismuth vanadate or Bi4V2O11 (abbreviated as BIVOX) andelated compounds with aliovalent substitution for vanadiumBIMEVOX), including Bi4(CuxV1−x)2O11 (BICUVOX), haveigh oxygen conductivity that makes them potential candidatesor solid oxide fuel cells (SOFCs), gas sensors and gas separationembranes.1–4 The oxygen conductivity of such compounds

s two orders of magnitude greater than that of state of thert stabilized zirconia, over a wide range of temperatures.5

ome BIMEVOXes, however, can exhibit an increase in electronransference number and phase instability with reduced oxygenartial pressure at high temperature.6,7 Hysteresis in the low tem-

erature conductivity can occur, particularly with single crystalnd large grain ceramics.8–10 Strength and toughness of theeramic has also been questioned.11,12 These reports have raised

∗ Corresponding author at: New Mexico Institute of Mining and Technology,01 Leroy Pl., Socorro, NM 87801, USA. Tel.: +1 575 835 5497.

E-mail addresses: fuierer@nmt.edu (P. Fuierer), mmaier@umail.ucsb.eduM. Maier), Joerg.Exner@uni-bayreuth.de (J. Exner),alf.Moos@uni-bayreuth.de (R. Moos).

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955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved.ttp://dx.doi.org/10.1016/j.jeurceramsoc.2013.10.016

oubts about their utility in SOFCs. Research continues exper-menting with various ion substitutions, mainly to increase thehermodynamic stability.13,14 Over the years, increasing atten-ion has been given to the effects of ceramic processing and

icrostructure on the behavior of various BIMEVOXes,11,16–18

ut a satisfactory solution for their use as solid electrolytes hasot yet been found.

The crystal structure of BIVOX can be generally describeds an alternating stack of (Bi2O2)2+ layers and [(VO3.5–0.5)]2−erovskite sheets.19–21 The room temperature structure, �, hav-ng orthorhombic symmetry undergoes a series of transitionsith increasing temperature (α − β − γ) with the high tem-erature � phase having tetragonal symmetry and high ioniconductivity above about 570 ◦C. This �-phase can be stabilizedt lower temperature by a number of aliovalent substitutions.he high conductivity is attributed to a high concentration ofisordered oxygen ion vacancies (V=

0 ) and a circular diffusionathway, particularly on and between the O2 and O3 sites withinhe perovskite-like slabs.21,22

Due to the layered structure, the conductivity of BIMEVOXompounds is expected to be anisotropic when comparing theirection parallel to the (Bi2O2)2+ layers (perpendicular to-axis) and perpendicular to the layers (parallel to c-axis). Kim

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44 P. Fuierer et al. / Journal of the Euro

nd Miyayama23 confirmed the strong anisotropy using sin-le crystals of both BIVOX and BICOVOX (two and threerders of magnitude respectively). This anisotropy in conduc-ivity might be exploited through grain orientation in ceramics.

uller et al.24 were the first to obtain texture in BICOVOX viaressing and application of strong magnetic field. An orientationf 50% was achieved, so anisotropy was about half that of sin-le crystals. Shantha and Varma25 produced partially oriented∼75%) BIVOX by liquid phase-aided sintering, and reportedhe dielectric and ferroelectric properties.

The solution to the BIMEVOX problems mentioned aboveay be a combination of (i) composition tuning using dou-

le substitution for stability, (ii) achieving ultra-high densityor maximizing strength and minimizing degradation withreducing) gas exposure, and (iii) grain orientation in ordero maximize bulk ionic conductivity and lower the operationemperature even further. Some of the highest conductivityalues have been reported for double substitution compoundi2Cu0.05Ti0.05V0.9O11 (BICUTIVOX).26–28 Hot forging (HF)

s a means by which preferred grain orientation has beenchieved in layered-structure ferroelectric ceramics,29 andt may be a way to optimize the electronic properties ofIMEVOX. HF can produce ceramics with maximal densitynd therefore near zero open porosity, improved thermochemi-al stability, and increased strength/toughness. HF also providesodel polycrystalline specimens to which other sample types,

uch as aerosol or spray deposited thick films,30,31 can beompared.

