proceedings 2015

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| 11 Her

any, 9

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TECHNICAL UNIVERSITY OF KOŠICEFACULTY OF METALLURGY

h – 11th June 2015 Herľany, Slovakia

9th – 11th June 2015 Herľany, Slovakia 9th – 11th June 2015 Herľany, Slovakia

PREPARATION OF CERAMIC MATERIALS

PROCEEDINGS OF EDITED CONTRIBUTIONS

Slovak Silicate SocietySlovak Glass Society

2015 | 11ISBN 978-80-553-2122-6

TECHNICAL UNIVERSITY OF KOSICE

FACULTY OF METALLURGY,

DEPARTMENT OF CERAMICS

Institute of Geotechnics Slovak Academy of Sciences, Kosice

Slovak Silicate Society

Slovak Glass Society

PREPARATION OF CERAMIC MATERIALS Proceedings of the XI. International Conference

Herľany

9 th – 11 th June, 2015

Internet: Full text available online at http://web.tuke.sk/seminar_PKM/

Proceedings - XI

th International Conference

PREPARATION OF CERAMIC MATERIALS Herľany, 9

t h – 11

t h June, 2015

Editor: Organizing Committee of International Conference PREPARATION OF CERAMIC MATERIALS 2015

Faculty of Metallurgy, Department of Ceramics

Copyright Technical University of Kosice 2015 June 2015, I. edition , Number of Copies – 101; Number of Pages – 153

The publication of proceedings was supported by

SLOVAK SILICATE SOCIETY- ASSTS

SLOVAK GLASS SOCIETY

TECHNICAL UNIVERSITY OF KOSICE, FACULTY OF METALLUGY

Edited by:

Assoc Prof. RNDr František Lofaj, DrSc., Assoc. Prof. Ing. Miroslav Hnatko, Ph.D., Assoc. Prof. Ing. Beatrice Plešingerová, CSc., Assoc. Prof. Ing. Gabriel Sučik, Ph.D,, Ing. Diana Horkavcová, Ph.D., RNDr. Martin Fabián, Ph.D.,

ISBN: 978-80-553-2122-6

Acknowledgements:

Members of Organization Committee thank all Companies and Institutions:

SILICON a.s., Dobšiná

PROMAT, s.r.o.

SLOVALCO, a.s. Žiar nad Hronom

SLOVAL, s.r.o. Žiar nad Hronom

KERKOTHERM, a.s., Košice,

SMZ, a.s, Jelšava,

RONA a.s., Lednické Rovne,

for support of XI. International Conference


Herľany, 9t h

– 11th

June, 2015

CONTENT

Herľany, Slovakia 9

th-11

th June, 2015

3

LECTURES - research and engineering

L1 F. Lofaj, L. Kvetková, R. Podoba

THE OPTIMIZATION OF THE PROPERTIES OF HITUS W-C BASED COATINGS 8

L2 I. Petríková, M. Parchovianský, G. Barroso, G. Motz, D. Galusková, D. Galusek

POLYMER DERIVED GLASS CERAMIC LAYERS FOR CORROSION PROTECTION

OF METALS

13

L3 M. Fabián, J. Briančin

MECHANOSYNTHESIZED Ce1-xYxO2-δ SOLID SOLUTIONS AND THEIR SINTERING 18

L4 K. Haladejová, R. Klement, L. Dvorská, A. Prnová, D. Galusek

THE EFFECT OF THERMAL TREATMENT ON LUMINISCENCE PROPERTIES

OF Ce3+ DOPED GLASS AND GLASS-CERAMICS IN THE SYSTEM Y2O3Al2O3

23

L5 J. Drga, J. Chovanec., S. Šimkovič

DENSITY OF STATES AND MAGNETIC PROPERTIES OF Al2O3, Cr2O3 AND CrO2 28

L6 M. Kašiarová, P. Tatarko, Z. Chlup, I. Dlouhý

CERAMIC MATRIX REINFORCED BY BN NANOFILLERS 33

L7 A. Prnová, R. Klement, K. Bodišová, L. Hric, E. Neubauer, D. Galusek,

E. Bruneel, I. Van Driessche

STUDY OF THERMAL BEHAVIOUR AND HOT PRESSING OF ALUMINATE GLASS

MICROSPHERES

37

L8 A. Kovalčíková, Z. Lenčéš, J. Sedláček, P. Šajgalík, J. Dusza, M. Ignácz,

M. Mihaliková

OXIDATION RESISTANCE OF SiC CERAMICS PREPARED BY DIFFERENT

PROCEESSING ROUTES

40

L9 H. Bruncková, Ľ. Medvecký, P. Hvizdoš, J. Ďurišin, V. Girman

PYROCHLORE LANTHANUM NIOBATE PREPARED BY SOL-GEL METHOD

IN DIFFERENT SOLVENTS

44

L10 Š. Csáki, V. Trnovcová, J. Ondruška, I. Štubňa

INFLUENCE OF MILLING ON DC CONDUCTIVITY OF ILLITE 48

L11 A. Švančárková, D. Galusková, P. Gaalová, J. Balko, M. Fides, D. Galusek

EFFECT OF HEAT TREATMENT ON MECHANICAL PROPERTIES AND

CORROSION RESISTANCE OF LITHIUM DISILICATE DENTAL GLASS CERAMICS

55

L12 Z. Pramuková, M. Kašiarová, M. Precnerová, M. Hnatko, P. Šajgalík

MECHANICAL AND MICROSTRUCTURAL CHARACTERIZATION OF POROUS

SILICON NITRIDE BIOMATERIALS

59

L13 P. Gaalová, D. Galusková, A. Švančárková, J. Balko, D. Galusek

CORROSION OF NATURAL AND SYNTHETIC BIOMATERIALS IN ACIDIC MEDIA

AND ITS EFFECT ON MECHANICAL PROPERTIES

64

L14 A. Fedoročková, P. Raschman, G. Sučik

HIGH-GRADE SILICA PREPARED FROM SERPENTINITE TAILINGS USING

TWO-STAGE LEACHING AND PRECIPITATION

68

CONTENT

XIth

International Conference PREPARATION OF CERAMIC MATERIALS

4

L15 Ž. Dohnalová, P. Šulcová, N. Gorodylova

IMPACT OF WAY OF PREPARATION ON THE QUALITY OF CERAMIC PIGMENTS 73

L18 B. Plešingerová, N. Jádi, A. Doráková

THE REMOVAL OF HEAVY METALS FROM WATER BY NATURAL LIMESTONE

AND CONCRETE

86

L19 J. Bounziová, I. Nagyová, M. Černík

CASTABLES IN U.S.STEEL KOŠICE STEEL LADLES 91

LECTURES - industrial and education activities

L20 R. Hirjak, J. Bounziová, D. Chudíková, J. Parnahaj, M. Mingyár

APPLICATION OF CASTABLES PRODUCED BY RMS, a.s. KOŠICE

IN THE PUSHER FURNACE

97

L21 D. Hršak, A. Štrkalj, L. Slokar, Z. Glavaš

THE IMPACT OF INDUSTRY ON CLIMATE CHANGES FROM HOLISTIC

ENVIRONMENTALISM POINT OF VIEW

105

L22 K. Jesenák

COMMENTS ON THE CURRENT RELATIONSHIP BETWEEN GEOLOGY

AND CHEMISTRY AT SLOVAK UNIVERSITIES

109

POSTERS - Abstracts

P1 R. Štulajterová, Ľ. Medvecký, M. Giretová, T. Sopčák

PREPARATION AND CHARACTERIZATION BIOCERAMICS PREPARED FROM

TETRACALCIUM PHOSPHATE-NANOMONETITE CEMENT

112

P2 T. Sopčák, Ľ. Medvecký, R. Štulajterová, J. Ďurišin

INFLUENCE OF THE pH VALUE ON PHASE COMPOSITION AND MORPHOLOGY

OF CaO-SiO2-P2O5 BIOACTIVE GLASSES SYNTHESIZED BY SOL – GEL

PRECIPITATION

114

P3 D. Horkavcová, P. Novák, M. Černý, I. Fialová, E. Jablonská, Z. Zlámalová Cílová,

A. Helebrant

NEW TYPES OF TiSi ALLOYS WITH TITANIA AND SILVER SOL-GEL COATINGS

116

P4 J. Luxová, P. Šulcová

MALAYAITE PIGMENTS DOPED BY Fe 118

P5 K. Těšitelová, P. Šulcová

SYNTHESIS AND STUDY OF COMPOUNDS BASED ON Bi-Zn-Ce-Nb 120

P6 L. Pluhařová, L. Svoboda, P. Bělina

NONTRADITIONAL METHODS OF SYNTHESIS OF SPINEL COMPOUNDS 122

P7 P. Švančárek, R. Klement, L. Dvorská, D. Galusek

WILLEMITE BASED FLUORESCENT MATERIALS 125

CONTENT


th-11

th June, 2015

5

P8 D. Hršak, V. Bermanec, A. Štrkalj, L. Slokar, Z. Glavaš, D. Mašinović

THE POSSIBILITIES OF HYDROMETALURGICAL USING OF MANGANESE ORE

FROM BIHAĆ REGION

128

P9 R. Sedlák, A. Kovalčíková, Z. Pramuková, P. Rutkowsk, J. Dusza

MICROSTRUCTURE AND BASIC MECHANICAL PROPERTIES OF Si3N4 +

GRAPHENE PLATELETS COMPOSITES

131

P10 R. Klement, P. Švančárek, D. Galusek

THE EFFECT OF Mn2+ CONCENTRATION AND ZnO/SiO2 RATIO ON

LUMINESCENCE INTENZITY AND LUMINESCENCE DECAY IN GREEN EMITTING

PHOSPHOR Zn2SiO4:Mn2+

134

P11 L. Dvorská, K. Haladejová, R. Klement, A. Prnová, D. Galusek

EMISSION PROPERTIES OF Eu3+/Eu2+ PHOSPHORS IN THE SYSTEM Y2O3Al2O3 137

P12 D. Horkavcová, D. Rohanová, I. Březovská, P. Bozděchová, A. R. Boccaccini

EFFECT OF TRIS AND HEPES BUFFERS ON THE GLASS-CERAMIC SCAFFOLD 138

P13 I. Veverková, I. Lovětinská-Šlamborová, P. Exnar

INORGANIC-ORGANIC NANOFIBROUS MATERIALS FOR BIOMEDICAL

APPLICATIONS

140

P14 I. Danilová, I. Lovětinská-Šlamborová, P. Holý

IMMOBILIZATION OF BIOMOLECULES ON SILICA NANOFIBERS 142

P15 M. Fides, P. Hvizdoš, M. Novák

NANOINDENTATION OF (Ti,Ta) CARBONITRIDE CERMETS WITH VARIOUS

COBALT BASED BINDERS PREPARED BY MECHANOCHEMISTRY

145

P16 J. Urbánek, J. Macháček, J. Kutzendörfer, J. Hamáček

RHEOLOGICAL PROPERTIES OF CASTABLES 148

P17 M. Hubeňák, M. Kadlečíková, J. Breza, K. Jesenák, M. Kolmačka

NANOCOMPOSITE BASED ON CARBON NANOTUBES AND MONTMORILLONITE 150

XIth


6


th-11

th June, 2015

7

LECTURES

– research and engineering –

L1

XIth


8

THE OPTIMIZATION OF THE PROPERTIES OF HITUS W-C BASED

COATINGS

František Lofaj, Lenka Kvetková, Rudolf Podoba

Institute of Materials Research of SAS, Watsonova 47, 040 01 Košice

e-mail: [email protected]

The optimization of the high target utilization sputtering (HiTUS) made W-C based coatings

resulted in the increase of the hardness from around 20 GPa in the commercial coatings of

this type up to 33 GPa. HiTUS parameters offer possibilities for the control of the C/W ratio

and subsequent control of the resulting properties.

Keywords: HiTUS, W-C coatings, nanohardness

Introduction

The W-C based coatings are a part of metal doped diamond-like carbon coatings

covering wide range of properties depending on their composition and techniques of

deposition. The hardness up to 40 GPa is obtained in the case of small contents of

free C and decreases to <10 GPa, when free C prevails [1]. The coefficient of friction

(COF) exhibits opposite behavior – it is around 0.8 in low C content coatings and

decreases to ~0.1 when the amount of free C is higher than 50 % [2]. W-C coatings

are usually prepared by conventional DC and/or RF magnetron sputtering [1-3] but

recently, the first attempts of high power impulse magnetron sputtering (HiPIMS)

have been reported [4]. The latest development in the ionized sputtering

technologies patented by PQL Limited, UK, in 2005 is „High Target Utilization

Sputtering“ (HiTUS) [5]. In HiTUS, high density plasma is generated in a remote RF

plasma source and it is extracted into the chamber over the target area. The target

is negatively biased and its sputtering occurs in a conventional manner. HiTUS is

suitable for the deposition of single-phase or multi-component coatings, including

W-C coatings. Therefore, the aim of the work is to investigate the possibilities of

HiTUS technology in the optimization of hardness in W-C coatings.

Experimental Procedure

The HiTUS W-C based coatings were deposited on the polished bearing steel (STN

14109) substrates with the diameter of 25 mm in a sputtering system S 500 (Plasma

Quest Ltd., United Kingdom) from WC target (76.2 mm diameter). Thin Ti interlayer

between the substrate and coating has been deposited after plasma cleaning of the

substrate to improve adhesion. The principal variables included Ar flow (working

pressure), flow of additional C2H2 and RF power (bias) on the target. All the

parameters but one was kept constant to obtain the corresponding dependencies.

The deposition time was 1.3 h which resulted in the thickness of around 1.2 µm.

The structure and composition of the studied W-C coatings have been observed by

FESEM (Auriga Compact, Zeiss) with EDS system and Raman microscope

mailto:[email protected]

L1 THE OPTIMIZATION OF THE PROPERTIES OF HITUS W-C BASED COATINGS

F. Lofaj, L. Kvetková, R. Podoba


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9

2

WC1-x 20 -1316

Fe

Fe (substrate)

(XploRA, Horiba Yvon Jobin). Their thickness was measured by optical

interferometer (Sensofar Neox Plu, Spain) and SEM. The composition depth profiles

were obtained from GDOES measurements (GD Profiler 2, Horiba Yvon Jobin). The

instrumented hardness of the coatings was measured on nanoindenter (G200,

Agilent in constant strain rate (0.05 s-1

) regime and sinusoidal mode to obtain

hardness depth profiles. The resulting hardness is an average from 25 indents made

on each coating.

Results and discussion

The studied W-C coatings have been deposited in two series: at constant bias (50 W

RF power) to investigate the effect of acetylene content; at three levels of bias for

different acetylene contents and Si or Ti sublayers to improve adhesion. SEM

observations of the coating fractured cross sections showed featureless structure

(Fig. 1a) and X-ray diffraction (Fig. 1b) revealed only wide peaks suggesting

amorphous or nanocrystalline structure corresponding to WC1-x phase (PDF 20 -

1316). Limited HRTEM observation confirmed only amorphous structure up to now.

Fig. 1: a – The structure of W-C coating on the fracture cross section; b – X-ray diffraction indicating the presence of WC1-x phase.

Fig. 2a illustrates the GDOES depth profiles of W and C and Fig. 2b their summary as

a function of acetylene content. The C/W ratio of the atomic concentrations

increases from 0.7 until 1.5 with the increase of C2H2 content from 0 vol. % to

4.0 vol. %. Obviously, the addition of acetylene is a way to control the composition

(and properties) of W-C coatings. When stoichiometric WC formation is assumed,

the ratio C/W just slightly above 1.0 would be required to achieve nanocomposite

structure with high hardness. However, only WC1-x formed (see Fig. 1b), therefore the

optimum C/W ratio depends on x value. Total content of carbon reached the

maximum at 3.2 at.% of C2H2 also at 50 and 100 W bias power. It may suggest

limited ability to form WC1-x and possibility for the excess of carbon. The presence

of free carbon can be detected by Raman spectroscopy. Indeed, Raman spectra

L1

XIth


10

contained D and G peaks corresponding to sp2 bonds of disordered carbon (Fig. 3a).

In agreement with GDOES measurements, their intensities were low in the coatings

deposited with 0 and 0.8 vol.% of C2H2, which indicated minor content of free C.

Moreover, the tendency of a shift of G peak toward higher frequency was observed

as a function of C2H2 content and RF power bias (Fig. 3b).

0 200 400 600 800 1000 1200 1400 1600 1800 2000

0

10

20

30

40

50

60

Co

nc

en

tra

tio

n (

at.

