Structure and properties of plasma sprayed BaTiO3 coatings after thermal posttreatment

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Structure and properties of plasma sprayed BaTiO3 coatings after thermal posttreatment Pavel Ctibor, Josef Sedlacek, Zdenek Pala

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S0272-8842(15)00284-9 http://dx.doi.org/10.1016/j.ceramint.2015.02.065 CERI9987

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Cite this article as: Pavel Ctibor, Josef Sedlacek, Zdenek Pala, Structure and properties of plasma sprayed BaTiO3 coatings after thermal posttreatment, Ceramics International, http://dx.doi.org/10.1016/j.ceramint.2015.02.065 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Structure and properties of plasma sprayed BaTiO3 coatings after thermal posttreatment

Pavel Ctibor1,2, Josef Sedlacek2, Zdenek Pala1

1 - Institute of Plasma Physics, ASCR, Za Slovankou 3, 182 00 Praha 8, Czech Republic 2 - Department of Electrotechnology, Faculty of Electrical Engineering, Czech Technical University, Technická 2, 166 27 Praha 6, Czech Republic

Corresponding author. Tel: +42-266053717; fax: +42-286586389. E-mail address: [email protected] (Ctibor)

Abstract Previously published results on electrical and mechanical properties of BaTiO3 coatings prepared by atmospheric plasma spraying showed anomalies in their dielectric response. This paper provides a study of electrical and mechanical properties of BaTiO3 coatings after thermal posttreatment. The spraying was carried out by a direct current gas-stabilized plasma gun. BaTiO3 was fed into the plasma jet as a feedstock powder prepared by reactive sintering of micrometer-sized powders of BaCO3 and TiO2. In the next step the coatings were annealed in air. Microstructure and phase composition are reported and discussed in relation to electric and mechanical properties. Dielectric properties are reported for the radio frequency (RF) range.

Keywords: BaTiO3, Plasma spraying, Annealing, Electrical properties, Microstructure

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1

Introduction

Barium titanate, BaTiO3, is a multifunctional oxide that exhibits complex phase appearance. Due to its high relative permittivity, BaTiO3 is used in multilayer ceramic capacitors, dynamic random access memories, piezoelectric sensors, thermistors and resistors with high positive thermal coefficients [1,2]. Between 120°C (393K) and 1457°C (1730K) BaTiO3 has a cubic perovskite structure that consists of corner linked oxygen octahedra containing Ti4+, with Ba2+. Cooling below 120°C results in small displacements in the positions of cations in the unit cell resulting in a polar ferroelectric phase existing in the temperature interval between 5°C (278K) and 120°C [3]. Perovskite barium titanates BaTiO3 and (Ba,Sr)TiO3, due to high relative permittivity, are used frequently as multilayer capacitor components and sensors. However, it has been found that with respect to the electrical properties, BaTiO3 in the form of thin films does not reach the qualities of the bulk material. This difference was explained by a combination of the intrinsic dead layer effect, stress effect, effect of microstructure within the thin film, and the effect of stoichiometry [4]. In particular, the relative permittivity of films decreases if the film thickness is reduced [5]. The optimal dielectric characteristics are obtained for the sintered bulk BaTiO3 sample with a density of about 5300 kg.m-3 [6]. Another category of coatings that have many applications in electronics industries are coatings with thicknesses between 10 µm and 1 mm, known as thick films [7,8]. The properties of these films are usually compared with bulk materials. One of the most important methods for producing thick films is thermal spraying [9]. Any deviation from the stoichiometric Ba/Ti ratio leads to suppression of the high relative permittivity of the ferroelectric barium titanate [10]. In general, there are differences in behavior of barium titanate in the form of a sintered bulk material and a thin film [3,10]. BaTiO3 itself has seldom been plasma sprayed until now, and an understanding of its behavior