Previously, we described the improved processability of BIC-TIVOX, and the fabrication of hot forged ceramic.32 In the

urrent manuscript, we present the direction dependent ioniconductivity and thermal stability of samples machined fromhese forged specimens, and compare to conventionally sinterederamic. Direction dependent thermal expansion is also char-cterized, since matching to electrode supports is critical formplementation of new electrolytes.

. Experimental

.1. Synthesis procedures

Conventional Bi4(Cu0.05Ti0.05V0.90)2O11 (BICUTIVOX)eramic samples, referred to in this paper as “ordinaryred” (OF), were fabricated using established solid-stateowder preparation and pressureless sintering, as describedreviously.32 Bar-shaped compacts with approximate dimen-ions 17 mm × 7 mm × 3 mm were pressed and sintered in air atemperature of 775 ◦C for a period of 10 h, followed by slow fur-ace cooling (∼2 ◦C/min). For the hot-forged (HF) specimen, aylindrical sample (25 mm diameter, 15 mm height) was pressednd bisque-fired at 550 ◦C prior to forging. Hot-forging was

erformed by loading (300 kg) the specimen between two pol-shed stabilized zirconia ram surfaces inside a vertical clamshellurnace at a soak temperature of 750 ◦C.

A1pt

Ceramic Society 34 (2014) 943–951

.2. Microstructural characterization

For microstructural imaging, samples were sectioned with aiamond saw, surfaces were ground with decreasing grit SiCaper, followed by a 1 �m colloidal alumina polish. They werehen thermally etched at 700 ◦C for 10 min. The surfaces werexamined without metallic coating using a variable pressurecanning electron microscope (Hitachi, S-3200).

.3. Physical property measurement

Two bar shaped samples were cut from the forged specimen,ne with its long dimension perpendicular to the forging axisreferred to as HFperp), and one with its long dimension par-llel to the forging axis (referred to as HFpara). These namesre used to make a distinction and avoid confusion with labelssed elsewhere in the literature. For example �⊥ and �||, haveeen used to represent the conductivities measured perpendic-lar and parallel to the “intergrowth structure” in single crystalICUVOX.8 In our case, however, the HF sample is known toave only modest texture, and therefore we choose to take theorging axis (FA) as the directional reference.

All bars were hand machined and surface ground (1000 gritnish) to create rectangular prisms with high degree of paral-

elism. Dimensions were measured with error less than 10 �msing a precision micrometer. Sample volumes for densitiesere determined by both dimensional measurements and by pyc-ometry (AccuPyc 1330, Micromeretics, GmbH). A dilatometerL75, Linseis) was used to measure thermal expansion of the barpecimens. Dimensional change was measured both on heatingnd cooling, at rate of 5 ◦C/min.

.4. Electrical characterization

The bars were then prepared for 4-probe impedance mea-urements using a unique electroding procedure. Two shallowemi-circular grooves (∼0.050 mm depth) were precisionachined around the perimeter of the sample using a YAG:Nd

aser in order to accommodate tight and straight wrapping ofhe 0.100 mm diameter platinum wire used for the inner con-acts. The locations of these grooves were at 1/3 and 2/3 alonghe length of the specimen. Platinum paste (Heraus CL11-5100)ith a tailored viscosity (7 kcps) was applied to the wire, filling

he trench and creating continuous contact between the ceramicample and the platinum wire. The paste was applied to thends of the bar samples in a uniformly thick film and thenintered at 750 ◦C for 5 h to a porous but continuous electri-ally conducting film with thickness of about 50 microns. Thisethod produced straight and uniform electrodes with well-

efined interelectrode distance to meet the objective of obtainingccurate impedance/resistance measurements (see Fig. 1).