%)

Depth, [nm]

W-C coating substrate

steel 14109C

W

a) b)

Fig. 2: a - GDOES depth profiles of C and W in HiTUS W-C coating; b - C and W concentrations as a function of C2H2 at constant Ar flow (120 sccm) in HiTUS W-C coatings.

a) b)

Fig. 3: The intensities of D and G peaks of free C in HiTUS W-C coatings as a function of C2H2 flow – a; the shift of D and G peak positions for different acetylene flows and RF power applied to substrate (bias) - b.

steel substrate

Ti sublayer

0 W (bias)

600 800 1000 1200 1400 1600 1800 2000

0

50

100

150

200

250

300

350

G

4 % C2H

2

3,2 % C2H

2

2,4 % C2H

2

0,8 % C2H

2

0 % C2H

2

Ra

ma

n In

ten

sity

Raman shift (cm-1)

D

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

1330 1340 1350 1360 1540 1550 1560 1570

0

20

40

60

80

100

1330 1340 1350 1360 1540 1550 1560 1570

0

20

40

60

80

100

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

1330 1340 1350 1360 1540 1550 1560 1570

0

20

40

60

80

100

1330 1340 1350 1360 1540 1550 1560 1570

0

20

40

60

80

100

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

1330 1340 1350 1360 1540 1550 1560 1570

-10

0

10

20

30

40

50

60

70

80

90

100

110

Su

bstr

ate

RF

po

we

r -

Bia

s (

W)

0 %

0,8%

2,4%

3,2%

4%

G peak (cm-1)D peak (cm

-1)

0 1 2 3 4

35

40

45

50

55

60

40

42

44

46

48

50

52

54

56

58

60

C [a

t. %]

W c

on

ce

ntr

atio

n [a

t. %

]

C2H

2 flow [ vol. %]

0 W power bias

CW

L1 THE OPTIMIZATION OF THE PROPERTIES OF HITUS W-C BASED COATINGS

F. Lofaj, L. Kvetková, R. Podoba


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11

Such shift indicated an increase of ordering of the initially disordered sp2 carbon

toward graphitic carbon [1]. The coating hardnesses at different acetylene contents are compared for different bias powers in Fig. 4. Without acetylene, hardness increased from ~27 GPa to ~33 GPa with the increase of RF power bias. Addition of just 0.8 % of acetylene into Ar flow resulted in the shift of hardness by 1.5 – 3.0 GPa toward higher values. This seemed to be attributed to the formation of WC1-x with smaller x value, i.e. closer to hard stoichiometric WC. However, 2.4 % acetylene addition caused hardness degradation close or even below 20 GPa which was related to the increased content of soft free graphitic carbon indicated by Raman and GDOES measurements. RF power applied to the substrate slightly improves the hardness values at zero and 0.8 vol.% of C2H2 but it results in its strong degradation at higher contents of acetylene and higher (100 W) power.

0 25 50 75 100

15

20

25

30

35

Ti sublayer

5 sccm C2H

2 - delaminated at 100 W

4 sccm C2H

2

3 sccm C2H

2

1 sccm C2H

2

0 sccm C2H

2

Na

no

ha

rdn

ess, H

IT [G

Pa

]

Substrate RF power - Bias [W] a) b)

Fig. 4: Nanohardness, Hit as a function of substrate RF power (bias) – a) and acetylene flow – b) in the HiTUS W-C coatings.

Indentation moduls exhibited similar behavior with the highest indentation modulus

values of 310 – 350 GPa and degradation at higher acetylene contents and RF

power.

Conclusions

HiTUS is suitable for the reactive deposition of W-C based coatings. The addition of

acetylene and RF power applied to the substrate provide a possibility to control the

amount of free carbon, the level of its ordering and subsequently, also the resulting

hardness of the coatings. Depending on the conditions, the W-C coatings with low

amount of free carbon (C/W ≤ 1.0) exhibit higher hardness (up to 33 GPa) whereas

the increase of the content of C (C/W > 1.0) results in significant reduction of the

hardness and indentation modulus.

0 1 2 3 4 5

15

20

25

30

35

Ti sublayer

Na

no

ha

rdn

ess, H

IT [G

Pa

]

C2H

2 flow [vol. %]

###

###

###

0 0,8 1,6 2,4 3,2 4

0 W Bias

50 W Bias

100W Bias

L1

XIth


12

Acknowledgment

The contribution of M. Ferdinandy with technical help is greatly appreciated. The financial support provided by the projects APVV-0520-10, VEGA 2/0098/14 and VEGA 2/0187/15 is acknowledged.

References

[1] S. El Mrabet, M.D. Abad, J.C. Sánchez-López, Surf. Coat. Technol., 206, 1913–1920 (2011). [2] M.D. Abad, M.A. Muñoz-Márquez, S. El Mrabet, et al., Surf. Coat. Technol., 204, 3490 (2010). [3] A.A. Voevodin, J.P. O'Neill, S.V. Prasad, J.S. Zabinski, J. Vac. Sci. Technol., A17, 986 (1999). [4] F. Lofaj, P. Hviščová, L. Kvetková, Proc. Vrstvy a povlaky 2014, Trenčín, Knihviazačstvo, 141-146 (2014).

ISBN 978-80-970824-3-7 [5] http://www.plasma-quest.com/attachments/article/115/BACKGROUNDTO_THE_OPERATION_OF_THE_PQL

_PLASMA_TECHNOLOGY.pdf

http://www.plasma-quest.com/attachments/article/115/BACKGROUNDTO_THE_OPERATION_OF_THE_PQL

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13

POLYMER DERIVED GLASS CERAMIC LAYERS FOR CORROSION

PROTECTION OF METALS

Ivana Petríková1, Milan Parchovianský

1, Gilvan Barroso

2, Gunter Motz

2,

Dagmar Galusková1, Dušan Galusek

1

1Vitrum Laugaricio – Joint Glass Center of the IIC SAS, TnU AD, and FCHFT STU, Študentská 2,

911 50 Trenčín, Slovakia 2University of Bayreuth, Ceramic Materials Engineering (CME), D-95440 Bayreuth, Germany


The work was aimed at development of a relatively thick, protective, dense and well adherent

coating system on steel. For that purpose, a double layer coating consisting of a polymer

derived ceramic (PDC) bond coat, and a PDC top coat with glass and ceramic fillers were

prepared. The phase composition and microstructure of the prepared glass microspheres and

coatings was investigated by X-ray diffraction studies and SEM.

Keywords: glass microspheres, polymer derived ceramics, coatings

Introduction

Environmental barrier coatings are widely applied to enhance the resistance of

various materials against aggressive environments. Non-oxide and oxide ceramic

coatings are promising candidates for improving corrosion and oxidation resistance

of metals. The coatings are mainly based on silicon containing precursors like

polysilazanes. A novel environmental barrier coating system for oxidation/corrosion

protection of steel consisting of a perhydropolysilazane (PHPS) derived bond coat

and a polysilazane-based glass/ceramic composite top coat has been under

investigation.In order to increase the application temperature of the coatings, which

is usually limited by the glass melting temperature of the used glass frit, special

aluminosilicate glasses (SiO2-Al2O3-ZrO2) with high melting temperature were

prepared in the form of microspheres by flame synthesis. The addition of fillers was

aimed at elimination of shrinkage in the course of polymer-to-ceramic conversion,

which could result in crack formation and delamination of the polymer derived

protective coating from the metallic substrate.

Experimental part

Two commercially available polysilazanes PHPS (perhydropolysilazane) and

Durazane 1800 (both AZ Electronic Materials GmbH, Germany) were used as

precursor materials. The glass microspheres were prepared by combination of a

modified Pechini sol-gel method and flame synthesis. Precursor powders for the

flame synthesis of glass microspheres were prepared by mixing of suitable amounts

of aluminium nitrate dissolved in distilled water with zirconium oxychloride also

dissolved in distilled water. Citric acid and ethylene glycol were added to the mixture

and heated in oil bath for 2 h at the temperature 85 - 90 °C. Simultaneously, a SiO2


POLYMER DERIVED GLASS CERAMIC LAYERS FOR CORROSION PROTECTION OF METALS L2

I. Petríková, M. Parchovianský, G. Barroso, G. Motz, D. Galusková, D. Galusek

XIth


14

sol was prepared from TEOS. TEOS was mixed with pure ethanol and then

deionized water together with hydrogen chloride were added dropwise. The silica sol

was stirred for 2 h at room temperature. Finally the SiO2 sol was added into the

Al2O3-ZrO2 sol, and the solvent was evaporated under continuous stirring. The

product was dried, calcined at 800 °C for 6 h to remove organic residua, and finally

sieved through a 42 µm mesh screen. The synthesized precursor powders were

used for the preparation of glass microspheres by the flame synthesis technique.

Five compositions with different SiO2-Al2O3-ZrO2 weight ratios were prepared (Tab.

1). The precursor powders were fed into methane-oxygen flame using methane as a

carrier gas. Molten particles were quenched by spraying them with deionized water.

The glass microspheres were collected in a container, sedimented, separated and

dried.

Flat stainless steel (AISI 441) plates were cut into sheets with the dimensions of

4.5 cm × 4.5 cm, ultrasonically cleaned in acetone and dried. The bond coat was

prepared from a pre-ceramic polymer precursor (polysilazane PHPS, AZ Electronic

Materials GmbH) that forms a ceramic layer after heat treating in air. The steel

sample substrates were dip-coated with a hoisting apparatus and the coating

thickness was adjusted by concentration of the solution and variation of the hoisting

speed. Annealing of the bond coated sheets was performed in air at the temperature

of 500 °C for 1 h with heating and cooling rates 3 K/min. The top coat was prepared

by mixing defined volume fractions of ceramic filler particles (YSZ, ZrSi2), glass

microspheres (SiO2-Al2O3-ZrO2) of various compositions and a liquid polysilazane

Durazane 1800. The compositions of top coats are listed in the Table 2. The YSZ,

ZrSi2 and glass microspheres were separately dispersed in a solution of di-n-

butylether and a dispersant (DISPERBYK, Wesel, Germany), followed by 30 min

ultrasonic treatment and 48 h stirring. Then the liquid polysilazane Durazane 1800

was added to the mixture, and the resulting mixture was applied by spray coating

technique. The pyrolysis of the composite coatings was performed in air (electric

oven Nabertherm® LH 60/14, Nabertherm,Germany) at the temperature 1000 °C

with the heating and cooling rates of 3 K/min and a holding time of 1 h.

Table 1: The compositions of glass microspheres

Compositions SiO2 (wt.%) Al2O3 (wt.%) ZrO2 (wt.%)

SAZ_M_80-10-10 80 10 10

SAZ_M_60-30-10 60 30 10

SAZ_M_50-20-30 50 20 30

SAZ_M_50-30-20 50 30 20

SAZ_M_50-40-10 50 40 10

Table 2: The composition of top coats

YSZ (vol. %) ZrSi2 (vol. %) Durazane 1800 (vol. %) Glass (vol. %)

64.8 5.4 26.3 3.4

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The amorphous nature of prepared glass microspheres was confirmed by X-ray

powder diffraction (Fig. 1). No diffraction peaks indicating the presence of crystalline

phases were present. The results of DSC analysis of glass microspheres are shown

in the Fig. 2. The glass transition temperature of prepared glasses lies in the interval

between 850 and 920 °C, the onset of crystallization temperatures range between

915 and 990 °C, and the temperatures of the maxima of exothermic peaks are in the

range from 940 to 1020 °C. The observed differences were attributed to various

compositions of studied glasses, especially different contents of silica as the

component, which was expected to increase the stability of glasses in terms of their

resistance against crystallization.

10 20 30 40 50 60 70 80

2

SAZ_M_50-20-30

SAZ_M_50-30-20

SAZ_M_50-40-10

SAZ_M_60-30-10

SAZ_M_80-10-10

0 200 400 600 800 1000 1200

SAZ_M_50-20-30

SAZ_M_50-30-20

SAZ_M_50-40-10

SAZ_M_60-30-10

T/°C

SAZ_M_80-10-10

Fig. 1: XRD patterns of the glass microspheres Fig. 2:The results of DSC analysis of glass

microspheres

The SEM examination of prepared glasses revealed spherical particles (Fig: 3). More

detailed inspection by SEM revealed that the microspheres had smooth surfaces,

which indicate that the temperature of methane-oxygen flame was sufficient for

complete melting of precursor powders, as confirmed also by the results of X-ray

diffraction. No un-melted residua of precursor powder were found. The results of

SEM confirmed the results of particle size analysis, showing the microspheres with

diameters ranging from 1 to 30 μm. The major fraction contained the microspheres

with the diameters within the interval 1 – 5 μm and 5 – 10 μm, and only small part

accounted for the microspheres with larger diameters, i.e. from 10 to 30 μm.

The thickness of the coatings of various compositions is shown in the Fig. 5. The

coating thickness was adjusted between 10 and 20 μm. The results of thickness

measurements confirm that relatively thick coating systems were obtained. The

phase composition and microstructure of the coatings was also investigated by X-

ray diffraction studies and SEM. XRD patterns of coatings are displayed in the Fig. 6.

The glass filled polysilazane-based coatings are crystalline after pyrolysis in air at

POLYMER DERIVED GLASS CERAMIC LAYERS FOR CORROSION PROTECTION OF METALS L2

I. Petríková, M. Parchovianský, G. Barroso, G. Motz, D. Galusková, D. Galusek

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16

the temperature 1000 °C. For the glass and zirconia filler coatings, two crystalline

phases were detected in the coatings, namely monoclinic and cubic ZrO2. The

phases created by crystallization of glass fillers or the Durazane 1800-derived

phases were below the detection limit of the X-ray diffraction.

Fig: 3: SEM micrograph of glass microspheres

(SAZ_M_60-30-10)

Fig. 4: The SEM micrographs of the C60-

30-10 top coat

5

10

15

20

25

Th

ickn

ess/

m

C811 C631 C541 C532 C523

0 10 20 30 40 50 60 70 80 90

ZrO2 monoclinic

ZrO2 cubic

C 50-20-30

C 50-30-20

C 50-40-10

C60-30-10

2

C80-10-10

2

Fig. 5: The average thickness of coatings of various composition

Fig. 6: XRD patterns of the coatings

Fig. 4 shows the SEM micrograph of a C 60-30-10 coating on stainless steel: the

SEM examination was focused on evaluation of homogeneity, adhesion and

possible failures of the coatings. The composite coatings were uniform and well

adherent. ZrO2 particles were distributed homogeneously in the matrix. The PHPS-

based bond coat increased the adhesion of top coatings and acted as a diffusion

barrier against oxidation during the pyrolysis of the coating system. However, the

composite coatings were not fully dense, contained small closed pores with a

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17

diameter up to 100 nm. In some cases, cracks and occasionally also delamination

were detected both within and at the interface of the approximately 10 μm thick

coatings, due to high volume shrinkage of the precursor during pyrolysis.

Conclusion

In this work, a double layer polysilazane-based environmental barrier coating system

for steel was prepared by simple dip- and spray-coating techniques. Due to the

reactivity of the PHPS, direct chemical bonds between the substrate and the coating

were formed. The results of SEM analysis confirm that the combination of PDCs with

glass and ceramic fillers enable the processing of relatively dense, uniform and well

adherent composite coating system with thickness between 10 and 20 μm.

Acknowledgments

Financial support of this work by the grant VEGA 2/0058/14, and the Alexander von Humboldt Foundation in the frame of the institutional cooperation grant scheme is gratefully acknowledged. This publication was created in the frame of the project "Centre of excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056, based on the Operational Program Research and Development funded from the European Regional Development Fund.

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MECHANOSYNTHESIZED Ce1-xYxO2-δ SOLID SOLUTIONS

AND THEIR SINTERING

Martin Fabián, Jaroslav Briančin

Institute of Geotechnics, Slovak Academy of Sciences, Košice, Slovakia


One-step mechanochemical approach was used to synthesize nanocrystalline Ce1-xYxO2-δ

(x ≤ 0.35) solid solutions. X-ray powder diffraction (XRPD) was utilized to characterize

crystalline phases of as-prepared products. It was shown that the presence of yttrium

suppresses the crystal growth. The effect of yttrium content on the grain size before and after

sintering of green bodies at high temperatures was evaluated by means of scanning electron

microscopy (SEM). The combination of one-step mechanochemical process and relatively low

sintering temperatures allowed for preparation of ceria-based ceramics with wide scale of

applications.

Keywords: ceria, yttrium, solid solutions, mechanosynthesis, sintering

Introduction

It is well known, that CeO2 has been extensively investigated in various applications

including catalysators, fuel cells, polishing materials, ultraviolet absorbers,

phosphors, oxygen sensors, etc. [1]. The CeO2 possess a cubic fluorite structure

(S.G. F m 3 m) which is known to tolerate a considerable reduction in size without

phase change especially at high temperatures. An aliovalent doping makes the

CeO2 promising oxygen-conductive electrolyte for intermediate-temperature

SOFC´s. It is generally accepted that Gd3+

or Sm3+

doped ceria exhibits the highest

conductivity due to the small association enthalpy between dopant cation and

oxygen vacancy in the fluorite lattice [2]. However, relatively high cost of these

elements faces difficulties for their commercialization and thus environmentally more

friendly with comparatively cheaper price elements present promising dopants for

CeO2 based SOFC´s. The synthesis of nanosized doped ceria powders is of

immense importance to get dense sintered product at a lower temperature. This is

because the CeO2 based materials prepared by the conventional solid-state

synthesis are characterized by difficult density below 1600 °C [3]. Many methods

have been used to synthesize nanocrystalline CeO2 doped powders [see e.g. 4, 5

and references therein]. Most of these methods involve complicated steps and

expensive raw materials to fabricate nanosized ceria-based ceramic powders. On

the other hand, relatively simple one-step ball milling towards preparation of

aliovalent doped CeO2 presents an alternative method to produce lower-cost

powders in the nanometer range [6].