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in the form of a sprayed coating is not satisfactory. For coatings with the thickness of about 100 µm, values of relative permittivity 50 and loss factor 0.08 were reported [11]. Much higher relative permittivity (about 2500) was reported [2] more recently but this was for the room temperature and it was combined with a loss factor of about 0.45. The dielectric properties of the plasma sprayed BaTiO3 were related to the degree of crystallinity [11]. Coatings containing more crystalline material have higher relative permittivity. The relative permittivity was affected also by cracks and splat interfaces within the coating [11,12]. The reported value of relative permittivity of the plasma sprayed coating with a thickness of about 1 mm is higher [12] but even remarkably higher values are typical for a bulk BaTiO3 [13]. Additional heat treatment seems to be a promising method for microstructure restoration towards a bulk character. However, research on the effect of postspray treatment on the microstructure modification and dielectric response of plasma sprayed BaTiO3 films is very rare. Regarding mechanical properties of plasma sprayed BaTiO3 with a post-spray treatment; an even lower quantity of data is available. The goal of our thermal posttreatment experiments was to achieve coatings with stable relative permittivity and stable loss factor versus frequency of the electric field, while simultaneously maintaining sufficient mechanical quality. The focus of the present paper is on selected aspects of the dielectric as well as structural and mechanical characteristics of the annealed barium titanate coatings.

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2

Experimental

2.1 Material and specimen preparation BaTiO3 feedstock powder was obtained by crushing and sieving of sintered coarse agglomerates. These agglomerates were prepared by reactive sintering of micrometer-sized BaCO3 and TiO2 powders used as starting materials. The sintering was carried out in the company Keramicke kondenzatory, Hradec Kralove, Czech Republic. After the sieving of the feedstock the size distribution was between 20 and 63 µm with an average of 40 µm [12]. Gas-Stabilized Plasma gun (GSP) was used to perform Atmospheric Plasma Spraying (APS) of BaTiO3 [12] using the conventional direct current plasma gun F4 (Metco, Westburry, NY, USA). Carbon steel was used as a substrate. The samples were grit blasted with alumina particles to roughen the surface for an improvement of the adhesion strength. The plasma argon/hydrogen gas mixture was used with flow rates of 45/15 slm, and in certain cases 53/7 slm [12]. The powder was injected perpendicularly to the plasma jet axis with argon as a carrier gas (5 slm at 0.3 MPa) through an injector located downstream of the torch nozzle exit. Barium titanate was sprayed with the input power around 30 kW and the deposition time was about five minutes to reach a thickness of 0.9 to 1 mm. Our as-sprayed samples were identified via X-ray diffraction as partly amorphous [12]. To restore crystallinity, thermal posttreatment (annealing) was applied. The samples were heated in a laboratory furnace in the air atmosphere with heating and cooling rates 3 °C/min and with the dwell time of 30 min on the maximum temperature. Two temperatures were chosen: 500 °C in one case and 750 °C in another experiment with the goal to find the lowest temperature high enough for substantial changes in the microstructure, and especially in dielectric properties. The annealing was applied to the whole coating-substrate sandwich, profiting from the advantage that its thermal expansion is similar. The linear thermal expansion coefficient

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of BaTiO3 has an approximately linear run between 10.8x10−6 per °C at 350 °C to 17.5x10−6 per °C at 1050 °C [14]. The linear thermal expansion coefficient of carbon steel is between 10.4x10−6 per °C and 15.1x10−6 per °C according to common databases.

2. 2 Characterization techniques Polished cross sections of the coatings were prepared for microscopic observation and for microhardness measurement. For a better description of pores, additional criteria besides the percentual porosity were introduced: N.V. denotes the “Number of Voids per unit area“ of the cross section and Equivalent Diameter (E.D.) of voids represents their size distribution. All parameters were calculated for 10 images. Resolution of light microscopy used in this work is sufficient to provide quantification of the interlamellar pores but not the ultra-fine vertical cracks that are also common in the plasma sprayed coatings [15]. The coatings annealed for 30 min at 500 and 750 °C respectively were analyzed by powder Xray diffraction (PXRD) with CuKα radiation. We used a D8 Discover diffractometer (Bruker AXS, Germany) with parallel beam geometry and 1D detector. The obtained PXRD patterns were subjected to Rietveld refinement in order to ascertain the weight fraction of the identified phases, refine their lattice parameters and calculate the average values of coherently scattering domains or the so-called crystallite size. Moreover, the coating in as-sprayed and annealed (500°C) form was measured at elevated temperatures and these high temperature Xray diffraction (HTXRD) experiments were performed with CoKα radiation, X’Pert PRO MPD diffractometer in Anton Paar HT chamber. To detect the tetragonal BaTiO3 phase by Xray diffraction, the split of peaks of (002) and (200) reflection is a well-established indication [16-18].