A computer controlled test station integrating a small tubeurnace and high resolution dielectric analyzer (Novocontrol

lpha-Analyser) was used to measure complex impedance from07 Hz to 0.1 Hz. A sample holder with shielded current andick-up leads and a spring-tensioned pusher rod to make con-act to the end electrodes was employed. The most significant

P. Fuierer et al. / Journal of the European Ceramic Society 34 (2014) 943–951 945

Fig. 1. (a) Edge view of HFperp after machining showing the 100 �m diameter laser trenches. (b) Surface view of HFperp after wrapping with platinum wire andfi rainedh Fparaf

siks

Y

Y

wcpfbe

atftfRiAclfiw

3

3

fTsaoTsm

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3

ring platinum paste. The elongated blemishes are oriented domains of fine gighlighted example of an oriented domain of fine-grains. (d) Surface view of Horging axis (FA) is indicated.

ource of error in a 4-wire connection is the parallel stray capac-tance, Csp, of the sample holder,33 which was minimized byeeping cables as short as possible, and was subtracted from thepectra as a parallel imaginary admittance according to:

∗s = Y∗

tot − iY ′′sh (1)

′′sh = 2πfCsp (2)

here Y∗s , sample complex admittance, Y∗

tot, total measuredomplex admittance, and Y ′′

sh, imaginary admittance of the sam-le holder. Measurements were made over the temperature rangerom 100 ◦C to 700 ◦C in 50 ◦C intervals, with a heating rateetween steps of 2.5 ◦C/min, and equilibration time of 20 min atach temperature.

With lower temperature data, Nyquist plots were constructednd Winfit software (Nova Control) used to fit impedance spec-ra. With higher temperature data (>400 ◦C), the relaxationrequencies were beyond the measurement range, so intersec-ion of the complex impedance plot with the real axis at lowrequency was taken to be the total bulk sample resisitance.esistance values were then converted to material conductiv-

ty using sample crossectional area and inner winding distance.rrhenius plots were used to calculate activation energies for

onduction. Thermal stability was then examined using extended◦ ◦

ow temperature (425 C and or 450 C) anneals, immediately

ollowed by another impedance measurement cycle, both heat-ng and cooling. For these measurements, temperature stepsere 100 ◦C.

pr

material. (c) Surface view of HFperp, showing in-laid platinum wire, and a, showing fine-grain domains with orthogonal orientation. The direction of the

. Results and discussion

.1. Microstructure

Fig. 2(a) and (b) shows typical views of the microstructureor the HF sample. Grain size (d) ranges from 1 to 12 �m.he surface perpendicular to the forging axis (FA) tends tohow slightly larger particle size than the cross-sectional par-llel surface. Fig. 2(c) and (d) shows the corresponding surfacesf the OF sample, with respect to the dry pressing direction.he HF sample appears to have a wider size distribution withome smaller grains, but over a majority of the samples, theicrostructures are similar with apparent normal grain growth.There are, however, regions in the HF sample that are atypi-

al, revealing some preferred orientation with respect to the FAFig. 3). This is in agreement with 11% grain orientation, mea-ured by X-ray diffraction.32 One also finds distributed ellipticalomains of very fine grained material (d < 1 micron), with shortxis of the domain aligned with the FA. These can be seen inig. 1(b–d). A higher magnification image of such a domain washown previously.32 In three dimensions, these domains are thehape of an oblate ellipsoid. We surmised that these fine-grainedegions accommodate the contraction experienced during theorging process as a consequence of superplastic deformationia grain boundary sliding.

.2. Physical properties

Pycnometer measurement of solid bodies can yield highlyrecise volumes, however open porosity in a sample will giveise to a large error in measured bulk density. As in careful

946 P. Fuierer et al. / Journal of the European Ceramic Society 34 (2014) 943–951

Fig. 2. Dense BICUTIVOX ceramics. Hot forged (HF) sample; view of (a) surface perpendicular to the forging axis (FA); (b) surface parallel to FA. Ordinary fired(OF) sample; (c) plan view of large surface of pressed disk, and (d) cross-sectional view.

Table 1Bulk densities (BD) of BICUTIVOX, hot-forged (HF) and ordinary fired (OF). Relative density is based on theoretical (X-ray) density of 7.82 g/cm3.26

BD, g/cm3(geometric) BD, g/cm3(pycnometer) Relative density Uncertainty

HF perp 7.575 7.782 99.6% 1.1%HF para 7.631 7.645

OF 7.266 7.488

Fig. 3. An atypical cross-sectional view (surface parallel to FA) of HF BICU-Tg

wt

dhdt

spboer

α

afinding by Paydar et al. a hysteresis of about 10 C was

IVOX containing both partial orientation of large grains and a region of finerains.

ork with BICOVOX by Steil et al.9 geometric densities ofhe machined samples were determined to verify pyconmeter

oro

97.9% 1.6%95.8% 1.7%

ensity values (see Table 1). We conclude that the OF sampleas a density of about 96%, while the HF samples are 98–99%ense. This suggests an absence of open porosity, especially inhe HF samples.