In this work, Ce1-xYxO2-δ (x= 0 – 0.35) solid solutions have been prepared via on step

mechanochemical route. The main purpose in this case was to evaluate the effect of

L3 MECHANOSYNTHESIZED Ce1-xYxO2-δ SOLID SOLUTIONS AND THEIR SINTERING

M. Fabián, J. Briančin


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preparation method and dopant concentration on properties and morphology of

sintered specimens.

Experimental methods

The solid precursors, cerium oxide (CeO2, 99.9 % purity; Aldrich) and yttrium oxide

(Y2O3, 99.99 % purity; Aldrich) were used for the mechanosynthesis of Ce1-xYxO2-δ.

5 g of the (1–x)CeO2 + (x/2)Y2O3 mixtures (x = 0.1 – 0.35) were milled for various

times (up to 90 min) in a high-energy planetary ball mill Pulverisette 7 Premium line

(Fritsch). A grinding chamber (80 cm3 in volume) and balls (10 mm in diameter)

made of tungsten carbide were used. The ball-to-powder weight ratio was 40:1.

Milling experiments were performed in ambient atmosphere at 600 rpm. For

sintering study, as-prepared powders were compacted in a uniaxial hydraulic press

at 200 MPa containing 5 wt.% of cellulose. The green pellets were sintered at

1250 °C for 4 hours under static air.

The X-ray powder diffraction (XRPD) patterns were collected using an D8 Advance

diffractometer (Brucker, Germany) (for phase evolution analysis) and using an

X´pert MPD (Phillips, Netherlands) (for sintered samples) both with the Cu Kα1,2

wavelength in the Bragg-Brentano configuration. The generators were set-up at

40 kV and 40 mA and patterns were collected in the range of 20 – 85° 2θ with a step

of 0.05° 2θ and a time per each step of 5 s. The microstructural analysis of the as-

prepared as well as sintered samples was carried out by field emission-scanning

electron microscope FE-SEM, Mira 3 (Tescan, Czech Republic) coupled with EDS

analyzer (Oxford Instruments, UK). PCCS measurements were done using

Nanophox (Sympatec, Germany).


The X-ray diffraction patterns of the Ce1-xYxO2-δ (x= 0 - 0.3) solid solutions prepared

via 90 minutes of high-energy ball milling are shown in Fig. 1. It can be seen that

effective incorporation of yttrium into the fluorite-type structure of ceria results in

continuous shift of all characteristic XRPD reflections to the higher angular positions

(lower values of d-spacing), indicating a lattice contraction of Ce1-xYxO2-δ with the Y3+

doping and incorporation of Y3+

in the CeO2 structure. The broadening of XRPD

peaks reflects the nanocrystalline character (decrease in crystalline size) and

increase in lattice strain. This is confirmed by results obtained by profile analysis of

diffraction peaks (see Tab. 1) utilizing TCH-pV function comprising the size and strain

calculations based on integral breadths [7].

Based on the calculation of crystallites size it can be concluded that the presence of

yttrium in the ceria samples suppresses their crystallization (crystallite growth); i.e.,

the average crystallite size of mechanosynthesized Ce1-xYxO2-δ (x = 0.1 – 0.35)

ranges from ~ 14 to ~17 nm – see Tab. 1. This result is in accordance with the

findings on Sm3+

doped CeO2-δ, see, e.g., Ref. [8] and citations therein as well as

with recent work on Eu3+

doped CeO2-δ system [9]. The observed phenomenon can

be partly explained by the absence of oxygen (by the presence of oxygen

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20

vacancies) in the as-prepared Ce1-xYxO2-δ causing the breaking of bonds between

the atomic planes and, thus, resulting in a relatively small grain size, as it has been

suggested by Colis et al. [10]. However, relatively big agglomerates resulted from

high-energy ball milling with particles size distribution in the range of ~250 - 300 nm.

The tendency to form agglomerates as well as homogeneous distribution of

particular elements is further confirmed in Fig. 1. These results are in the good

accordance with PCCS analysis summarized in Tab. 1. Moreover, it can be clearly

seen that all investigated elements are uniformly distributed over the as-prepared

sample confirming the homogeneity of the solid solution.

Fig. 1: (left) The XRPD pattern of Ce1-xYxO2-δ (x= 0 - 0.35) solid solutions synthesized via high-energy ball milling; (right) The particles morphology and elements distribution in mechanosynthesized Ce0.7Y0.3O2-δ.

In order to obtain the nanocrystalline ceramic powder dense sample, the

mechanosynthesized nanopowders were sintered at 1250 °C for 4 h in atmospheric

conditions. The Ce1-xYxO2-δ diffraction patterns are essentially unchanged implying

that no parasitic phases are present and no phase related transformation had

occurred during high-thermal treatment.

Table 1: The effect of yttrium content on the crystallites and particles size.

Y3+ content Crystallite size/nm (XRPD) Particles size/nm (PCCS)

0.0 21 ̶

0.1 17 280

0.2 17 300

0.3 14 260

0.35 17 283

L3 MECHANOSYNTHESIZED Ce1-xYxO2-δ SOLID SOLUTIONS AND THEIR SINTERING

M. Fabián, J. Briančin


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21

Fig. 2: (left) The XRPD pattern of sintered Ce1-xYxO2-δ (x = 0.1 - 0.3) solid solutions; (right) The SEM micrographs of Ce1-xYxO2-δ pellets sintered for 4 h at 1250 °C; (a) x = 0, (b) x = 0.1 and (c) x = 0.3.

The SEM micrographs of three sintered samples with nominal composition of CeO2,

Ce0.9O0.1O2- δ and Ce0.7O0.3O2- δ, are also shown in Fig. 2 respectively. It was found,

that the concentration of small closed pores present in the sample increases with

increase of yttrium content.

As it was reported [11], the grain size is very important for the crystal structure

because at high operating SOFC temperatures high grain sizes result in cracks. The

nano-sized fine-grained oxide structure is advantageous. This is so because the

diffusion can easily take place and release the accumulated stresses at high

temperatures for large grain sized crystalline ceramic structures [12]. On the other

hand, it was found that the highest values of electronic conduction are reached with

dopant concentration up to 10 mol.% [13]. It also implies that the conductivity is

influenced by presence of small pores in the ceria based solid solutions.

Conclusions

The one-step mechanosynthesis has been successfully employed to prepare

Ce1-xYxO2-δ (x = 0.1-0.35) solid solutions with a crystallite average size below 20 nm.

Yttrium incorporation into the F-type ceria structure was confirmed by XRPD. The

effect of yttrium concentration on structural and microstructural parameters of

sintered samples was investigated by SEM. It was found, that increase in yttrium

content results in formation of small pores distributed in solid solution.

Acknowledgement

The financial support by the Slovak Research Grant Agency (project VEGA 2/0097/13) is

gratefully acknowledged.

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References

[1] X. Song, N. Juang, Y. Li, D. Xu, G. Qiu, J. Rare Earth. 25, 428 (2007). [2] G. Laukaitis, J. Dudonis, J. Alloy. Compd. 459, 320 (2008). [3] H. Xu, H. Yan, Z. Chen, J. Power Sources 163, 409 (2006). [4] A. Arabaci, M.F. Öksüzömer, M.F., Ceram. Int. 38, 6509 (2012). [5] C. Goulart, E. Djurado, J. Eur. Ceram. Soc. 33, 769 (2013). [6] P. Baláž, M. Achimovičová, M. Baláž, P. Billik, et al., Chem. Soc. Rev. 42, 7571 (2013). [7] J. Rodriguez-Carvajal, Physica B 192, 55 (1993). [8] L. Živković, V. Lair, O. Lupan, A. Ringuedé, Thin Solid Films 519, 3538 (2011). [9] A. Kremenović, D. K. Bozanić, A. M. Welsch, B. Jancar, et al., J. Nanosci. Nanotechnol. 12, 8893 (2012). [10] S. Colis, A. Bouaine, G. Schmerber, C. Uhlaq-Bouillet, et al., Phys. Chem. Chem. Phys. 14, 7256 (2012). [11] I. Uslu, A. Aytimur, M.K. Öztürk, S. Koçyiğit, Ceram. Int. 38, 4943 (2012). [12] Z.Y. Liu, W. Gao, K. Dahm, F.H. Wang, Scr. Mater. 37, 1551 (1997).

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THE EFFECT OF THERMAL TREATMENT ON LUMINISCENCE

PROPERTIES OF Ce3+ DOPED GLASS AND GLASS-CERAMICS

IN THE SYSTEM Y2O3Al2O3

Katarína Haladejová, Robert Klement, Lívia Dvorská, Anna Prnová,

Dušan Galusek

Vitrum Laugaricio – Joint Glass Centre of the IIC SAS, TnU AD and FChPT STU, Trenčín, Slovakia

e-mail: [email protected], [email protected]

The Ce3+

doped glass, Y40A60Ce0.5, with composition of 76.67 mol.% (60 wt.%)

Al2O3, 23.08 mol.% (40 wt.%) Y2O3 and 0.25 mol.% Ce2O3 was prepared by flame-

straying technique. The thermal properties of prepared glass were studied by DSC.

The DSC record revealed two exothermic effects observed at peak maxima

temperatures 936 and 1000 C, respectively, corresponding to crystallization of YAG

phase. On the basis of thermal properties studied by DSC, the glass was heat-

treated at several temperatures for 3h and the effect of the thermal treatment on

luminescence intensity was studied in detail. The emission photoluminescence

intensity was found to increase with the heat-treatment temperature of samples,

reaching the maximum for the sample treated at 1200 C.

Keywords: photoluminescence (PL), Ce3+

doped glass, yttrium aluminate glasses, LED, phosphor, YAG, Y3Al5O12

Introduction

Yttrium aluminium garnet (Y3Al5O12, YAG) has been recognized as one of the best

phosphor host materials for rare earth ions because of its optical transparency in the

range of ultraviolet to infrared [1]. YAG is a host with excellent structural

compatibility. The Y3+

and Al3+

ions can be substituted by many kinds of cations with

different sizes and valence in a certain extent [2]. The oxide garnet Y3Al5O12 (YAG),

when substituted with a few percent of the activator ion Ce3+

, is a luminescent

material that has nearly ideal photoluminescence properties for excitation by a blue

solid-state light source. YAG:Ce continues to be a phosphor with widespread use in

solid-state lighting due to high quantum efficiency, excellent chemical and thermal

stability, high mechanical strength and excellent optical properties [3]. It is still

attractive to develop simple and reliable synthetic methods for phosphor particles

with controlled morphology and structure. The one of the possible method is to

prepare polycrystalline material from glass; crystalline particles can be obtained by

controlled crystallization of glasses. In this case we can combine the

photoluminescence properties of luminophore hosted in polycrystalline and glass

matrix (e.g. YAG and residual Al2O3 enriched glass). In controlled crystallization,

nucleation and crystal growth kinetics are controlled by the diffusion of related

THE EFFECT OF THERMAL TREATMENT ON LUMINISCENCE PROPERTIES L4


K. Haladejová, R. Klement, L. Dvorská, A. Prnová, D. Galusek

XIth


24

components from glass matrix to the crystal surface. The size of the crystals can

readily be adjusted from several nanometers to hundreds of micrometers by

controlling the heat-treatment temperature and time [4]. The crystal morphology and

microstructure is strongly affected by the glass composition and thermal annealing

procedure [5,6]. In the present work we report the results of a preliminary study on

the preparation and characterization of Ce3+

doped yttrium aluminate glass,

containing 76.67 mol.% (60 wt.%) Al2O3 and 23.08 mol.% (40 wt.%) Y2O3,

denoted as Y40A60Ce0.5. The doping level was 0.25 mol.% Ce2O3 corresponding

to 0.5 at. % of Ce3+

. The effect of thermal treatment on luminescence intensity has

been studied in detail.

Experimental

The glass Y40A60Ce0.5 was prepared by flame-spraying technique from precursor

powders synthesized by the Pechini method in the form of transparent glass

microspheres with diameters from a few to several tens of micrometers. The glass

microspheres were characterized by scanning electron microscopy (SEM, JEOL

JSM-7600 F/EDS/WDS/EBSD), powder X-ray diffraction (XRD, Panalytical

Empyrean, CuKα radiation, at ambient temperature in the 2θ range of 10–80°). The

DSC traces were recorded using Netzsch STA 449 F1 Jupiter (TG/DTA/DSC) at the

heating rate of 10 °C/min in the temperature range from 25 °C to 1300 °C. The

emission and excitation fluorescence spectra were measured by Fluorolog 3 (FL3-

21, Horiba) fluorescence spectrometer in front-face mode. The Xe-lamp (450W) was

used as excitation sources for steady-state measurements. The time-temperature

regime for heat treatment of the sample Y40A60Ce0.5 was selected on the basis of

nucleation and crystallization experiments performed on STA instrumentation.

10 20 30 40 50 60 70 80

B

YAG (JCPDS No. 88-2047)o

+

+ o+

oo

oo

o

o

o o

ooo o

o

o

o

++

+

++

Inte

ns

ity

(a

.u.)

2

+

o

Al2O

3 (JCPDS No. 82-1399)

A

Fig. 1: SEM image of the as prepared microspheres in the system Y40A60Ce0.5

Fig. 2: XRD pattern of Y40A60Ce0.5 glass microspheres (A) and sample

crystallized at 1500 C for 5 h (B).

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Results and Discussion

The obtained glass microspheres with the composition Y40A60Ce0.5 were first

examined by optical microscopy and SEM (Fig. 1) that revealed spherical particles

with the diameter ranging from a few to several tens of microns. Closer inspection by

SEM shows regular features (faces) at the surface of some microspheres indicating

that some of them were at least partially crystalline. The composition of prepared

glass Y40A60Ce0.5 determined by SEM-EDS was found to be close to the

theoretical composition (41 wt.% Y2O3, 59 wt.% Al2O3); the content of Ce3+

was not

possible to determine reliably due to its low concentration. The SEM-EDS mapping

revealed that the distribution of Ce3+

ions in the sample is homogeneous. The

sample of glass microspheres was found to be XRD amorphous (Fig. 2A).

The thermal properties of the sample Y40A60Ce0.5 were examined by DSC

analysis. The DSC record of the glasses microbeads is shown in Fig. 3; two

crystallization peaks were observed. The glass transition temperature (Tg), onset of

crystallization temperature (Tx1, Tx2), and the temperature at maxima of the

exothermic crystallization peaks (Tp1, Tp2) were estimated from the DSC curve

(Fig. 3). The Tg of studied glass (endothermic effect) was found to be in the range

875 – 901 C (875 C onset, 893 C inflection and 901 C end of the endothermic

effect). The characteristics of exothermic effects corresponding to crystallization of

the sample Y40A60Ce0.5 are: Tx1 = 923 C, Tp1 = 936 C, Tx2 = 989 C and

Tp2 = 1000 C, respectively.

The time-temperature regime for heat treatment of the sample Y40A60Ce0.5 was

selected on the basis of the nucleation and crystallization experiments performed on

STA instrumentation. The selected temperatures are depicted as points on Fig. 3, the

dwell time for heat treatment at each temperature was 3h. The phase composition of

the sample after heat treatment was examined by XRD. The powder diffraction

traces revealed that the major phase is YAG up to 1200 C. The second phase, -

Al2O3, was only reliably observed in the XRD of sample treated at 1500 C for 5 h

(Fig. 2B).

The luminescence spectra of glass microspheres and of the heat-treated samples

are shown in Fig. 4. The excitation spectra (not shown) exhibit two broad bands

(absorptions) centred at 343 and 456 nm, respectively, that belong to the Ce3+

4f-5d

configuration. The emission spectra consist of a peak centred at 535 nm and a

shoulder at longer wavelength side 575 nm, which can be attributed to electronic

transitions of the Ce3+

. The former is assigned to the 5d1

2F5/2 transition and the

latter to the 5d1

2F7/2 transition, respectively. The emission (PL) intensity increases

with the heat-treatment temperature of the samples reaching the maximum for the

sample treated at 1200 C. This effect is most likely due to the higher packing

densities of Ce3+

ions in the crystallites than in glass [5]; crystallinity of the samples

increases with increasing treatment temperature, as documented by XRD and SEM.

THE EFFECT OF THERMAL TREATMENT ON LUMINISCENCE PROPERTIES L4


K. Haladejová, R. Klement, L. Dvorská, A. Prnová, D. Galusek

XIth


26

800 850 900 950 1000 1050 1100

Tx2

T

x1

Tf1

Tg

Tp2

Tp1

Ex

o

Temperature (°C)

Tp1

= 937oC

Tp2

= 1000oC

450 500 550 600 650 700 750

Y40A60Ce0.5_1200°C/3h

Y40A60Ce0.5_1050°C/3h

Y40A60Ce0.5_1000°C/3h

Y40A60Ce0.5_955°C/3h

Y40A60Ce0.5_935°C/3h

Y40A60Ce0.5_not treated

Y40A60Ce0.5_880°C/5h

exc

= 420nm

Wavelength [nm]

PL

in

ten

sit

y [

a.u

.]