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Microhardness of the coatings was measured with an optical microscope equipped with a Hanemann head and Vickers indenter using a 1 N load applied over 15 seconds. The mean value of microhardness was calculated as an average from 20 indentations. Electrical measurements were performed after removing the coatings from the metallic substrate. The surface of specimens was ground to eliminate surface roughness. Layers of aluminum as thin film electrodes were sputtered under reduced pressure on both sides of each sample. A three-electrode measurement fixture was used to evaluate dielectric parameters of the samples. The electric field was applied parallel to the spraying direction (i.e., perpendicular to the substrate surface). Capacity was measured in a frequency range from 80 Hz to 1 MHz using a programmable impedance analyzer (4284A, Agilent, USA). Applied voltage was 1V AC [12]. Relative permittivity εr was calculated from measured capacities CP and specimen dimensions since εr is directly proportional to CP according to the Eq. 1. CP = ε0 x εr x 1/k

Eq. 1

where ε0 = 8.854x10-12 F m-1; 1/k [m] is defined as the ratio between the guarded surface and the thickness of the sample. This technique is accurate enough (CP ± 0.01 pF) for dense materials (d > 95 % dth) [19]. The same arrangement and equipment was used for the loss tangent measurement at the same frequencies as capacity. Electric resistance was measured with a special resistivity adapter – Keithley model 6105. The electric field was applied from a regulated high-voltage source and the values read by a multi-purpose electrometer (617C, Keithley Instruments, USA). The applied voltage was 100 ± 2 V DC and the time of exposure 10 min. Volume resistivity was calculated from the measured resistance and specimen dimensions. The diffuse reflectance was measured by an UV-VIS-NIR scanning spectrophotometer (Shimadzu, Japan) with a multi-purpose large sample compartment. The reflectance curves

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obtained between 250 and 2000 nm were then converted to absorbance and recalculated [20] to the bandgap energy Ebg.

3

Results and discussion

Cross sections of individual feedstock particles showed that some of them were pore-free and crack-free, whereas others included pores [12]. According to XRD, the feedstock powder was composed solely from the tetragonal phase of barium titanate [12]. Density of the as-sprayed coating measured by pycnometry was 5365 ± 2 kg.m-3 [12]. The obtained PXRD patterns of the annealed coatings are juxtaposed in Fig. 1 with an indication of phase identification as well. The results of Rietveld refinement are summarized in Tab. 1, and Fig. 2 illustrates the correspondence between measured data (in blue) and a refined structural model (in red), which included tetragonal (CIF with ICSD code 94436 (Aoyagi, Kuroiwa, Sawada, Yamashita, & Atake, 2002)) and hexagonal (CIF with ICSD code 75240 (Akimoto, Gotoh, & Oosawa, 1994)) phases of BaTiO3 and also an amorphous material. Both patterns included the so-called amorphous halo in the region from 22 to 32 °2θ indicating the presence of an amorphous material in the irradiated volume. The calculated crystallinity, or amount of the crystalline material within the irradiated volume, was estimated by amorphous halo fitting by pseudo-Voigt function and with the presumption that the amorphous material has the same elemental composition as the crystalline part. As evident from Tab. 1, the tetragonality, i.e. the c/a ratio of BaTiO3, is very close to unity and, thus, implies Pm3 m cubic lattice instead of tetragonal P4mm. Indeed, when the cubic structure was assumed in Rietveld refinement, the difference between measured data and the structural model increased slightly only. For the HTXRD, the differences between Rwp factors for cubic and tetragonal lattice were virtually nonexistent. The diffraction profiles of