Relative thermal expansion (dL/L0) of the three samples ishown in Fig. 4(a). HFperp and HFpara have strain measurederpendicular and parallel to the FA respectively. Anisotropyetween the two directions is measurable, even with the low levelf grain orientation. Fig. 4(b) shows the physical linear thermalxpansion coefficient (α), corresponding to the derivative of theelative expansion curve, and defined as:

(T ) = 1

L0

(∂(dL)

∂T

)(3)

This tends to accentuate anomolies such as the peak occuringt the �′ to � phase transition. Similar to previous DTA/TGA

27 ◦

bserved with cooling (not shown for clarity sake). Therefore,eported transition temperatures (Tt) are taken to be the midpointf heating and cooling (Table 2). With linearity over a specified

P. Fuierer et al. / Journal of the European Ceramic Society 34 (2014) 943–951 947

Table 2Thermal coefficients of expansion, TCE, (ppm/◦C) of hot-forged (HF) and ordinary fired (OF) BICUTIVOX above and below the �′ − � phase transition temperature(Tt). Also included is a hypothetical randomly oriented ceramic using a simple weighted average of the HFperp and HFpara values according to Eq. (5).

BICUTIVOX sample Tt�′ to �-phase 150–400 ◦C�′ phase 500–700 ◦C�-phase

HF perp 443 ◦C 16.3 19.9HF para 445 ◦C 15.5 17.1Random (calculated) 16.0 19.0OF 436 ◦C 16.2 19.8

10

15

20

25

30

700600500400300200

αα(ppm

/°C)

Temperature (°C)

HF perp

OF

HF para

(b)

0.2

0.4

0.6

0.8

1

700600500400300200

Rel

ativ

e Ex

pans

ion

(%)

Temperature (°C)

HF perp

OF

HF para

(a)

Fig. 4. (a) Relative thermal expansion and (b) physical linear thermal coeffi-cs

rc

T

Hief

T

0

0.5

1

1.5

0 0.5 1 1.5 2 2.5 3

norm

aliz

ed -

Zs"

(kΩ

·cm

)

normalized Zs' (kΩΩ·cm)

(b) 250°C

6

66 5

4

75

457

HF para

OF

HF perp

0

10

20

30

0 20 40 60 80norm

aliz

ed -

Zs"

(kΩΩ

·cm

)

normalized Zs' (k ΩΩ·cm)

(a) 150°C

5

HF perp

HF para

OF5

2

3

4

6 4 3

5

Fig. 5. Nyquist plots on the complex impedance plane for the three BICUTIVOXsamples at a temperature of (a) 150 ◦C and (b) 250 ◦C. Data are normalized usingta

rt

3

sf

F

sbTe

ient of expansion of BICUTIVOX ceramics. Samples machined from the HFpecimen with orthogonal orientations display measureable anisotropy.

ange of temperatures (T1 − T0), average technical thermaloefficients of expansion (TCEs) are defined as:

CE = (dL/L0)T1− (dL/L0)T0

T1 − T0(4)

The TCEs of HFperp are about 15% and 5% higher thanFpara in the � and �′ phase regimes respectively (Table 2). This

s the first report of TCE anisotropy in BIMEVOX ceramics. Thexperimental TCEs for the OF sample lie between, as expectedrom a weighted averaging:

CErandom = 2TCEHFperp + TCEpara

3(5)

ccl

he geometric factor, FG, of each sample for direct comparison. Small numberslong each curve indicate the log of the measurement frequency.

Values are similar to those reported by Yaremchenko34 forare-earth doped BICUVOX (∼15 ppm/◦C and ∼17 ppm/◦C forhe low and high temperature regimes respectively.