Y40A60Ce0.5_1500°C/5h

Fig. 3: DSC trace of as prepared microspheres in the system Y40A60Ce0.5

Fig. 4: PL emission spectra of crystallized glass microspheres in the system Y40A60Ce0.5. The spectra shown in order of increasing temperature.

However, the PL intensity of the sample treated at 1500 C for 5h was found to be

lower than for the sample treated at 1200 C, despite the fact that also another

phase, -Al2O3, that could accommodate Ce3+

, crystallizes from the glass. This

decrease of PL intensity arises most likely from the oxidation of Ce3+

ions under the

applied conditions (higher temperature, longer time). Reduction of the samples in

H2/N2 atmosphere at identical conditions is essential for reliable comparison of the

PL intensities. Moreover, the light scattering may also play the role.

Conclusion

The Ce3+

doped glass Y40A60Ce0.5 was successfully prepared by flame-straying

technique. The thermal properties were studied by DSC technique and characteristic

temperatures have been estimated for two exothermic effects observed on DSC

trace. On the basis of thermal properties, the glass was heat-treated at several

temperatures for 3h. The emission (PL) intensity has been found to increase with the

heat-treatment temperature of the samples reaching the maximum for the sample

treated at 1200 C. This effect is most likely due to the higher packing densities of

Ce3+

ions in the crystallites than in glass.

Acknowledgements

The financial support of this work by the projects SAS-NSC JRP 2012/14 and VEGA

1/0631/14, is gratefully acknowledged. This publication was created in the frame of the project

"Centre of excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056,

based on the Operational Program Research and Development funded from the European

Regional Development Fund.

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References

[1] Xianju Zhou, Kaining Zhou, Yingmao Li, Zhiliang Wang, Qiaochun Feng, J. Lumin. 132 (2012) 3004. [2] Zhongfei Mu, Yihua Hu, Haoyi Wu, Chujun Fu, Fengwen Kang, J. Alloy Compd. 509 (2011) 6476. [3] Yujun Liang, Dongyan Yu, Wen Zhu Huang, Mengfei Zhang, Guogang Li, Chunjie Yan, Materials Science in

Semiconductor Processing 30 (2015) 92–97. [4] G. He, L. Mei, L.L. Wang, G.H. Liu, J.T. Li, Cryst. Growth Des. 11 (2011) 5335. [5] L. Wang, L. Mei, G. He, G. Liu, J. Li, L. Xu, J. Lumin. 136 (2013) 378. [6] D.M. Frederick, A.E. Gash, R.L. Landingham, J.H. Satcher, Z.A. Munir, J. Non-Cryst. Solids 363 (2013) 64.

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DENSITY OF STATES AND MAGNETIC PROPERTIES OF Al2O3,

Cr2O3 AND CrO2

Jozef Drga1, Jozef Chovanec

2, Stanislav Šimkovič

3

1Alexander Dubček University of Trenčin, Študentská 2, 911 50 Trenčín, Slovak Republic 2Vitrum Laugaricio – Joint Glass Centre of the IIC SAS, TnU AD, FChPT STU, a.s, Študentská 2, 911

50 Trenčín, Slovak Republic 3J. Branecký High School of Piarists, Palackého 4, 911 01 Trenčín, Slovak Republic

e-mail:[email protected]

This contribution deals with theoretical investigation of density of states of three ceramic

materials. We used Al2O3 as diamagnetic, Cr2O3 as antiferromagnetic and CrO2 as

ferromagnetic material. The density of states, Fermi levels and magnetic moment were

calculated using software VASP.

Key words: ceramic materials, magnetic properties, density of states, Fermi level

Introduction

Many magnetic materials are used for ceramics production. Three materials were

selected for simulation of density of states and magnetic properties. Density of

states (DOS) of a system describes the number of states per interval of energy that

are available to be occupied. For investigation a density of states the Density

functional theory (DFT) was used [1].

Al2O3, Cr2O3 and CrO2 were considered as important ceramic materials. The Al2O3 is

a diamagnetic material. Diamagnetic effect is caused by a reaction of electron and

Lorentz force. The superconductors cooled below the critical temperature are perfect

diamagnets [2].

Al2O3 powders are often used in nanotechnology area, in preparation of special

ceramic materials etc. Its porous structure, high surface area, and high catalytic

surface activity belong among its benefits. This material is considered a suitable

alternative to silicon for producing semiconductor non-volatile random access

memories for future applications [3]. Investigation of the atomic and electronic

structures of γ-Al2O3 has also received much attention over recent years [4, 5]. This

material is widely used as a catalyst, an adsorbent, and a support for industrial

catalysts [6]. Structural variations of α−Al2O3 and γ-Al2O3 are the most prominent of

the transition aluminas [7]. In nanocrystalline form, γ-Al2O3 is more stable than

α-Al2O3 due to its lower surface energy [8]. Calculations of density of states were

published in various papers [9, 10].

Paramagnetism is caused by unpaired electron spin. The spins do not interact with

each other and in the absence of external field they are oriented randomly – contrary

to ferromagnetic materials. The ferromagnetic material shows the paramagnetic

behaviour above the Curie temperature can be explained by thermal vibrations of

lattice.

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J. Drga, J. Chovanec, S. Šimkovič


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Neighbouring spins are oriented in opposite directions in the antiferromagnetic

material. Several properties of Cr2O3 such as high hardness (29.5 GPa) and high

melting point (~2300 ºC) combined with chemical inertness, low friction coefficient,

high wear resistance, high solar absorption coefficient make it an important material

in a variety of industrial applications [11]. Several studies [12,13] have been devoted

to its utilization in spintronic applications. Cr2O3 is useful in the preparation of

fluorescent magnetic nanocomposite material [14]. The Cr2O3 is a magnetoelectric

material with antiferromagnetic insulator with the Néel temperature of 308 K [15].

The Cr2O3 behaves as a paramagnetic material over Néel temperature. It is an

insulator with a bandgap of 3.31 eV in anti-ferromagnetic phase [16].

Chromium dioxide (CrO2) is at present one of the most important ferromagnetic

ceramic materials. It is used as magnetic media in audio and video recording tapes.

In a ferromagnetic material an interaction between adjacent dipoles takes place. The

interactions form magnetic domains in material. After the biasing magnetic field is

removed, the material remains magnetic.

CrO2 provides practical application in technology of spintronics, tunneling devices,

magnetic reading heads and magnetic field sensors, because of its optical,

electronic, magnetic and structural properties [17, 18, 19]. 390 K is the Curie

temperature of CrO2. One publication states the magnetic moment m = 1.772 µB

[20]. Recent publications present this quantity as m =2.00 µB [18, 21, 22].

Computational method

In the theoretical investigation, the Vienna ab initio simulation package (VASP) was

used [23, 24] using projector augmented-wave PAW pseudopotentials (PPs) [25, 26]

and the Perdew–Burke–Ernzerhof (PBE) exchange-correlation functional [27, 28].

The computations were based on DFT theory and we used the Conjugate gradient

algorithm [29]. An attempt to solve time – free Schrodinger equation was made:

„Expected“value of Hamiltonian was proposed and diagonalized.

The residual vector was minimized

,

where

Two steps were considered. In the first stage the structure of materials was

calculated with the ionic and electronic relaxation. The number of K-Points was 6.

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The second step was DOS calculation. All calculations were performed at zero

temperature 0 K. In this case the number of K-Points increased to 18.


The α-Al2O3 unit cell has a hexagonal cell structure, with a rhombehedral primitive

cell. In the first calculation of the ionic and electronic relaxation of Al2O3 12 atoms of

aluminium and 18 atoms of oxygen were used. The basic structural parameters for

computing were obtained from the authors Toebbens et. al. [30]. The basic primitive

cell of Al2O3 is shown in the Figure 1. Cr2O3 consists of 12 atoms of chrome and 18

atoms of oxygen. The structure is similar to Al2O3 (rhombehedral) and the lattice

parameters were published by Baster et. al. [31]. The primitive cell of CrO2 is smaller

than the Al2O3 and Cr2O3. It consists of 2 atoms of chrome and 4 atoms of oxygen.

The structure of CrO2 is tetragonal and the basic information about the structure was

published by Cloud [32].

Fig. 1: Snapshot of Al2O3 primitive cell

The Fermi level and Magnetic moment of Al2O3 was calculated. The value of Fermi

level is Ef = 4.922 eV and magnetic moment per unit cell was calculated as

m = -0.0001 µB: the parameters for Cr2O3: Ef = 7.579 eV and m = 36 µB were

computed. CrO2 was used as the last material. Calculated values are: Ef = 6.055 eV

and m = 4 µB. In Figure 2 the density of states of Al2O3 can be seen. The DOS for

bulk Alumina (Fig. 2) shows that it is plainly a very strong insulator, with a very wide

band gap. The symmetry of DOS shows that the material has compensated spins.

The Cr2O3 DOS structure indicated antiferromagnetism in absolute zero

temperature, because the maxima of values of adverse up and down curves have

different values. The paramagnetic behaviour occurs over the Néel temperature

308 K [33]. The CrO2 behaves as ferromagnetic due to strong asymmetry of DOS

near Fermi level (Fig. 3).

The computer cluster which consist 32 processors was used for DFT simulations of

DOS. The simulations are very exacting of computer time. The simulation of one

material takes several days. If the multiplied cell had been used, a little more

accurate results may have been acquired. However, the obtained results were

sufficient; therefore there was no need for multiplying the cell. [18, 21, 22].

L5 DENSITY OF STATES AND MAGNETIC PROPERTIES OF Al2O3, Cr2O3 AND CrO2

J. Drga, J. Chovanec, S. Šimkovič


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Fig. 2: DOS of Al2O3

Fig.3: DOS of Cr2O3

Fig. 4: DOS of CrO2

Conclusion

Density of states of Al2O3, Cr2O3 and CrO2 materials was analysed. These materials

are used for preparation of ceramics. The basic properties of these materials; Fermi

energy levels and Magnetic moment per unit cell were calculated. Following values

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were obtained by optimising the structure and subsequently calculating DOS for

Al2O3: Ef = 4.922 eV and m = -0.0001 µB; Cr2O3: Ef = 7.579 eV and m = 36 µB. CrO2:

Ef = 6.055 eV and m = 4 µB. These values indicate the correct initial conditions.

The DOS curve of Al2O3 is symmetric, contrary to other oxides which were analysed.

The symmetry of curve shows non ferromagnetic behaviour. The small (negative)

value of magnetic moment indicates diamagnetic state in Al2O3 material. On the

Cr2O3 DOS curve shows the antiferromagnetistic behaviour. The asymmetric shape

of CrO2 DOS function signalise the typical ferromagnetic material.

Referecnes

[1] D.S. Sholl, J.A. Steckel, (2009) Density functional theory. A Practical Introduction, ISBN 9780470373170,

[2] C. Carter, M. Norton (2007) Ceramic materials – Science and Engineering, ISBN 0387462708 [3] G. Molas, M. et al. Reliability of charge trapping memories with high-k control dielectrics (invited paper).

Microelectronic Engineering, 86:1796–1803 (2009). [4] G. Gutiérrez, A. Taga, B. Johansson: Physical Review B, 65:012101 (2001). [5] H.P. Pinto, R.M. Nieminen, S.D. Elliott: Physical Review, B 70:125402 (2004). [6] I. Halasz, A. Brenner, M. Shelef, Applied Catalysis B: Environmental, 2:131-146 (1993). [7] S. Wang, A.Y. Borisevich, S.N. Rashkeev, M.V. Glazoff, K. Sohlberg, S.J. Pennycook, S. Pantelides, Nat.

Mater. 3(3), 143 (2004). [8] J.M. McHale, A. Auroux, A.J. Perrota, A. Navrotsky, Science 277, 5327 (1997) [9] Yazdanmehr et al., Nanoscale Research Letters 7:488 (2012). [10] C. Loyola · E. Men´endez-Proupin, G. Guti´errez J. Mater. Sci. 45, 5094 (2010). [11] P.M. Sousa, A.J. Silvestre, O. Conde, Thin Solid Films, 519, 3653–3657 (2011). [12] P. Borisov, A. Hochstrat, V.V. Shvartsman, W. Kleemann, P.M. Hauck, Integrated Ferroelectrics: An Int. J.,

99, 69–76 (2008). [13] X. He, W. Echtenkamp, Ch. Binek, Ferroelectrics, 426, 81-89 (2012). [14] Ch. Borgohain, K.K. Senapati, D. Mishra, K.Ch. Sarma, P. Phukan, Nanoscale, 2, 2250-2256 (2010). [15] N. Iwata, T. Asada, S. Ootsuki and H. Yamamoto: Growth, Ferroelectrics, Multiferroics, and Magnetoelectrics,

1034 (2008). [16] Y. Guo, S.J. Clark, J. Robertson, J. Phys. Condens. Matter 24, 325504 [17] E. Kulatov, I. Mazin, J. Phys: condnens Matter, 2, 343 (1990). [18] P. Lewis, B. Allen, T.Sasaki, Phys rev B., 55 10253( 1997). [19] B Maddox, et al., Phys rev B13, 144111 (2006). [20] A V. Nikolaev, B V Andreev, Phys Solid State 35, 603 (1993). [21] V Srivastava et al., Indian Journal of Pure and Applied physics, 46, 397-399 (2008). [22] K.H Schwarz, J. Phys F. Met Phys. 16, L211 (1986) [23] G. Kresse, J. Furthmuller, Comput. Mater. Sci. 6, 15 (1996). [24] G. Kresse, J. Furthmuller, Phys. Rev. B 54, 11169 (1996). [25] P. Blöchl, Phys. Rev. B 50, 17953, 17 (1994). [26] G. Kresse and D. Joubert, Phys. Rev. B 59, 1758 (1999). [27] J. Perdew, K. Burke, M. Ernzerhof, Phys. Rev. Lett. 77, 3865, 19 (1996). [28] J. Perdew, K. Burke, M. Ernzerhof, Phys. Rev. Lett. 78, 1396 (1997). [29 D.M. Bylander et al. Phys. Rev.B 42, 1394 (1990), [30] D.M. Toebbens, N. Stuesser, K. Knorr, H.M. Mayer, G. Lampert, Materials Science Forum 378, 288-293 (2001). [31] M. Baster, F. Bouree, A. Kowalska, Z. Latacz, Journal of Alloys Compd. 296, 1-5 (2000). [32] W.H. Cloud, D.S. Schreiber, K. R. Babcock, Journal of Applied Physics; 33,1193-1194 (1962). [33] Mu Sai, A. L. Wysocki, K. D. Belashchenko, Phys. Rev. B; 87 (2013) .

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CERAMIC MATRIX REINFORCED BY BN NANOFILLERS

Monika Kašiarová1, Peter Tatarko

2, Zdeněk Chlup

3, Ivo Dlouhý

3

1Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, Košice, Slovak Republic 2Nanoforce Technology Limited, Queen Mary University of London, E1 4NS, London, United Kingdom

3Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Žižkova 22, Brno,

Czech Republic


Hollow “cylindrical” and “bamboo-like” boron nitride nanotubes (1 wt.%) have been used to

reinforce 3Y-TZP zirconia ceramics via spark plasma sintering. No significant influence of

different morphologies of BNNTs on the mechanical properties at the macro-scale (elastic

modulus, hardness, and fracture toughness) has been observed. The addition of the BN

nanotubes into 3Y-TZP zirconia resulted in the improvement of the fracture toughness of

about 83 - 97 % compared to the reference ZrO2.

Keywords: 3Y-TZP zirconia, boron nitride nanotubes, spark plasma sintering

Introduction

The outstanding mechanical properties of tetragonal zirconia stabilized with 3 mol.%

yttria (3Y-TZP) ceramics are related to a stress shielding effect by the transformation

toughening that provides their excellent mechanical properties [1]. However, the

fracture toughness of the submicrometric or nanocrystalline zirconia is reduced

because the transformation toughening mechanism is lost [2]. Therefore, the

fracture toughness of the submicrometric or nanocrystalline zirconia can be further

increased by the incorporation of secondary phases into the zirconia matrix, such as

nanotubes, nanofibres or nanoplatelets, which eliminate the fracture toughness

reduction due to the small zirconia grain size. Well-known carbon nanotubes are not

the only type of nanotubes that have been investigated as a reinforcement, but there

are other types of ceramic-based nanotubes with even better properties. Boron

nitride nanotubes (BNNTs) can be a new type of reinforcement, being found to be as

strong, flexible and elastic as CNTs due to the same crystal structure [3]. Similar to

CNTs, they exhibit high tensile strength of approximately 30 GPa [4] and an

experimentally determined Young`s modulus of 1.22 TPa [5]. Their main advantage

over the CNTs is that they are much more thermally and chemically stable than

CNTs [4], which makes them more suitable for strengthening and toughening of

ceramics [3].