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(021) and (211) reflections of the cubic phase for as-sprayed and annealed coating measured at 100 and 500 °C are compared in Fig. 3. The content of the amorphous phase in the as-sprayed BaTiO3 was calculated earlier as 20 % [21]. During plasma spraying, because of the rapid cooling and low melting point of BaTiO3 (1650°C) compared to plasma jet temperatures, the formation of an amorphous phase is possible. As shown in Tab. 1, after heat treatment the amorphous material was changed into a crystalline one progressively with the annealing temperature but a totally crystalline sample was not obtained. Figure 4 shows the microstructure of the as-sprayed coating’s polished cross sections. In the thermal spray process particles first penetrated into the plasma jet, and then they transformed into a melted and semi-melted state because of the high temperature in the plasma jet. When the melted particles collided with the substrate, they cooled rapidly and made so-called splats. This process continued until the final coating was produced. Figure 4 shows quite a dense structure with a few cracks and predominantly rounded fine pores. Moreover, a certain variance in grayscale appeared corresponding to the shapes of individual splats. This is most probably associated with a subtle change in stiochiometry of Ti versus O [12] or also with the amorphous fraction. Figure 5 presents light micrographs of heat treated samples, where we can see individual splats in cross section. The structure is again rather dense; however, some vertical cracks are present. The variation in grayscale of individual splats signalizes incomplete restoration of stoichiometry for the 500°C annealing, but seemingly full restoration for the 750°C annealing. The main characteristics of porosity are summarized in Table 2. Generally, porosity produced in the structure of a thermal sprayed coating is divided into two categories: 1) porosity caused by gas entrapment through plasma spraying with a spherical shape [22]. 2) The other type of porosity has irregular and angular shapes, which are similar

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to casting contraction defects. These kinds of porosity can be created by formation of fragmented splats [22]. Also the residual stresses resulting from the rapid cooling of splats on impact, which caused curling of splat edges and therefore weak adhesion along the rim of each splat, could be another reason for porosity formation.

Figures 4 and 5 show that some of the defects start to sinter and heal. According to the image analysis results, the porosity was reduced by heat treatment from 4.7 % for the as-sprayed sample to 3.3 % for the 500°C annealing and finally to 1.5 % for 750°C. When such a porosity reduction happens, coating properties are usually improved by the heat treatment [23]. Heat treatment typically creates bridging of interfaces between splats in their immediate vicinity, and lead to a decrease in the quantity of pores and micro-cracks and an increase in the contact area between splats [24]. It should be mentioned; however, that the heat treated coatings cannot reach the theoretic density. Some inter-splat interfaces, porosity and large voids remain relatively unaffected by the heat treatment.

The N.V. parameter of pores (Tab. 2) is even higher than in the as-sprayed state after the 500°C annealing, but goes below the as-sprayed value after the 750°C annealing. The E.D. parameter of pores becomes larger with the increasing annealing temperature. A combination of these two parameters shows that the pores become larger and more frequent if the structural restoration starts at 500°C. With a further increase of the annealing temperature the pores are even larger but less frequent because they would coagulate [25] into individual objects, whereas small defects start to annihilate (therefore the E.D. parameter is even higher at 750°C).

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Mean Vickers microhardness at 1 N load is 6.6 ± 1.2 GPa for the as-sprayed, 6.0 ± 1.3 GPa after the 500°C annealing and 7.2 ± 1.5 GPa after the 750°C annealing. In combination with porosity, these results indicate that the 500°C annealing represents an unstable state where the microstructural changes worsen the mechanical properties.

Figure 6 shows the relative permittivity εr and loss factor tg δ measurement results. The dependence of permittivity on frequency exhibits a progressive decrease. This fact can be associated with the presence of pores, cracks and imperfections at the splat boundaries and grain boundaries; defects such as weak adhesion along the splat rim and at the inter-splat interfaces contribute to a cumulation of a freely movable charge. Dipoles created at these interfaces are active mainly at low frequencies. The permittivity values for the as-sprayed sample exhibit a pronounced growth at frequencies as low as 100 Hz. However, at frequencies between 1 kHz and 1 MHz the permittivity is under 400 and simultaneously the loss factor is under 0.4. As proved earlier [21], there was no sharp change of permittivity neither at any studied frequency, nor at any temperature. We consider that the amorphous fractions, as well as microstructural imperfections along the splat boundaries, are responsible for such a response of permittivity; i.e., without a clearly defined Curie temperature [21]. Generally, at a low frequency the relative permittivity of the ceramics decreased with increasing frequency. The relative permittivity of ceramics is attributed to four types of polarizations: interfacial, dipolar, atomic and electronic. At low frequencies, where all four types of polarization are involved, the relative permittivity and the dielectric loss are very high. The relative permittivity and the dielectric loss gradually decrease due to the increase in the frequency, which is due to the presence of only small-object polarizations (atomic and electronic) and the disappearance of the other two types. The interfacial polarization occurs up to frequencies of around 3 MHz. That is why the losses did not decrease below the value of

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about 0.2 at 1 MHz for the as-sprayed sample. With further frequency increase the relative permittivity, as well as dielectric loss, would be saturated because the electronic exchange cannot follow the AC field beyond a certain critical frequency [26].