.3. Impedance spectroscopy and conductivity

HFperp and HFpara were prepared from the same hot forgedpecimen for electrical measurements with similar geometricactors (FG = 0.215 and 0.191 cm respectively).

G = wt

l(6)

In Eq. (6), wt is the cross-sectional area and l is the electrodeeparation distance. Measured impedances, Z*, were multipliedy FG to construct the normalized Nyquist plots shown in Fig. 5.he overlapping half-circles in the complex impedance plane arexplained using the brick layer model (BLM) for polycrystalline

eramics with resistive grain boundaries, with the high frequencyircle (left) attributed to grain resistance and capacitance, and theow frequency (right) attributed to grain boundary.35 Total bulk

948 P. Fuierer et al. / Journal of the European

-5.00

-4.00

-3.00

-2.00

-1.00

0.00

0.8 1. 0 1. 2 1. 4 1. 6 1. 8 2. 0 2. 2 2.4

log

(σσ(S

/cm

))

1000/T (1/K)

OF

HF perp

HF para

Fig. 6. Arrhenius-like plot of conductivity vs temperature for BICUTIVOXca

rttmegtsst

aToritui

ticTit(tagvcbOHoc

t

B“stsd

ttHfactHtTtchtttfdnd

3

daWptpitto

aBTcarptOcoa long-term aneal at 425 C (just below the transition) results in

eramics. OF, ordinary fired; HF perp, hot-forged, perpendicular to the forgingxis. HF para, hot-forged, parallel to the forging axis.

esistivity is given by the intersection with the real axis. Conduc-ivity of the HFperp is higher than HFpara by a factor of 2 overhe entire range from 50 ◦C to 700 ◦C (Fig. 6). Error bars (esti-

ated at ±3% from dimensional and temperature measurementrror) are included on the HFpara data. Even though the overallrain orientation is low, careful impedance measurements revealhe anisotropy. In the lower temperature �′ phase regime, the OFample exhibits a slightly higher conductivity than the HFperpample. In the � phase, the difference in conductivity is withinhe limits of measurement error.

Activation energies for conduction are given in Table 3, andre in agreement with previous reports for BICUTIVOX.26,27

he higher value of 0.7 eV in the �′ phase is attributed to anrdering of oxygen vacancies. The transition temperature, Tt,eported in Table 3 for our experiments was determined from thentersection of the two straight line fits (� and �′) to the conduc-ivity. In previous reports, smaller temperature intervals weresed, revealing a temperature range of transition, with a thirdntermediary slope change.27,32 This explains the discrepancy.

A closer examination of the Nyquist plots reveals addi-ional information about BICUTIVOX ceramics. Fitting of thempedance curves using a series connection of parallel RC cir-uits according to the BLM yields the parameters given inable 4. A constant phase element (CPE) was used as the capac-

tance element in the equivalent circuit to produce best fitso the experimental data. Thus the reported relaxation timestime constant, τ) are average values. Capacitance values forhe high frequency (HF) and low frequency (LF) semicirclere on the order of 100 pF and 10 nF, appropriate for bulk andrain boundary contributions respectively, and similar to pre-ious BICUVOX analysis.15 We attempt interpretation withaution, as non-ideal and inhomogeneous microstructures haveeen shown to cause significant deviations from the BLM.36

ne might speculate that the small tail at low frequencies on theFpara plots is an additional polarization related to the domainsf small grains. These data were not included in the equivalent

ircuit model.

Larger bulk resistance for the HFpara can be attributed tohe expected anisotropy according to the layered structure of

ada

Ceramic Society 34 (2014) 943–951

iMEVOx compounds. The HFperp sample corresponds to thehigh oxygen ion conductivity” direction (a–b plane). We canpeculate that the lower RC time constant in HFpara may be dueo the platelet shape of some grains, the smaller grain dimen-ion in this direction, and therefore a shorter relaxation time foriffusing oxygen ions.