The purpose of the present work is to investigate and clarify the influence of two

different inner morphologies of boron nitride nanotubes, hollow “cylindrical” and

“bamboo-like”, on the mechanical properties of zirconia ceramics. The main interests

are focused on the microstructure evolution and fracture toughness of

CERAMIC MATRIX REINFORCED BY BN NANOFILLERS L6 M. Kašiarová, P. Tatarko, Z. Chlup, I. Dlouhý

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nanostructured tetragonal zirconia stabilized with 3 mol.% yttrium prepared by spark

plasma sintering.

Experimental procedure

The starting powder was a high purity 3 mol.% yttrium-stabilized tetragonal zirconia

(3Y-TZ, Tosoh Co. Ltd.) with an average particle size of 60 nm and two different

kinds of boron nitride nanotubes: hollow “cylindrical” (BNNT-C) and “bamboo-like”

(BNNT-B). In order to obtain a homogenous powder mixture, the ZrO2 and 1 wt.%

BNNTs powders were ball milled in ethanol for 12 h using zirconia balls. After that,

the slurry was dried in a furnace and then powder was screened using a 100-mesh

sieve, followed by drying at 300 °C for 3 h to remove the ethanol. Then the powder

mixture was sintered in graphite die with diameter of 30 mm using SPS at 1250 °C

under pressure of 50 MPa for 5 min in vacuum. The bulk density of the sintered

samples was measured by the Archimedes method in distilled water. The elastic

modulus was determined on polished test bars 2 x 3 x 20 mm3 by the resonance

method using GrindoSonic Mk5i. The hardness was determined by the

measurement of length of indent diagonals created by the Vickers indentations at

load of 49.05 N for 10 s. The chevron notched beam technique (CNB) was used for

the fracture toughness determination based on ASTM C1421 standard.

Results

The grain size distributions of monolithic zirconia and ZrO2 composite reinforced by

the bamboo-like BN nanotubes are shown in Fig.1. Reference zirconia contains large

proportion of very fine zirconia grains, around 100 150 nm with the maximum grain

size up to 350 nm. After the addition of BNNTs, the distribution was shifted to a

larger grain sizes and becomes wider. Peak of the distribution is still around 100

150 nm, but large grains (around 600 nm) also appeared. Unusual growth of grains

is connected with the presence of metal catalysts inside the BNNTs which form

during sintering metal oxide and changed the sintering behaviour from solid-phase

to liquid-phase sintering.

Fig. 1: Grain size distribution in the pure ZrO2 and ZrO2-BNNT-B composite

0,0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,80

5

10

15

20

25

30

Re

lati

ve

fre

qu

en

cy

[%

]

ZrO2 grain diameter [m]

d = 157 ± 58 nm

ZrO2

0,0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,80

5

10

15

20

25

30

d = 226 ± 105 nm

ZrO2 grain diameter [m]

ZrO2-BNNT-B

Re

lati

ve

fre

qu

en

cy

[%

]

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According to Table 1, all the investigated ceramic materials are fully dense and the

incorporation of the BNNTs did not lead to a decrease in the density of the material.

Table 1 shows that the elastic modulus of the ZrO2-BNNTs composites slightly

decreased with the increasing amount of both kinds of the nanotubes. On the other

hand, the hardness values increased after the addition of 1 wt.% of both kinds of the

nanotubes. The increase in the hardness for these materials is due to higher density

of the materials as the densification has been promoted by the addition of the

BNNTs.

Table 1: Density, grain size and some mechanical properties of ZrO2-BNNTs composites

Material Relative

density [%] Average grain

size [nm] Elastic

modulus [GPa]

Hardness

[GPa]

KIC

[MPa.m1/2]

ZrO2 pure 98.7 157 ± 58 210 ± 1.5 12.5 ± 0.2 3.95 ± 0.09

1.0 wt.% BNNT-C 99.5 231 ± 99 202 ± 2.0 13.5 ± 0.3 7.81 ± 0.79

1.0 wt.% BNNT-B 99.7 226 ± 105 209 ± 1.5 13.9 ± 0.2 7.26 ± 0.17

The fracture toughness significantly increased after the addition of BNNTs reaching

a very similar improvement in the range of 83 97 %. Similarly to the hardness and

the elastic modulus, no significant influence of different structures of BN nanotubes

on the fracture toughness has been observed. This result indicates that the different

morphologies of BNNTs do not have any significant effect on the mechanical

properties at macro-scale.

Fig. 2: Fracture surfaces after CNB fracture toughness measurements: a) pure ZrO2; b) ZrO

2 + 1 %

BNNT-B

Fig. 2 shows the fracture surfaces of the monolithic zirconia and the ZrO2 composite

reinforced by the bamboo-like BN nanotubes after the CNB fracture toughness tests.

It is obvious that in the case of monolithic zirconia the fracture is mostly intergranular

(Fig. 2a), on the other hand, after the addition of the BN nanotubes the fracture

morphology is changed from intergranular to transgranular (Fig. 2b). Grains without

a) ZrO2 b) ZrO

2+1% BNNT

CERAMIC MATRIX REINFORCED BY BN NANOFILLERS L6 M. Kašiarová, P. Tatarko, Z. Chlup, I. Dlouhý

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any nanotubes around are broken in the intergranular mode, but when the grains are

surrounded by the nanotubes, the fracture surface is obviously transgranular. All of

these observations clearly confirm that the BN nanotubes strengthen the grain

boundaries of zirconia ceramics. Moreover, the presence of nanotubes causes

significant toughening of composites.

Conclusions

Two different inner structures of boron nitride nanotubes, hollow “cylindrical” and

“bamboo-like”, have been used to reinforce the nanostructured tetragonal zirconia

stabilized with 3 mol.% yttrium by SPS. It has been observed that the different

structure of the BNNTs do not have any significant effect on the mechanical

properties at the macro-scale. The elastic modulus slightly decreased with the

increasing amount of the BNNTs. Hardness values increased after the addition of

1 wt.% BNNTs. The addition of 1 wt.% BNNTs into the zirconia matrix produced

about 83 97 % increase in the fracture toughness compared to the monolithic

zirconia.

Acknowledgement

The work was supported by Slovak Research and Development Agency under the contract

No. SK-CZ-2013-0194 and APVV-0161-11 and project of Slovak Grant Agency No.

2/0189/15.

References

[1] Hannink RHJ, Kelly PM, Muddle BC. Transformation toughening in zirconia-containing ceramics. J Am Ceram Soc. 83 (2000) 461-87.

[2] Garmendia N, Grandjean S, Chevalier J, Diaz LA, Torrecillas R, Obieta I. Zirconia-multiwall carbon nanotubes dense nano-composites with an unusual balance between crack and ageing resistance. J Eur Ceram Soc. 31 (2011)1009-1014.

[3] Xu JJ, Bai YJ, Wang WL, Wang SR, Han FD, Qi YX, Bi JQ. Toughening and reinforcing zirconia ceramics by introducing boron nitride nanotubes. Mat Sci Eng. A 546 (2012) 301-6.

[4] Golberg D, Bai XD, Mitome M, Tang CC, Zhi CY, Bando Y. Structural peculiarities of in situ deformation of a multi-walled BN nanotube inside a high-resolution analytical transmission electron microscope. Acta Mater. 55 (2007)1293-98.

[5] Chopra NG, Zettl A. Measurement of the elastic modulus of a multi-wall boron nitride nanotube. Solid State Commun 105 (1998) 297-300.

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STUDY OF THERMAL BEHAVIOUR

AND HOT PRESSING OF ALUMINATE GLASS MICROSPHERES

Anna Prnová1, Robert Klement

1, Katarína Bodišová

2, Lubomír Hric

3,

Erich Neubauer3, Dušan Galusek

1, Els Bruneel

4, Isabel Van Driessche

4

1Vitrum Laugaricio – Joint Glass Centre of the IIC SAS, TnU AD, and FChPT STU, Trenčín, Slovakia, 2Faculty of Chemical and Food Technology, Slovak University of Technology; Bratislava, Slovakia,

3 RHP-Technology GmbH, Seibersdorf, Austria, 4Department of Inorganic and Physical Chemistry, Ghent University, Ghent, Belgium


Yttrium-aluminate glass microspheres with the eutectic composition in the pseudo-binary

system Y3Al5O12-Al2O3 (i.e. 76.8 mol.% Al2O3, 23.2 mol. % Y2O3) were prepared by

combination of flame spraying and sol-gel Pechini method. Optical microscopy, SEM, XRD,

DSC, and HT XRD analysis were used for characterisation of prepared system. The glass

microspheres were hot-pressed at the pressure of 80 MPa, and temperatures 900, 1050,

1150, 1300, 1600 °C with holding time 0 - 30 min prepare bulk materials. Sintered samples

were characterised by XRD, SEM, SEM EDX analysis and Vickers hardness and densities of

sintered bodies were measured. A bulk polycrystalline material with fine grained

microstructure and the hardness HV = 17.7 GPa was prepared by hot pressing at 1300 °C for

30 min.

Keywords: glass microspheres, flame synthesis, hot-pressing

Introduction

Yttrium-aluminate glasses with high alumina content are potential candidates for

various applications (transparent ballistic protections, infrared transparent windows,

materials for solid state lasers), because of their outstanding mechanical and optical

properties. Also high structural stability (comparable to those of sapphire single

crystal or polycrystalline alumina) under extreme mechanical, electrical, thermal and

chemical influence is expected [1, 2]. The combination of sol-gel method and flame

synthesis and subsequent hot-press sintering of prepared glass microbeads appears

to be a suitable method for preparation of glass and glass ceramic materials with

high degree of homogeneity. In this work, yttrium-aluminate glass microspheres with

eutectic composition A60Y40M (76.8 mol.% Al2O3, 60 wt.% Al2O3), were prepared

by combination of sol-gel – Pechini synthesis and the flame-spraying method.

Optical microscopy, SEM, XRD, DSC, HT XRD were used for characterisation of

prepared glass particles. The hot-press (HP) experiments were performed on the

basis of DSC, HT XRD analysis and the results of our previous work [3]. Vickers

hardness and density measurements of sintered bodies were performed and SEM

analysis was used to study the microstructure of bulk materials prepared by hot

pressing.

STUDY OF THERMAL BEHAVIOUR AND HOT PRESSING OF ALUMINATE GLASS MICROSPHERES L7 A. Prnová, R. Klement, K. Bodišová, L. Hric, E. Neubauer, D. Galusek, E. Bruneel, I. Van Driessche

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Experimental part

The Al2O3-Y2O3 precursor powder was prepared by modified Pechini sol-gel method

[4]. The aluminium nitrate solution in deionized water and yttrium nitrate solution

prepared by dissolution of yttrium oxide in diluted nitric acid were mixed together.

Then, an aqueous solution of citric acid and ethylene glycol (in the molar ratio 1 : 1)

was added to the solution of nitrates. The resulting solution was refluxed at 85 C for

2 hours and then heated to 150 C to promote polymerization (solution viscosity

rapidly increased) and solvent evaporation until an aerated resin was formed.

Finally, the organic compounds were removed by heating to 800 °C for six hours.

Prepared powders were fed into methan-oxygen flame and molten particles were

cooling by distilled water, collected, separated, dried and calcined at 650 °C for 4 h

to remove of residua from flame synthesis. The morphology of prepared glass

particles was studied by optical microscopy (Nicon ECLIPSE ME 600) and SEM

(Zeiss EVO 40HV at accelerating voltage 20 kV). X-ray powder diffraction analysis

(Panalytical Empyrean, CuKα radiation, 2Θ range 10-80) was used to obtain

qualitative information on the phase composition of prepared glass particles. The

thermal behaviour of prepared microspheres was studied by DSC (Netzsch STA 449

F1 Jupiter) and HT XRD (Panalytical Empyrean, equipped with high temperature cell

Anton Paar HTK 16). The HP experiments at temperatures 900 °C (sintering by

viscous flow), 1050, 1150, 1300, 1600 °C (sintering by diffusion flow), different

holding time (0 - 30 min) and pressure 80 MPa were performed. The sintered

samples were polished and etched in boiling H3PO4 and microstructure was studied

by SEM. Finally hardness (Vickers microhardness equipment - AFFRI WIKI 200),

and density (liquid pycnometry) of prepared samples was measured.


Prepared glass particles were spherical, transparent, fully re-melted and XRD

amorphous. The study of morphology by SEM revealed the presence of fully

amorphous (glassy) microspheres. For detailed study of thermal behaviour of

prepared system HT XRD analysis in the temperature interval 750 – 1250 °C with

the step of 5 °C was performed. Formation of YAG as the only crystalline phase was

observed in a wide temperature range from 750 °C to 1200 °C, with the maximum

increase of the YAG phase content in intervals 930 – 935 °C and 995 – 1000 °C.

This is in good agreement with the results of DSC, where two exothermic peaks with

the maxima at 935 and 999 °C (crystallization of YAG in two steps) were observed.

The conditions of HP experiments (selected on the basis of DSC, HT RTG and the

results of our previous work) and the values of Vickers hardness, relative density

and XRD quality of sintered samples are summarized in Tab.1. The XRD

examination of hot pressed ceramic bodies revealed YAG as a major phase in all

samples. The materials sintered at lower temperatures were amorphous (at 900 °C)

or partially crystalline, containing YAG phase (at 1050 and 1150 °C). The samples

sintered at 1300 and 1600 °C were polycrystalline and contained YAG and α-Al2O3.

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The sample A6Y1 hot pressed at 900 °C was soft and porous, and Vickers hardness

(HV) was not possible to measure. The HV of 17.7 GPa and very fine-grained

microstructure was achieved in the sample sintered at 1300 °C (pressure 80 MPa)

for 30 min. This higher value of hardness in comparison to our previously published

results may be attributed to better homogeneity of glass microspheres, as confirmed

by SEM. The sample sintered at 1600 °C for 30 min shows lower HV value

(12.7 GPa). The reduction is probably caused by excessive grain growth in the

sample during sintering.

Table 1: The conditions of selected HP experiments and HV values of sintered samples

Sample name

Temperature

[°C]

Press

[MPa]

Dwell time

[min]

Hardnes

[GPa]

XRD

(phases)

ςrel.(%)

A6Y1 900 80 0 n.m amorphous 84.5

A6Y2 1050 80 0 13.1 YAG 96.8

A6Y3 1150 80 0 12.2 YAG 98.6

A6Y4 1300 80 30 17.7 YAG α-Al2O3 104

A6Y5 1600 80 30 12.7 YAG α-Al2O3 98.8

Conclusion

The yttrium-aluminate glass microspheres with eutectic composition were prepared

and HP sintered at different temperatures. The HV 17.7 GPa was achieved for

sample sintered at pressure 80 MPa, temperature 1300 °C and holding time 30 min.

Also, SEM analysis of H3PO4 etched surface shown very fine microstructure of this

prepared sample.

Acknowledgements

The financial support of this work by the projects SAS-NSC JRP 2012/14, and VEGA 1/0631/14, is gratefully acknowledged. This publication was created in the frame of the international academic agreement (IIC SAS, Slovakia, University of Gent, Belgium), with a financial contribution of FWO and KVAB (Belgium) institutions and the project "Centre of excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056, based on the Operational Program Research and Development funded from the European Regional Development Fund.

References

[1] Y. Mizutani, H. Yasuda, I. Ohnaka, N. Maeda, Y. Waku, Coupled growth of unidirectionally solidified Al2O3–YAG eutectic ceramics, J. Cryst. Growth, 244, 384-392 (2002).

[2] T.A. Parthasarathy, T. Mah, Deformation behavior of an Al2O3-Y3Al5O12 eutectic composite in comparison with sapphire and YAG, J. Am. Ceram. Soc,. 76, 29 (1993).

[3] A. Prnová, D. Galusek, M. Hnatko, J. Kozánková, I. Vávra, Composites with eutectic microstructure by hot pressing of Al2O3-Y2O3 glass microspheres, Ceramics-Silikáty, 55, 3, 208-213(2011).

[4] P. Pechini, “Method of Preparing Lead and Alkaline-Earth Titanates and Niobates and Coating Method Using the Same to Form a Capacitor,” U.S. Pat. No. 3 330 697, July 11, (1967)

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OXIDATION RESISTANCE OF SiC CERAMICS PREPARED

BY DIFFERENT PROCEESSING ROUTES

Alexandra Kovalčíková1, Zoltán Lenčéš

2, Jaroslav Sedláček

2, Pavol Šajgalík

2,

Ján Dusza1, Martin Ignácz

3, Mária Mihaliková

3

1Institute of Materials Research Slovak Academy of Sciences, Watsonova 47, Košice, Slovakia 2Institute of Inorganic Chemistry Slovak Academy of Sciences, Dúbravská cesta 9, 845 36 Bratislava,

Slovakia 3Faculty of Metallurgy, Technical University Košice, Letná 9, 04001 Košice, Slovakia


The oxidation behaviour of SiC ceramics was investigated as a function of the processing

route at oxidizing condition at 1400 °C/0 204 h. The oxidation resistance increases with

increasing temperature of heat treatment from 1650 °C to 1850 °C. Rapid hot pressed SiC

materials without sintering additives show better oxidation resistance compared to HP LPS

SiC. Oxidation always followed parabolic rate law indicates diffusion as the rate limiting

mechanisms.