After annealing, both the permittivity and the loss factor dropped and became much less frequency-dependent. The drop was more pronounced with the 500°C-annealed sample than with the 750°C-annealed one. That should be associated with progressive crystallization, (Tab. 1) as well as with stoichiometry restoration.

Electric DC resistivity grew from the order of 105 Ωm for the as-sprayed to as high as the order of 1010 or 1011 Ωm for the annealed samples. This indicated a strong decrease in a vacancy concentration after annealing. The change of the coating color induced by the annealing was also associated with this attribute, which was documented in Fig. 7. With the increasing temperature, the color of the coating surface, which had originally been dark due to the chemically reduced state, became brighter. At 750°C this process seemed to be finished because the coating was already a bit brighter than the bulk sintered sample (and had a similar color as the original powder – not shown). The compared bulk sample was sintered by heating in air at 1300°C for 2 h.

The diffuse reflectance, Fig. 8 (a) is depicted versus the BaSO4 standard that has a constant reflectivity equal to 100 % on the y-axis [27]. Below 350 nm (UV region) the reflectivity of all samples is similar, but above that wavelength it starts to differ; this difference is most pronounced in the infrared region. In the visible region the reflectance is lowest for the coating, medium for the sintered bulk sample and highest for the annealed coating. The

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coating has, due to plasma-induced instabilities in the bandgap, an enhanced chemical reactivity and becomes more oxidized than the sintered BaTiO3 sample. The bandgap evaluation [20] of the as-sprayed coating and one annealed coating (750°C) in comparison with a sintered bulk sample is shown in Fig. 8 (b). The bandgap estimation was done assuming a direct transition within the bandgap [20]. Both coating samples exhibited a blue shift versus the sintered sample. This shift was more markedly pronounced for the annealed coating. The dark blue color of the strongly reduced BaTiO3 (in our case: assprayed) was attributed to a small-polaron hopping of electrons of Ti3+–Ti4+ charge transfer type and gives evidence of the strongly disordered nature of the material [28]. The sintered BaTiO3 sample exhibited main absorption edge with a bandgap energy Ebg at 2.67 eV. The BaTiO3 coating exhibited main absorption edge with a larger bandgap energy Ebg at 2.72 eV and the annealed coating exhibited main absorption edge with a markedly higher bandgap energy Ebg at 3.11 eV. This blue shift can be attributed to oxidation of the originally reduced coating at heating in air and to a decrease in vacancy concentration after annealing. In connection with all plasma sprayed titanates, we can speak about plasma induced metastability that represents the existence of anomalous electronic states within the bandgap [21,25] and that is also responsible for the photocatalytic effects on the surface of the assprayed BaTiO3 [29].

Various papers deal with the effect of a hydrogen-rich atmosphere on BaTiO3 during its synthesis or subsequent annealing, as it was summarized in our earlier work [12]. An interstitial H atom was found to bind to one of the O atoms of BaTiO3 structure forming the OH group. Such a defect in the BaTiO3 structure contributes to a decrease in resistivity as well as a marked rise in the loss factor because the material starts to behave like an n-type semiconductor. This is because in such defective cells Ti4+ is reduced to Ti3+. Another

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question is why annealing cannot bring a completely bulk-like behavior of the thick film from the dielectric viewpoint, namely the appearance of ferroelectricity and piezoelectricity. The reason is that the plasma sprayed disordered microstructure (porosity, stress, cracks) as well as disordered lattice structure (phase mixture, anomalous electronic states within the bandgap) is so far from the thermal equilibrium that polarization is partly blocked and instead conductivity takes place due to weakly bounded electrons.

4

Conclusions

The goal of this study was to elaborate BaTiO3 dielectrics via plasma spraying and post-spray treatment of the coatings. We achieved nearly 1 mm thick coatings and investigated the main factors governing the differences between sintered bulk and annealed coatings. Among the major factors responsible for it we must include: i) not fully crystalline structure of our annealed coatings, ii) disordered microstructure with its defects as coagulated pores and short as well as longer cracks, and iii) presence of anomalous electronic states within the bandgap associated with an extraordinary sensitivity of the material to the atmosphere applied during the spraying and annealing. Our BaTiO3 prepared by plasma spraying plus subsequent annealing at 750°C for 30 min exhibited (at room temperature and RF range) relative permittivity between 150 and 200 and loss factor just under 0.1. Comparing the existing literature describing relevant material and technology combinations, our results represent an interesting compromise. We preclude that a higher annealing temperature and longer dwell time will lead to an even more pronounced pore coagulation and grain (crystallite) coarsening that will cause large problems in properties tailoring. Perhaps a pressure-assisted post-spray sintering of the thermally sprayed coatings can overcome such difficulties.