The HF specimen clearly has a larger grain boundary resis-ance (RLF), particularly in the HFpara configuration. At 150 ◦C,he grain boundary contribution to total resistance is 36% forFpara, 26% for HFperp, and 14% for OF. This contribution

rom the grain boundary diminishes with increasing temper-ture for all samples, until a temperature of 300 ◦C, whereapacitances converge and the two semicircles overlap such thathey cannot be resolved. The difference between HFperp andFpara is attributed to a greater parallel conductance shunting

he series connection of grains and grain boundaries in HFperp.he higher conductivity of the OF sample in the �′ regime can

hen be attributed to even greater parallel connectivity, as indi-ated by its lower RLF value. These results support our previousypothesis concerning BICUTIVOX.32 Namely, that the highhrough-thickness 2-point conductivity of conventionally sin-ered BICUTIVOX pellets (in previous work) is due to a lack ofhe grain orientation effect compared to other BIMEVOX. Weurther hypothesize that high-angle grain boundaries limit con-uction in the HFperp sample with only 11% orientation. Thiseeds to be tested with measurements on samples with greateregree of orientation.

.4. Thermal stability

Previous studies have reported hysteresis in total con-uctivity between heating and cooling runs with BICUVOXnd BICOVOX8–10,15 particularly with large grained ceramic.atanabe10 suggested that the lower conductivity at lower tem-

eratures on initial heating is due to the ordered �-phase, whichransitions to conducting �-phase on heating, and cools to the �′hase, which is actually metastable. The �′ back to � conversions very sluggish, but can be accelerated by long term annealing atemperatures of about 450 ◦C. Piravono’s work15 suggests thathis transition to the ordered �-phase may be avoided by controlf the microstructure (i.e. by maintaining small grain size).

In the case of our dense BICUTIVOX, such hysteresis andging was not observed. Aging experiments were conducted onICUTIVOX, OF at 425 ◦C, and HFperp at 425 ◦C and 450 ◦C.hese temperatures were chosen based on both previous workited above, as well as the transition temperature (435–445 ◦C)pparent from our dilatometry experiments. Comparing withesults reported by Watanabe10 on BiCUVOX (conventionallyrepared) ceramics, we see virtually no degradation in conduc-ivity (Fig. 7). Figs. 8 and 9 illustrate the thermal stability of theF and HFperp samples with temperature cycling. In both cases,

ooling from the high temperature �-phase always retraces theriginal heating curve. The only measureable difference is that

◦

n increase in the low temperature conductivity. The differenceecreases approximately linearly with increasing temperature,nd is non-existent once the transition temperature is reached

P. Fuierer et al. / Journal of the European Ceramic Society 34 (2014) 943–951 949

Table 3Conductivity and phase transition parameters for BICUTIVOX compared to previous literature.26,27 Activation energies are calculated from Arrhenius ln(σT) plots.

�@500 ◦C (cm)−1 T (�′ − �) (◦C) Ea, �′ phase (eV) Ea, �-phase (eV)

HFperp 8.61 × 10−2 510 0.701 0.436HPpara 4.07 × 10−2 510 0.701 0.438OF 9.35 × 10−2 500 0.694 0.411(literature) (8.6 × 10−2) (460–480) (0.71) (0.43)

Table 4Estimated equivalent circuit parameters (normalized using geometric factor, FG) including average RC time constant (τ) from impedance spectra taken at 150 ◦C.

High frequency semicircle (intragranular) Low frequency semicircle (grain boundary)

Normalized RHF (k cm) Normalized CHF (nF/cm) τHF(�s) Normalized RLF(k cm) normalizedCLF(nF/cm) τLF(�s)

HF perp 26.1 0.54 14.1 8.75 21.6 190HF para 46.9 0.072 3.39 25.6 13.8 353OF 23.8 0.75 17.8 4.09 28.2 115

-2.6

-2.4

-2.2

-2

-1.8

-1.6

-1.4

-1.2

-1

0 20 40 60 80

log

( σσ(S

/cm

))

time (h)

HF perp 450°C

HF perp 425°C

OF 425°C

BICUVOX, 430°CWatanabe ref[10]

Fig. 7. Annealing dense BICUTIVOX at temperatures just below and just abovethe detected � − �′ phase transition (435–440 ◦C) appears to have little effecton conductivity compared to previous tests with BICUVOX.10

-8

-7

-6

-5

-4

-3

-2

-1

0

1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0

log

(σσ(S

/cm

))

1000/T (1/K)

run 1 heatin g

run 2 heating

run 2 coolin g

Fig. 8. Thermal cycling has little effect on the conductivity of OF BICUTIVOX.Run 1 is the first heating cycle. Run 2 took place immediately after cooling(∼1 ◦C/min) from the extended 425 ◦C anneal. Measurements were taken in100 ◦C increments during heating to 650 ◦C, then cooling back to 50 ◦C. Theonly change is a small increase in low temperature conductivity following the425 ◦C anneal.