Keywords: silicon carbide, sintering, oxidation resistance

Introduction

Silicon carbide has been recognised as important structural ceramic because of its

good combination of mechanical and thermal properties. SiC ceramics show good

wear, oxidation and creep resistance at high temperatures, but relatively low fracture

toughness [1]. Oxidation behaviour of SiC ceramics depends on many factors as

partial pressure of oxygen, temperature, grain boundary phase composition or

amount of sintering additives. There are two effective methods for improving the

high-temperature properties of silicon carbide materials to increase the

refractoriness of the grain boundary phase and to promote crystallization of the

amorphous phase. This can be achieved by the selection appropriate additives and

chemical compositions. A problem associated with the use of additives is the

degradation of high temperature properties due to residual grain boundary phase.

Sintering additives such as Al2O3Y2O3 and AlNY2O3 were extensively used to

obtain dense LPSSiC ceramics [2]. These oxides react with SiO2 which is always

present at the surface of SiC and AlN particles and form an oxynitride melt.

SiCAlNRE2O3 systems showed favourable oxidation resistance due to the higher

eutectic temperature of the grain boundary phases, RE2Si2O7 (RE = Er, Lu, Sc)

compared to oxynitride phases. The aim of this study is to investigate the influence

of different sintering on oxidation resistance of SiC ceramics.


L8 OXIDATION RESISTANCE OF SiC CERAMICS PREPARED BY DIFFERENT PROCEESSING ROUTES

A. Kovalčíková, Z. Lenčéš, J. Sedláček, P. Šajgalík, J. Dusza, M. Ignácz, M. Mihaliková


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Experimental material and methods

Two types of materials were prepared. In first case, -SiC powder was mixed with

3 wt.% Al2O3 and 6 wt.% Y2O3.The powder mixtures were ball milled in isopropanol

with SiC balls for 24 h. The suspension was dried and sieved through 25 m sieve

screen in order to avoid hard agglomerates. The samples were hot pressed at

1850 C/1h under mechanical pressure of 30 MPa in N2 atmosphere. The hot

pressed samples were subsequently annealed at 1650 °C/5 h or 1850 °C/5 h (Table

1).

In second case, Crystalline SiC powder was mixed with 1.5 wt.% of defloculant,

2 wt.% of plastificator and 1.6 wt.% of binder by ball milling for 24 h with SiC milling

balls in methanol. The suspension was subsequently freeze granulated in liquid

Nitrogen to granules with the diameter of 300 600 m and lyophilized at -60 °C

under vacuum. The granulated powder was poured into a graphite die and densified

by rapid hot-press at 1850 °C/20 or 60 min under mechanical pressure 30 MPa in

vacuum.

Dried clean samples were weighed and the exact dimensions were measured in

order to calculate the surface area. The oxidation experiments were performed in a

furnace at 1400 °C in air, for time in the range 0 204 h. At intervals, the bars were

removed from the furnace, weighed to compute the specific mass change (m), and

returned for further oxidation. After oxidation tests the surface of oxidized samples

was observed by SEM with EDS analysis.


A plot of weight gain per unit surface area as a function of time at 1400 °C is on Fig.

1. The materials with fine globular microstructure (SiCHP, SiC1650) [3] had a

specific weight gains on the order of 3.24 -3.17 mg/cm2 after 204 h. Specimens with

plate-like microstructure (SiC - 1850) showed lower mass gain of 1.73 mg/cm2.

Square of the specific weight gain as a function of oxidation time is shown in Fig. 2.

According to Eq. (1), the slope of the straight lines corresponds to parabolic

oxidation rate constants:

QtkW p .2 (1)

where 2W is the weight gain per unit surface area, kp is the rate constant of

parabolic oxidation, t is the exposure time and Q is an additive constant (ideally = 0).

Table 1 contains the parabolic oxidation rate constants for all experimental samples,

as well as the total specific weight gains after 204 h. As can be seen there exist

appreciable differences between the oxidation kinetics of material hot pressed, heat

treated at 1650 1850 °C and rapid hot pressed. For HP ceramics the specific mass

gain and rate constant of parabolic oxidation decrease with increasing of annealing

temperature. On the other site, RHP samples had better oxidation resistance

compared to HP materials. The sample SiCRHP60 with the highest oxidation

resistance in this study had a specific rate constant of 4.6 x 10-5

mg2.cm

-4.h

-1 at

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1400 °C. This parabolic oxidation rate is about three orders of magnitudes lower

compared to HP samples, kp = 1.47 5.15 x 10-2

mg2.cm

-4.h

-1. Nevertheless, there is

always mass gain throughout the oxidation period indicating that oxidation is passive

because active oxidation is accompanied by mass loss.

Fig. 1: Specific mass gain versus oxidation time

Fig.2: Square of specific mass gain versus oxidation time

Table 1: Mechanical properties and parameters of oxidation of SiC ceramic

The cross-sections of the oxidized layers can be seen in Fig. 3. The oxide layer

thickness was determined to be ~32 m for materials SiCHP and SiC1650 (Fig.

3a) and around 5 m for materials SiC1850. A few of pores were observed at the

interface between the oxidized layer and SiC bulk, probably due to possible

reactions between SiC and SiO2, resulting in the formation of CO or CO2. Analysis of

the composition of the oxide layer of HP materials by SEM/EDS revealed that the

dark phase is predominantly composed of Si and O indicating the presence of SiO2

while the brightest phase is composed of Y, Si and O indicating the presence of

Y2Si2O7. In materials SiCHP and SiC-1650 an Aluminum was also analysed

indicating probably various aluminium silicates or yttrium aluminium silicates. SiC

would react first with O2 to form SiO2 and CO. Then, part of the SiO2 would react

with YAG (Y3Al5O12) to form Al2Si2O7 and Y2Si2O7. The formation of these silicates

result in a zone depleted of Y3+

and Al3+

cations immediately bellows the oxide

Samples Sintering HV10 KIC(Anstis) W kp

(GPa) (MPa.m1/2) (mg/cm2) (mg2/cm4h)

SiC-HP HP (1850 °C/1 h) 20.2 ± 0.9 2.90 ± 0.2 3.24 5.15.10-2

SiC-1650 HP+AN (1650 °C/5 h) 19.7 ± 0.3 3.35 ± 0.2 3.17 4.91.10-2

SiC-1850 HP+AN (1850 °C/5 h) 21.4 ± 0.3 4.54 ± 0.4 1.73 1.47.10-2

SiC-RHP20 RHP (1850 °C/20 min) 16.8± 0.7 3.66 ± 0.14 1.01 4.88.10-3

SiC-RHP60 RHP (1850 °C/60 min) 18.7± 0.3 3.38 ± 0.3 0.10 4.6.10-5

L8 OXIDATION RESISTANCE OF SiC CERAMICS PREPARED BY DIFFERENT PROCEESSING ROUTES

A. Kovalčíková, Z. Lenčéš, J. Sedláček, P. Šajgalík, J. Dusza, M. Ignácz, M. Mihaliková


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scale. It has been proposed that the parabolic oxidation behaviour indicates that the

rate-limiting step is a diffusive process, either by oxygen diffusing trough a growing

silica layer or by RE3+

(Y3+

) cations diffusing trough the oxide layer in the opposite

direction [4]. Only 2 m thickness of oxide layers were created in RHP materials.

The oxidation layer consisted of Si, O and C. SiC-RHP60 sintered 40 min longer

than SiC-RHP20 had the same chemical composition but higher density and better

oxidation resistance. The oxidation resistance of RPH SiC ceramics was significantly

improved due to absence of residual grain boundary phase.

Conclusion

The oxidation of LPS SiC is always passive and protective. The kinetic follows the

parabolic law indicated diffusion as the rate limiting mechanisms. The heat treatment

at higher temperature has a positive effect on oxidation behaviour. SiC materials

sintered by rapid hot pressing without sintering additives show better oxidation

resistance compared with HP SiC because absence of residual grain boundary

phase.

Acknowledgements

This Work was partly supported by Slovak Government through projects VEGA 2/0043/14,

VEGA 2/0189/15, and by APVV-0108-12.

References

[1] L. K.L. Falk, Microstructural development during liquid phase sintering of silicon carbide ceramics, J. Eur. Ceram. Soc. 17 (1997) 983-994.

[2] K. Biswas, G.Rixecker, F.Aldinger, Effect of rare-earth cation additions on the high temperature oxidation behaviour of LPS-SiC, Mater. Sci.Eng. A 374 (2004) 56-63.

[3] A. Kovalčíková, J, Dusza, P. Šajgalík, Influence of the heat treatment on mechanical properties and oxidation resistance of SiC-Si3N4 composites, Ceram. Int. 39 (2013) 7951-7957.

[4] K. Biswas, G. Rixecker, F. Aldinger, Improved high temperature properties of SiC-ceramics sintered with Lu2O3-containing additives, J. Eur. Ceram. Soc. 23 (2003) 1099-1104.

Fig.3: The cross-sections of the oxidized layers, (a) SiC-1650, (b) SiC-RHP60

a) b)

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PYROCHLORE LANTHANUM NIOBATE PREPARED BY SOL-GEL

METHOD IN DIFFERENT SOLVENTS

Helena Bruncková, Ľubomír Medvecký, Pavol Hvizdoš, Juraj Ďurišin,

Vladimír Girman

Institute of Materials Research, SAS, Košice, Slovak Republic


Lanthanum orthoniobate LaNbO4 (LN) precursors were prepared by new polymeric complex

sol-gel method using Nbtartrate in different solvent (ethanol or methanol) and calcination at

750°C. The monoclinic LaNbO4 and orthorhombic LaNb5O14 phases were formed after annealing at 900 °C from precursors synthesized in ethanol (LNet) and methanol (LNmet) solvents, respectively. From Raman spectra resulted different structure at 900 °C. The LN microstructures indicate strong effect of solvent on surface morphology. HRTEM verified the presence of the orthorhombic LaNb5O14 phase with ferroelastic domains.

Keywords: sol-gel, pyrochlore phase, LaNbO4, LaNb5O14, domains.

Introduction

Pyrochlore lanthanum orthoniobate LaNbO4 (LN) ceramics have been shown to be

promising candidates for intermediate temperature proton conducting electrolytes

with potentional applications in solid oxide fuel cells (SOFCs) and hydrogen sensors

[1]. LaNbO4 has a monoclinic structure and transforms into a tetragonal structure at

500 °C. Both polymorphs contain isolated NbO4 tetrahedra interlinked with Laions

of eight-fold coordination to oxygen [2]. The domain structure was suggested to be

related to the rubber-like behaviour of the monoclinic phase. Mokkelbost et al.

reported the preparation of high quality LaNbO4 ceramics by the spray pyrolysis,

calcination at 800 °C and sintering at 1200 °C [1]. Environmentally acceptable the

polymeric complex sol-gel method used for utilization of inorganic salts involves the

preparation of Nbcitrate complex in ethanol solvent [3] and subsequent formation of

viscous LN sol, which is transformed into a gel.

In the present work, we utilize tartaric acid (TA) as chelating agent to prepare

polymeric NbTA complex in the sol-gel process of LaNbO4 precursors with different

solvents (ethanol or methanol) and study their phase composition and nanostructure

after annealing at 900 °C.

Experimental

LaNbO4 precursors were prepared by sol-gel method from La(NO3).6H2O ethanol or

methanol solutions and polymeric Nbtartrate complex solution with stoichiometric

ratio of La : Nb = 1 [4]. LN powder precursors were obtained by the calcination of

xerogel at temperature of 750 °C for 1 hour. After an uniaxial pressing of powders

L9 PYROCHLORE LANTHANUM NIOBATE PREPARED BY SOL-GEL METHOD IN DIFFERENT SOLVENTS

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into pellet form, pellets were annealed in air in closed crucible at the temperature of

900 °C for 6 hours.

The phase composition of samples was analyzed by the X-ray diffraction analysis

(XRD, Philips X´ PertPro, Cu K radiation) and Raman spectra were collected by a

Raman spectroscopy (HORIBA BX 41TF). The microstructures of samples were

observed by the scanning electron microscopy (SEM), (JEOL JSM 7000F) and

transmission electron microscopy (JEOL JEM 2100F).


The XRD diffractograms of LN precursors after calcination at 750 °C and samples

prepared at 900 °C are shown in Fig. 1. XRD analyses verified the formation of

pyrochlore monoclinic LaNbO4 (JCPDS 71-1405) and orthorhombic LaNb5O14

(JCPDS 76-0263) phases and Nb2O5 (JCPDS 72-1484). After annealing at 750 °C,

the major LaNbO4 phase and minor Nb2O5 (both solvents) were revealed in LN

precursors. From the comparison of XRD diffractograms resulted that different final

phase compositions were formed in ethanol or methanol LNc(et) or LNc(met) after

annealing at 900 °C.

The Raman spectra of LN precursors and samples after annealing at 900 °C are

shown in Fig. 2 and exhibit the same bands for both LNp(et) and LNp(met) and sharp

peaks with different intensity for LNc(et) and LNc(met).

Fig. 1: XRD diffractograms of LNet and LNmet precursors after annealing at 750 °C and 900 °C.

The Raman peaks were assigned according to Laguna and Sanjuán [5]. A strong

differences in the Raman spectra of LNc(et) (LaNbO4 structure) and LNc(met)

(LaNb5O14 structure) given by various molecular arrangement result from the

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comparison of spectra at 900 °C. The Raman band of LNc(et) at 640 cm-1

(the NbO

stretching) matches that of the slightly distorted NbO6 octahedral present in Nb2O5

which was verified by XRD analysis. The intensity of peak at 810 cm-1

(structure

LaNbO4) and 670 cm-1

(structure LaNb5O14) was increased in LNc(et) and LNc(met)

respectively.

Fig. 2: Raman spectra of LNet and LNmet precursors after annealing at 750 °C and 900 °C.

Fig. 3 indicates a strong effect of solvent on LN surface morphology after annealing

at 900°C. LN micrographs show differences in particle morphologies, where LNc(et)

(Fig. 3a) had more spherical shape and were composed of large number of

nanosized spherical and more rectangular particles with size up to 100 nm contrary

to LNc(met) exhibit a high porosity and around 1 m irregularly shaped agglomerates

Fig. 3: SEM microstructures of (a) LNet and (b) LNmet after annealing at 900 °C.

L9 PYROCHLORE LANTHANUM NIOBATE PREPARED BY SOL-GEL METHOD IN DIFFERENT SOLVENTS

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(Fig. 3b) with sharp edges and a large fraction of nanoparticles do not exceed the

size of 200 nm.

Figs. 4a and 4b show TEM and HRTEM images, which reveal a heavily domained

structure of the LNmet after annealing at 900 °C. The Fig. 4a confirm elongated shape

of the particle clusters probably of LaNbO4 and LaNb5O14 structures interconnected

with their ferroelastic domains. In HRTEM image, domain walls of the orthorhombic

LaNb5O14 phase are showed.

Fig. 4: TEM image of LNmet prepared at 900 °C (a) particle clusters and (b) HRTEM image of ferroelastic domains of LaNb5O14.

Conclusion

Pyrochlore LaNbO4 (LN) precursors were prepared by new polymeric Nb-tartrate sol-

gel process in ethanol or methanol solvents and calcination at 750 °C. From XRD

analyses resulted, that the different mechanism of phase transformation was

determined from monoclinic LaNbO4 at 750 °C (in both solvents) to major LaNbO4

(in ethanol) and LaNb5O14 (in methanol) phases after annealing at 900 °C.

From Raman spectra and microstructures resulted different structure of lanthanum

niobates at 900 °C. HRTEM observation confirmed the orthorhombic LaNb5O14

phase and ferroelastic domains.

Acknowledgement

This work was supported by the Grant Agency of the Slovak Academy of Sciences through

project No. 2/0041/14.

References

[1] T. Mokkelbost, H.L. Lein, P.E. Vullum, R. Holmestad, T Grande, M.A Einarsrud, Ceram. Int. 35 (2009) 2877-2883.

[2] Prytz, O., Tafto, J.: Acta Mater. 53 (2005) 297-302. [3] Y.J. Hsiao, T.H. Fang, Y.S. Chang, Y.H. Chang, CH. Liu, L.W. Ji, W.Y. Jywe, J. Luminescence 126 (2007)

866-870. [4] H. Brunckova, L. Medvecky, P. Hvizdos, V. Girman, J. Sol-Gel Sci. Technol. 69 (2014) 272-280. [5] M.A. Laguna, M.L. Sanjuán, Ferroelectrics 272 (2002) 63-68.