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Acknowledgment This work was supported by the Academy of Science of the Czech Republic under the project AV0 Z 20430508. H. Ageorges (Limoges University, France) should be acknowledged for her assistance with spraying.

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Table 1: Results of Rieveld refinement of PXRD pattern taken on the surface of the coatings annealed at 500 °C and 750 °C. Annealing tetragonal BaTiO3

hexagonal BaTiO3

Lattice

Crystallite Weight

parameters

size [nm]

fraction parameters

[Å] 500 °C

750 °C

a = 4.0092

Lattice

Cryst.

Weight

size [nm]

fraction

linity

[Å] D = 33±1

64 %

a = 5.727 ± D = 14±1 24 %

± 0.0004

0.004

c = 4.0338

c = 14.02 ±

± 0.0005

0.02

a = 4.0088

Crystal-

D = 27±1

75 %

a = 5.733 ± D = 13±1 20 %

± 0.0006

0.006 c =

c = 4.0349

14.04 ±

± 0.0008

0.03

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88 wt.%

95 wt.%

Table 2: Image analysis results.

Porosity

E.D. (pore

N.V., i.e.

Dark “phase”

Parameter

[%]

size) [µm]

pores per mm2

[%]

as-sprayed

4.7 ± 0.5

4.5

5840

16.3 ± 3.7

3.3 ± 0.9

5.6

8750

n.a.*

1.5 ± 0.7

5.8

4500

0

annealed 500 annealed 750

* Impossible to set the threshold correctly

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Table 3: Vickers microhardness HVm at 1 N load and electric resistivity.

HVm

Resistivity

Sample

[GPa]

[Ωm]

as-sprayed

6.6 ± 1.2

1.77 x 105

6.0 ± 1.3

7.82 x 1010

7.2 ± 1.5

2.91 x 1011

annealed 500 annealed 750

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Figure captions Fig. 1 – Juxtaposition of measured XRD patterns and phase identification of the observed reflections.

Fig. 2 – Rietveld refinement of PXRD pattern (Rwp = 9.64) with three phases being refined, i.e. hexagonal and tetragonal BaTiO3 and amorphous material whose broad diffraction profile lies in the region from 22 to 32 °2θ.

Fig. 3 – Diffraction profiles of (021) and (211) cubic BaTiO3 measured at 100 °C and 500 °C in as-sprayed coating and in the coating after annealing at 500 °C.

Fig. 4 – Polished cross section of the as-sprayed BaTiO3 coating (a) SEM-SE and light microscopy image (b).

Fig. 5 – Light microscopy views of the polished BaTiO3 coating: (a) annealed at 500°C and (b) annealed at 750°C. Fig. 6 – Dependence of permittivity (a) and loss tangent (b) on frequency.

Fig. 7 – Colors of the as-sprayed and annealed coatings; the sintered sample is inserted separately.

Fig. 8 – Reflectivity of the as-sprayed, annealed (at 750°C) and sintered sample (for comparison) (a); estimation of the bandgap of the same samples (b).

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Counts

Intensity, arb. units

□ ■ - hexagonal BaTiO3 □ - tetragonal/cubic BaTiO3

2000

1000

□

■ ■

□

□

□ □

■

□ 500 °C

□

750 °C

0 20

30

40

50

2θ, deg Fig. 1

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60

70

Fig. 2

- 24 -

Intensity, arb. units

Counts 15000

10000

(211)

▬ HTXRD at 100 °C ▬ HTXRD at 500 °C

(021) annealed coating at 500 °C

5000

as-sprayed coating

0 60

65

2θ, deg

Fig. 3

- 25 -

(a)

(b)

Fig. 4

- 26 -

(a)

(b)

Fig.5

- 27 -

(a)

(b)

Fig. 6

- 28 -

Fig. 7

- 29 -

(a)

(b)

Fig. 8

- 30 -