-8

-7

-6

-5

-4

-3

-2

-1

0

1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0

log

(σσ(S

/cm

))

1000/T (1/K)

run 1 heatin g

run 2 heatin g

run 2 coolin g

run 3 heatin g

run 3pa cooling

Fig. 9. Thermal cycling also has little effect on the conductivity of HF BICU-TIVOX. Run 1 and Run 2 were conducted as described in Fig. 8. Run 3 beganseveral weeks after Run 2, with data points duplicating the initial Run 1. Run3pa (pa, post anneal) took place immediately after the extended 450 ◦C anneal,cf

(awtBots

4

ccmod

ooling first. The only change is a small increase in low temperature conductivityollowing the 425 ◦C anneal.

roughly 34% at 50 ◦C, 26% at 150 ◦C, 18% at 250 ◦C, 8%t 350 ◦C, and 0% at 450 ◦C). This increase in conductivityith low temperature aneal, consistent in both samples, is con-

rary to previous reports of single substitution BICUVOX andICOVOX. We surmise that in the case of Ti substitution, a dis-rdering rather than ordering of oxygen ions may occur aroundhe � − �′ transition temperature. This, however, is a metastableituation, and has only a very minor effect on the conductivity.

. Conclusions

BICUTIVOX ceramics with density approaching theoreti-al (98–99%) can be obtained by hot forging. In the presentase, a heterogeneous microstructure was obtained, with a

ajority of normal grain growth, but with regions of larger,

riented grains, and also very fine grained material located inistributed disk shaped domains. Partially textured specimens

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isplay anisotropic thermal expansion, with expansion greatererpendicular to the forging axis than parallel to it: 16.3 vs5.5 ppm/C in the low temperature �′ regime, and 19.9 vs 17.1n the high temperature � regime. This may have implicationsith respect to metallic anode supports.Impedance measurements revealed anisotropy in the con-

uctivity of hot forged BICUTIVOX, even when the preferredrientation is a modest 11%. Total conductivity is a factor of tworeater when measured perpendicular to the forging axis (paral-el to the perovskite-like layers in the crystalline structure) from0 ◦C to 700 ◦C. Impedance spectroscopy in the low temperatureegime below the � − �′ transition reveals details of the relativeontributions of bulk grains and grain boundaries to the totalesistance of BICUTIVOX ceramics. In HFpara, greater seriesonnectivity of the perovskite-like layers in the crystal struc-ure and of tabular grains in the microstructure yields higheresistance. Well-prepared ordinary fired ceramic displays a con-uctivity at least equal to the forged (HFperp) sample, due in parto greater parallel connectivity. It is suspected that high anglerain boundaries present in HFperp due to imperfect orientation,imit the conductivity.

Overall, we find that the combination of Ti-substitution tonhibit grain growth and high density can serve to stabilize theonic conductivity in BICUTIVOX. While small grained, dense,nd randomly oriented BIMEVOX may yield high conductivityue to disorder in the grain boundaries, it is not yet known whatole the grain boundaries play in properties such as high temper-ture red-ox stability and electron transport. These factors maye a reason to produce well crystallized material with a dense,riented microstructure. Work continues to increase the extentf texturing and also further assess thermal and atmospheretability.

cknowledgements

Funding for this project is provided through the NSF-DFGaterials World Network program, NSF grant # 1108466 andFG grant # MO 1060/16-1. The authors acknowledge help

rom Dr. Jarislaw Kita for laser trenching and Dr. Gunther Hagenor designing and building the impedance sample holder, bothn the Functional Materials Department at the University ofayreuth. Acknowledgment also goes to Mr. Michael Adam atinseis Corporation (Selb, Germany) for performing dilatome-

ry experiments.

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