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INFLUENCE OF MILLING ON DC CONDUCTIVITY OF ILLITE

Štefan Csáki1,2

, Viera Trnovcová1, Ján Ondruška

1, Igor Štubňa

1

1Department of Physics, Constantine the Philosopher University, A. Hlinku 1, 949 74 Nitra, Slovakia 2Department of Physics of Materials, Charles University in Prague, Ke Karlovu 5, 121 16

Praha 2, Czech Republic


The illitic clay (80 % illite, 4 % montmorillonite, 4 % orthoclase and 12 % quartz) was milled in

a ball mill for 0, 30, 60, 90, 120 and 180 min. From the milled clay, samples for

thermoelectrometry, TGA and DSC were prepared. It was found the dehydroxylation of illite is

a 2-step process, and it is shifted to lower temperatures with increasing time of milling. The

thermogravimetry measurements showed the dehydroxylation mass loss does not depend on

the milling time. The dependence of DC conductivity of green samples varies with different

milling time. Up to 250 °C, H+ and OH

- ions are dominant charge carriers. When physically

bound water is removed, the concentration of the H+ and OH

- ions decreases and dominant

charge carriers become mainly K+ and Na

+.

Keywords: illite, milling, thermoelectrometry, DC conductivity

Introduction

The processing of raw materials in ceramic products requires a preparation of

ceramic powders. The shape and size of the crystals of phyllosilicates influence

properties of ceramic materials significantly during preparing the green ceramic body

as well as during firing. The size of crystals of kaolinite, illite and other clay minerals

determines plastic qualities, shrinkage rate, forming potential and drying

characteristics. A larger part of bigger crystals in clay results in a less plasticity, less

shrinkage and faster drying. A larger part of smaller crystals shifts phase

transformations to lower temperatures. Therefore, the milling of hard minerals (such

as quartz, feldspar and others) as well as crushing of phyllosilicate clumps (mainly

from kaolinite and illite), are necessary technological steps in a manufacturing

traditional ceramics [1-3]. A reasonable milling time leads to an increase of the

specific surface area. After a long milling, changes in some clay mineral components

(e.g. muscovite) can be observed because of its transformation into illite-1M, quartz,

albite and others [4]. The dry milling of kaolinite is also a possible way how to obtain

a nanomaterial [5, 6].

A firing process and sintering are significantly influenced by crystallinity of the

phyllosilicates. The low crystallinity and small size of crystals help to decrease the

maximum firing temperature [7-10] as well as the temperature of phase

transformations [11]. It is important to know how the milling influences a firing

process and physical properties of the ceramic samples. Influence of a milling time

on XRD, TGA and DTA was studied for kaolin in [5, 6, 12, 13]. It was found that the

DTA effects of dehydroxylation shifted to lower temperatures, decreased in intensity

L10 THE INFLUENCE OF MILLING ON DC CONDUCTIVITY OF ILLITE

Š. Csáki, V. Trnovcová, J. Ondruška, I. Štubňa


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with increasing milling time. They disappeared completely after 120 min of milling

[5]. The activation energy of dehydroxylation became lower with the increase of the

milling time [12,14]. The temperature of the first exothermic effect, which is

characteristic for a release of the physically bound water (PBW), shifted slightly to

lower temperatures after milling [5,13]. It was also found that milling increased a

quantity of the mullite in porcelain [15] and mullite was transformed from kaolinite at

lower temperature [16]. The influence of milling on firing is relatively well known for

system kaolinite – quartz/alumina – feldspar that is connected with manufacturing of

porcelain. This influence was less studied for ceramic mixtures with illite. In this

contribution, the influence of milling on results of the TGA, DSC, and DC

conductivity of illitic clay are shown and discussed.

Experimental

Samples were made from the illitic clay supplied from the mine in Füzérradvány,

Hungary (Table 1). This clay is composed of 80 % illite, 4 % montmorillonite, 4 %

orthoclase and 12 % quartz.

Table 1: The chemical composition (in wt. %) of illite from Füzérradvány, Hungary

SiO2 Al2O3 Fe2O3 TiO2 CaO MgO K2O Na2O L.O.I.

58.0 24.0 0.6 0.05 0.38 1.70 7.85 0.10 7.3

To obtain a powder with small particles, the clay was crushed and dried at 120 °C

for 1 hour and then milled. In our experiments we used a mill RETSCH PM 100 with

alumina balls and vessel in which the crushed illitic clay was milled for different time

periods 0, 30, 60, 90, 120, and 180 min (samples are called as M0, M30, M60, M90,

M120, and M180, respectively). A possible way how to obtain a relatively

homogeneous batch of the milled clay is a sieving. Our experience [17] showed that

the coarser screen was used, the more quartz and feldspar was present on the

screen. To prevent differences between mineral compositions of the samples after

different milling times, we did not sieve the milled clay.

Samples for electrical conductivity were prepared from wet plastic mass as prisms

with dimensions of 10×10×20 mm with inserted parallel platinum wire electrodes.

The sample arrangement gave a good electrical contact between the measured

material and platinum electrodes during the whole measurement cycle as

experienced in [18, 19].

The measuring circuit was fed from the stabilized voltage source Tesla BS 525 with

10 V. The voltage was recorded by the multimeter Agilent 34972A and the current

was recorded by electrometer Keithley 6514 [19]. Measurements were performed up

to 1100 °C with heating rate 5 °C/min in the air: 1) the first run on the green

samples, 2) the second run on the samples used in the first run, it means that they

were previously heated to 1100 °C. To reach identical initial conditions for DC

conductivity measurements, the samples were placed into a measuring cell and

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heated up to 120 °C for 2 h. Then the heating was switched off and temperature

dependence of the DC conductivity was measured from the temperature 40 °C up to

1100 °C.

We used powder samples for TGA and DSC that were performed up to 1100 °C with

a heating rate 5 °C/min using analysers Mettler Toledo TGA/SDTA 851 and Netzsch

DSC 404.


The results of DTG obtained on powder samples of the mass 50 mg are visible in

Fig. 1. In spite of expectation, the sample M180 contains the smallest amount of the

(PBW). When we look at the mass loss at the temperature 200 °C (i.e. when most

PBW is removed), no regular dependence on the milling time is observed:

m/m0 = 2.65 % for M120, m/m0 = 2.25 % for M0, m/m0 = 2.05 % for M60 and

m/m0 = 1.85 % for M180. We observe 3 steps of the remove of PBW which are

represent with minima at 80 °C, 160 °C (which survives only for M0 and M60) and

260 °C.

Fig. 1: Thermogravimetric results (DTG) for samples M0, M60, M120 and M180

Dehydroxylation of illite is a typically 2-step process with 2 endothermic minima [20,

21]. The mean peak is shifted to lower temperatures with an increasing milling time.

We also observe that the longer milling the lower temperature of the start of

dehydroxylation. We suppose, in analogy with kaolinite, that the milling increases

the concentration of defects in illite and reduces the crystalline size. Both of them

decrease the activation energy of dehydroxylation as found for kaolinite [12, 14], we

can assume the same for illite. It shifts the dehydroxylation to lower temperatures. In

the cases the mass loss is 4 % due to dehydroxylation. Contrary to kaolinite [5],

our results do not show a clear decrease of the dehydroxylation mass loss with

increasing milling time. It can be explained by differences between illite and

kaolinite. The illite crystals are smaller than kaolinite ones and they often have




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51

irregular form and probably many defects [22]. Such crystals are less affected by the

milling than relatively well developed kaolinite crystals.

Qualitatively DSC curves (Fig. 2), which describe removing PBW and

dehydroxylation, are in an agreement with DTG: they also confirm the shift of

dehydroxylation to lower temperatures with increasing time of the milling, but the

lower-temperature peak’s temperatures and their shifts registered by these two

different methods are not the same, see Tab. 2.

In addition to endothermic minima in Fig. 2, which belong to removing PBW and

dehydroxylation, the another endothermic process is characterized by the minimum

between 750 °C and 1100 °C. Actually, sintering, vitrification and high-temperature

reactions (creation of Al-Si spinel and mullite) take place here [21, 23].

Fig. 2: DSC results for samples M0, M60, M120 and M180

Table 2: DSC and DTG peak temperatures in a dehydroxylation region

Milling time [min]

DSC peak temperature

[°C]

Shift of DSC peak [°C]

DTG peak temperature

[°C]

Shift of DTG peak [°C]

0 571 0 557 0

60 552 19 514 43

120 527 44 493 64

180 502 69 489 68

The results of the thermoelectrometry of the green illite samples are shown in Fig. 3.

Up to 250 °C, the DC conductivity is significantly influenced by PBW. In this

temperature range, H+ and OH ions are dominant charge carriers. When PBW is

removed by heating, a concentration of the H+ and OH ions decreases and

dominant charge carriers become mainly K+ and Na

+ [18]. The presence of these

ions is confirmed by chemical analysis (Tab. 1). We suppose that the most important

charge carriers are K+ ions because potassium is a component of illite.

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It follows from the temperature dependencies of DC conductivity that the DC

conductivity increases with increasing time of milling. The conduction activation

energy in this temperature region (250 °C – 440 °C) is 0.85 eV which is typical for

movement of alkali ions in silicate materials. We can suppose that the milling looses

the tightly bound alkali ions, not only from illite but also from feldspar, to sites where

these ions are only loosely bound and are prepared to participate in the conduction

process. The longer is the milling time the higher is concentration of these “free”

alkali ions, and, consequently, the higher is the DC conductivity. An anomaly in the

dependencies around 450 °C can be connected with beginning of dehydroxylation.

Above 900 °C, an influence of the milling on DC conductivity is negligible.

After heating up to 1100 °C, temperature dependencies of the DC conductivity,

measured in the temperature range 40 – 1100 °C, show a more simple character

(Fig.4).

Fig. 3: Temperature dependences of the DC conductivity, in the 1st run, for samples: M30 (○), M60 (), M90 (), and M180 ()

Fig. 4: Temperature dependences of the DC conductivity, in the 2nd run,for samples: M30 (○), M60 (), M90 (), and M120 ()




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Up to 450 °C, the conduction energy is equal to 0.7 eV. However, the dependence

of the conductivity on the milling time is not systematic. The conductivity is highest

after 120 min milling. For all milling times, DC conductivity of fired samples is much

lower than that of green samples. However, at 1050 °C, the DC conductivity of all

samples, green and fired, milled for arbitrary milling times, is equal.

Conclusions

The illitic clay (80 % illite, 4 % montmorillonite, 4 % orthoclase and 12 % quartz) was

milled in a ball mill for 0, 30, 60, 90, 120 and 180 min. From the milled clay, samples

for thermoelectrometry, TGA and DSC were prepared. It was found

Dehydroxylation of illite is a 2-step process as confirmed by DSC and DTG.

The dehydroxylation is shifted to lower temperatures with increasing time of

milling.

The dehydroxylation mass loss does not depend on the milling time.

The DC conductivity of green samples increases with increasing time of milling.

The DC conductivity of the fired samples is significantly lower than the DC

conductivity of the green samples.

Up to 250 °C, H+ and OH ions are dominant charge carriers. When physically

bound water is removed, a concentration of the H+ and OH ions decreases and

dominant charge carriers become mainly K+ and Na

+.

Acknowledgement

This work has been supported by the grant VEGA 1/0162/15. The authors are also indebted

to Mr. J. Biber (Inter-ILI Engineering Office, Hungary) for illitic clay.

References

[1] M.N. Rahaman: Ceramic processing and sintering. Marcel Dekker Inc. New York 2003.

[2] S. Ferrari, A.F. Gualtieri: The use of illitic clays in the production of stoneware tile ceramics. Applied Clay Science, 32 (2006) 73-81.

[3] S .Kavalci, E. Yalmar, S. Akkurt: Effects of boron addition and intensive grinding on synthesis of anortite ceramics. Ceramics International, 34 (2008) 1629-1635.

[4] L.A. Gömze: Mechanochemical phenomena during fine comminution of clay minerals for ceramic bricks and roof-tiles. Materials Science Forum 659 (2010) 19-24.

[5] R.L. Frost, E. Horvath, E. Mako, J. Kristof: Modification of low- and high- defect kaolinite surfaces: implications for kaolinite mineral processing. Journal of Colloid and Interface Science, 270 (2004) 337-346.

[6] R. Hamzaoui, F. Muslim, S. Guessasma, A. Bennabi, J. Guillin: Structural and thermal beghavior of proclay kaolinite using high energy ball milling process. Powder Technology, 271 (2015) 228-237.

[7] J. Dubois, M. Murat, A. Amroune, X. Carbonneau, R. Gardon: High-temperature transformation in kaolinite: the role of the crystallinity and of the firing atmosphere. Applied Clay Science, 10 (1995) 187-198.

[8] M. Nakahara, Y. Kondo, K. Hamano: Effect of particle size of powders ground by ball milling on densification of cordierite ceramics. Journal of the Ceramic Society of Japan, 107 (1999) 308-312.

[9] J. Ranogajec, M. Djuric, M. Radeka, P. Jovanic: Influence of particle size and furnace atmosphere on the sintering of powder for tiles production. Ceramics-Silikáty, 44 (2000) 71-77.

[10] S.N. Monteiro, C.M.F. Vieira: Solid state sintering of red ceramics at lower temperatures. Ceramics International, 30 (2004) 381-387.

[11] O.V. Andryushkova, O.A. Kirichenko, V.A. Ushakov, V.A. Poluboyarov: Effect of mechanical activation on phase transformations in transition aluminas. Solid State Ionics, 101-103 (1997) 647-653.

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[12] P. Ptáček, T. Opravil, J. Wasserbauer, J. Másilko, J. Baráček.: The influence of structure order on the kinetics of dehydroxylation of kaolinite. Journal of the European Ceramic Society 33 (2013) 2793-2799.

[13] P.J. Sánachez-Soto, M. Del Carmen Jiménez de Haro, L.A. Pérez-Magueda, I. Varona, J .L. Pérez-Rodriguez: Effects of dry grinding on the structural changes of kaolinite powders. Journal of the American Ceramic Society, 83 (2000) 1649-1657.

[14] I. Štubňa, G. Varga, A. Trník: Investigation of kaolinite dehydroxylation is still interesting. Épitoanyág, 58 (2006) 6-9.

[15] E. Elmas, K. Yildiz, N. Toplan, H.O. Toplan: Effect of mechanical activation on mullite formation in an alumina-quartz ceramics system. Materiali in Tehnologije, 47 (2013) 413-416.

[16] S. Koc, N. Toplan, K. Yildiz, H.O Toplan: Effects of mechanical activation on the non-isothermal kinetics of mullite formation from kaolinite. J. Thermal Analysis and Calorimetry, 103 (2011) 791-796.

[17] J. Ondruška, M. Jankula.: Influence of the sieving process on thermophysical properties of brick clay. In: Conf. Proceedings 18th Thermophysics 2013, Podkylava, SAV Bratislava 2013, 90-95.

[18] R. Podoba, I. Štubňa, V. Trnovcová, A. Trník: Temperature dependence of DC electrical conductivity of kaolin.J. of Thermal Analysis and Calorimetry, 118 (2014) 597-601.

[19] I. Štubňa, V. Trnovcová, L. Vozár, Š. Csáki: Uncertainty of the measurement of DC conductivity of ceramics at elevated temperatures. Journal of Electrical Engineering, 66 (2015) 33-38.

[20] A.F. Gualtieri, S. Ferrari: Kinetics of illite dehydroxylation, Physics and Chemistry of Minerals, 33 (2006) 490-501.

[21] T. Húlan, I. Štubňa, A. Trník, P. Bačík, T. Kaljuvee, L. Vozár: Thermomechanical analysis of illite from Füzérradvány. Materials Science (Medžiagotyra), submitted article.

[22] Z. Pécskay, F. Molnár, T. Itaya, T. Zelenka.: Geology and K-Ar geochronology of illite from the clay deposit at Füzérradvány, Tokaj mts., Hungary, Acta Mineralogica-Petrographica, 46 (2005) 1-7.

[23] R. Ori: The mineralogical and technological characterization of illite from Füzérradvány (Hungary) as a raw material for traditional ceramics. PhD thesis, University of Modena and Emilia Region, Modena, (2003) (in Italian).

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EFFECT OF HEAT TREATMENT ON MECHANICAL PROPERTIES

AND CORROSION RESISTANCE OF LITHIUM DISILICATE

DENTAL GLASS CERAMICS

Anna Švančárková1,2

, Dagmar Galusková1, Petra Gaalová

1, Ján Balko

3,

Martin Fides3, Dušan Galusek

1

1Vitrum Laugaricio – Joint Glass Center of the Institute of Inorganic Chemistry, SAS, Alexander Dubček

University of Trenčín, and Faculty of Chemical and Food Technology, Slovak University of Technology,

Študentská 2, 911 50 Trenčín, Slovakia 2Faculty of Chemical and Food Technology STU, Radlinského 9, 812 37 Bratislava, Slovakia

3Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, SK-043 53 Košice,

Slovakia


The effect of corrosion in 10 % citric acid and in 1 % citric acid mixture with tartaric acid

(5.9 g/L) on microhardness and wear resistance of lithium disilicate dental ceramics was

studied. Three types of lithium disilicate glass ceramics were observed before and after

corrosion process. Partially crystallized lithium metasilicate glass ceramics were sintered to

fully crystallized lithium disilicate dental ceramics using three different two-stage heat

treatment regimes: A – 500 ºC/2 h, 850 ºC/2 h; B – 500 ºC/1 h, 820 ºC/1 h; C – 850 ºC/1 h.

The temperatures of the two stage regime were determined on the basis of data obtained

from literature, where the first temperature was applied to nucleate lithium disilicate, while the

second temperature was used to grow the lithium disilicate crystals. This way the materials

with different microstructures and various mechanical properties and corrosion resistance

were prepared. The uncorroded samples A showed the lowest wear values. Corrosion

process in 10 % citric acid caused a decrease in wear resistance at all monitored samples.

Keywords: lithium disilicate glass ceramics, corrosion, microhardness, wear resistance

Introduction

Glass ceramics of lithium disilicate-type needed to be developed for CAD/CAM

applications. As lithium disilicate is very difficult to machine with diamond tools and

the base glass is too brittle, other procedures had to be explored in order to allow

this glass-ceramic to be machined with CAD/CAM equipment. This challenge was

met with the development of a material containing an intermediate phase, lithium

metasilicate, which is precipitated in a heat treatment process. The glass ceramics

produced in this way shows suitable machining properties. In its intermediate stage

the material has a bluish color but exhibits very low chemical durability. However,

these properties change significantly during the crystallization process at 850 ºC, in

which the lithium metasilicate is transformed into a durable lithium disilicate glass

ceramics with dental color [1]. Nucleation and crystallisation play a decisive role in

determining the resulting properties of a given glass after heat treatment [2]. The

chemical durability of ceramic materials is basically good, but it may by influenced


EFFECT OF HEAT TREATMENT ON MECHANICAL PROPERTIES AND CORROSION RESISTANCE L11

OF LITHIUM DISILICATE DENTAL GLASS CERAMICS A. Švančárková, D. Galusková, P. Gaalová, J. Balko, M. Fides, D. Galusek

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by many factors, such as the composition and microstructure of the ceramic

material, the chemical character of the corrosive medium, the exposure time, and

the temperature [3].The chemical durability of dental ceramics is an important

property affecting their clinical performance since they are constantly exposed to a

variety of aqueous environments [4].

Experimental part

The microhardness was determined on mechanically polished surfaces of dental

ceramics before and after corrosion at the load 0.2 N. Wear testing was carried by

tribometry measurements under dry conditions using ball-on-disc technique. The

tribological partner was a highly polished alumina ball with 6.35 mm diameter, the

applied load was 10 N, sliding speed 10 cm/s and sliding distance was 50 m. The

influence of corrosion process on Vickers hardness and wear resistance of dental

glass ceramics was monitored on the basis of results from the tests carried out

under dynamic and under static conditions. Both types of test were performed at the

temperature corresponding to the temperature of human body (37 °C). Under quasi-

dynamic conditions the samples were corroded in 10 % citric acid for 96 h with

exchange of the corrosion medium for fresh one in 12 and 24 h intervals. Under

static conditions the samples were tested for 96 h without a change of the corrosion

medium. In a mixture of 1 % citric and tartaric acid the samples were corroded for

12 h under dynamic conditions. The samples were also corroded for 96 h with

exchange of the corrosion medium for fresh one in 12 h intervals.


The highest microhardness was observed for the sample B (650 HV0.2), the lowest

for the sample C (627 HV0.2) (Fig.1). Wear rates of uncorroded samples A and B

were very similar (5.80 x 10-4

and 5.82 x 10-4

mm3/N.m), wear rate of the uncorroded

specimen C was slightly lower (5.54 x 10-4

mm3/N.m) (Fig.2). The Vickers hardness

decreased slighlty after 24 h of dynamic testing in all measured samples.

Fig.1: The change of hardness of corroded (10 % citric acid, 24 h dynamic test) and uncorroded lithium disilicate glass ceramics

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Decrease of hardness was very similar in all tested samples (Fig.1). For both types

of corrosion conditions (static and dynamic) in 10 % citric acid and for both types of

samples (A and B) the wear rates increased in comparison to uncorroded

specimens. Dynamic testing had higher impact on increase of wear of both types of

lithium disilicate ceramics than static testing (Fig. 3a, Fig. 4a).

Fig.2: Wear of uncorroded lithium disilicate glass ceramics

Fig.3: Wear rates of lithium disilicate glass ceramics type A: a) static (96 h) and dynamic test (12 h test) in 10 % citric acid, b) a comparison of uncorroded sample and the sample dynamically corroded (12 h) in 1 % citric acid (CA) + tartaric acid (TA) and in 10 % citric acid

The highest decrease of wear resistance of the lithium disilicate glass ceramics type

A was observed after corrosion in 10 % acetic acid within 96 h in 12 h dynamic test

(8.73 x 10-4

mm3/N.m) (Fig. 3b). Wear rate after the 12 h dynamic testing in the

mixture of acids was similar to uncorroded specimens (Fig. 3b).The corrosion in 1 %

acetic acid + tartaric acid caused increase of wear of the sample C relative to

corrosion process in 10 % citric acid (Fig. 4b).

EFFECT OF HEAT TREATMENT ON MECHANICAL PROPERTIES AND CORROSION RESISTANCE L11

OF LITHIUM DISILICATE DENTAL GLASS CERAMICS A. Švančárková, D. Galusková, P. Gaalová, J. Balko, M. Fides, D. Galusek

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Fig.4: Wear of lithium disilicate glass ceramics type C: a) static (96 h) and dynamic test (12 h test) in citric acid, b) a comparison of uncorroded sample and the sample corroded dynamically (12 h) in 1 % citric acid (CA)+ tartaric acid (TA) and in 10 % citric acid.

Conclusion

This study analyzed the effects of sintering time on the Vickers hardness and wear

before and after corrosion process of lithium disilicate glass ceramics. The sintering

time had impact on hardness of all tested samples, after corrosion in 10 % citric acid

hardness of these samples decreased.

Sintering time had no significant effect on wear of specimens A and B. Samples

sintered by regime C exhibited the lowest wear rates. Corrosion in 10 % citric acid

increased the wear of all sintered samples of lithium disilicate. Dynamic testing had

higher impact on increase of wear.

Acknowledgement

The financial support of this work by the grant VEGA 2/0058/11, and the grant APVV 0218-11

is gratefully acknowledged. This publication was created in the frame of the project "Centre of

excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056, based on

the Operational Program Research and Development.

References

[1] CH. Ritzberger, E. Apel, W . Höland, A. Peschke, V.M. Rheinberger: Properties and Clinical Application of Three Types of Dental Glass-Ceramics and Ceramics for CAD-CAM Technologies. Materials. 3 (2010) 3700-3713.

[2] S. Huang, B. Zhang, Z. Huang, W. Gao, P. Cao: Crystalline phase formation, microstructure and mechanical properties of a lithium disilicate galss-ceramic. J Mater Sci. 48 (2013) 251-257.

[3] W.B. White: Theory of corrosion of glass and ceramics. In: Clarke, D.E. and Zoitos, B.K., Eds. Corrosion of Glass, Ceramics and Superconductors. Noyes Publications, Park Ridge 2 (1992) 2-28.

[4] K.J. Anusavice, N. Zhang: Chemical durability of Dicor and lithia-based glass-ceramics. Dent Mater. 13 (1997) 3-19.

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MECHANICAL AND MICROSTRUCTURAL CHARACTERIZATION

OF POROUS SILICON NITRIDE BIOMATERIALS

Zuzana Pramuková1, Monika Kašiarová

1, Magdaléna Precnerová

2,

Miroslav Hnatko2, Pavol Šajgalík

2

1Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, Košice, Slovakia 2Institute of Inorganic Chemistry, Slovak Academy of Sciences, Dúbravská cesta 9, Bratislava, Slovakia


This work deals with mechanical and microstructural characterization of porous silicon nitride

biomaterials used as trabecular bone replacement. Ceramic composites (Si3N4 +

Hydroxyapatite (HA)) with defined porosity, pore sizes and mechanical properties similar to

human trabecular bone were prepared by different processing routes (air-sintering and

reaction-sintering). Microstructural (porosity, pore sizes, density) and mechanical (Young’s

modulus, hardness, compressive strength) properties of porous samples were investigated

and compared to human trabecular bone. Results showed that prepared porous ceramic

scaffolds had suitable matrix for bone replacement. The values of Young’s modulus and

compressive strength of reaction-sintered samples (SRBSN) were approximately two times

higher and hardness was higher about 21 % in comparison with the samples prepared by air-

sintering (ASSN). The samples coated with HA had about two times higher values of all

mechanical properties than the samples without HA.

Keywords: silicon nitride, biomaterials, trabecular bone, porous scaffold

Introduction

Among other, porous materials can be used as bone scaffolds in biomedical

engineering. The porous ceramics used as trabecular bone replacement should

have porosity more than 45 vol.% and the pore sizes should be larger than 100 µm

in the macro-scale (for cell ingrowth and migration) and lower than 50 µm in the

micro-scale (for transport of nutrients and oxygen). Moreover, three-dimensional

interconnected porous network with interconnected pores is required due to tissue

ingrowth and fixation of porous scaffolds on the skeleton [1].

Porous silicon nitride ceramic is ideal material for bone replacement due to its

superior combination of mechanical, biological and microstructure properties. Silicon

nitride is well known non-oxide ceramics owing with unique mechanical properties,

such as high strength, hardness, fracture toughness, chemical stability and wears

resistance [2]. Currently, silicon nitride is used in biomedical applications (e.g.

surgical screws/plates or prosthetic hip and knee joints [3]) due to its biocompatibility

and bioinertness. Moreover, Si3N4 is non-magnetic, which allows magnetic resonant

imaging and visible on plain radiographs as a partially radiolucent material.

The aim of this work is to present different ways for the preparation of porous silicon

nitride ceramics and achieve satisfactory biological, mechanical and microstructural

properties for human trabecular bone replacement by producing of composites.

MECHANICAL AND MICROSTRUCTURAL CHARACTERIZATION OF POROUS L12 SILICON NITRIDE BIOMATERIALS Z. Pramuková, M. Kašiarová, M. Precnerová, M. Hnatko, P. Šajgalík

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Experimental materials

The samples of human trabecular bone were taken from around 40 year old human

without any severe diseases in anamnesis, sterilized in 96 % alcohol and dried at

45 °C. The bones were obtained in the form of rectangular samples with dimensions

10 x 10 x 20 mm3. Porous ceramic samples were prepared by replication method of

polyurethane sponge as pore-forming agent. Reaction-bonded sintering and air-

sintering were used as different ways of preparation.

Air-sintered samples (ASSN) were prepared by infiltration of suspension (43 vol.% of

solid phase) with silicon nitride powder (Yantai, Tomley Hi-Tech Ind. & Tra. Co.,

China, Ltd.) into the polyurethane sponge, pre-sintered at 1100 °C for 30 minutes in

air, subsequently re-infiltrated with suspension of 10 vol.% Si3N4 and 15 wt.% SiO2

(LUDOX® TM-40, W.R. Grace & Co.-Conn, USA), calcined at 1100 °C for 30 min in

N2, again re-infiltrated with suspension of Si3N4 and SiO2 and sintered at 1200 °C for

2 hours in N2.

Reaction-sintered samples (SRBSN) were prepared by infiltration of suspension

(48 vol.% of solid phase) with mixture of silicon powder (Sicomill silicon powder,

grade 2D, Vesta ceramics AB, Ljungaverk, Sweden) and silicon nitride powder (SN-

E10, Ube Industiels, Japan) in mass ratio 4 : 1 into the polyurethane sponge.

Samples were nitrided at 1390 °C for 3 hours and sintered at 1650 °C for 1 hour in

nitrogen atmosphere.

Hydroxyapatite was prepared using sol-gel method ([Ca(NO3)2∙4H2O + Et-OH] +

[TEP + Et-OH + H2O]) at 85 °C/8 hours in oil bath. Sintered samples were two times

infiltrated by hydroxyapatite sol in vacuum for 0.5 hour and calcined at 900 °C/1 hour

in air.

Experimental methods

The microstructure was studied using scanning electron microscope Jeol JSM 7000

F. The average macro-pore sizes were evaluated through statistical analysis of

images obtained from a camera using the image analysis software (ImageJ).

Porosity and density were determined by Archimedes’ method in distilled water

according to ČSN 725010 standard [4]. Brinell’s hardness of studied porous samples

were carried out using universal mechanical tester UMT2 by BrukerNano with Al2O3

ball (6.35 mm). Maximum force Pmax = 100 N and holding time in maximum force thold

= 10 s was selected for measurements. Young’s modulus was measured by

resonant frequency technique using equipment Buzz-o-Sonic 5.9 by BuzzMac®

International. Compressive strength was determining using machine LR5Kplus by

LLOYD Instruments.


Porous silicon nitride scaffolds prepared using two ways of the preparation were

investigated. Fig. 1 shows characteristic macrostructures of studied porous Si3N4

ceramics and trabecular bone. It is obvious that samples had similar porosity and

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pore sizes in comparison with the studied trabecular bone and too each other, which

was also confirmed by the results summarized in Tab. 1.

Average macro-pore size of trabecular bone was 885 ± 404 µm and ranging from

425 ± 190 to 495 ± 283 µm for studied porous materials. Overall porosity of bone

was 71.5 ± 23.5 vol.% and ranging from 62.9 ± 2.81 to 72.4 ± 4.73 vol.% for

prepared samples (Tab. 1). Porosity of HA samples decreased in comparison with

the samples without HA coating, what was expected due to the filling of some micro-

pores by HA sol. It is clear that pores are relatively large and interconnected (Fig. 1)

what is necessary for bone cells ingrowths. It is effort to achieve interconnectivity in

the prepared porous materials.

Fig. 1: Typical microstructures of the studied human trabecular bone (a) and prepared porous Si3N4 materials: (b) air sintered sample and (c) reaction-sintered sample

Table 1: Microstructural characteristics of the studied porous ceramics and bone

Samples Density

(g/cm3)

Overall porosity

(vol.%)

Open porosity

(vol.%)

Macro-pore sizes

(µm)

ASSN 0.96 ± 0.09 69.9 ± 3.06 66.6 ± 3.42 464 ± 255

SRBSN 0.89 ± 0.15 72.4 ± 4.73 70.4 ± 6.90 495 ± 283

ASSN-HA 1.08 ± 0.07 66.3 ± 2.29 63.1 ± 2.60 425 ± 190

SRBSN-HA 1.18 ± 0.09 62.9 ± 2.81 61.0 ± 2.94 477 ± 210

BONE 0.45 ± 0.15 71.5 ± 23.5 64.1 ± 25.9 885 ± 404

a)

b) c)

MECHANICAL AND MICROSTRUCTURAL CHARACTERIZATION OF POROUS L12 SILICON NITRIDE BIOMATERIALS Z. Pramuková, M. Kašiarová, M. Precnerová, M. Hnatko, P. Šajgalík

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Comparison of the measured mechanical properties of studied materials with human

trabecular bone is summarized in Fig. 2. Hardness of trabecular bone was

9.26 ± 3.22 MPa, Young’s modulus was 2.43 ± 0.46 GPa and compressive strength

was 2.71 ± 1.31 MPa. Air-sintered samples (ASSN) had lower values of all

measured mechanical properties in comparison with the samples prepared by

reaction-sintering (SRBSN) due to lower sintering temperature and higher amount of

SiO2 in the air-sintered sample. The values of Young’s modulus and compressive

strength of reaction-sintered samples (SRBSN) were approximately two times higher

and the value of hardness was about 21 % higher in comparison with the samples

prepared by air-sintering (ASSN). The samples coated with HA had also about two

times higher values of all mechanical properties than the samples without HA. It was

caused by the filling of micro-structural defects (some voids and cracks) present in

the matrix of prepared porous ceramics, which are typical for replication method [5].

a) b)

c)

Fig. 2: Comparison of the measured mechanical properties of studied materials with trabecular bone: a) Hardness; b) Young’s modulus; c) Compression strength

The compressive strength of studied materials was close to the strength of

trabecular bone while average values of hardness and Young’s modulus were

slightly higher. Further optimization of the mechanical properties will be possible by

proper addition of SiO2 to the samples and by suitable temperature regime.

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Conclusions

Prepared porous ceramic materials had suitable matrix for bone replacement in view

point of microstructural and mechanical properties. The values of Young’s modulus

and compressive strength of reaction-sintered samples (SRBSN) were

approximately two times higher and hardness was higher about 21 % in comparison

with the samples prepared by air-sintering (ASSN). The samples coated with HA

had about two times higher values of all mechanical properties and slightly lower

porosity than the samples without HA coating. Porosity and pore sizes and the

compressive strength of studied samples were closed to human trabecular bone.

Acknowledgement

This work was realized within the frame of the projects APVV-0500-10; APVV-0218-11 and

COST Action MP1005 which is supported by the Operational Program “Research and

Development” financed through European Regional Development Fund.

References

[1] P. Sepulveda, J. R. Jones, L. L. Hench, Bioactive sol-gel foams for tissue repair, J. Biomed. Mater. Res. 59 (2002) 340-348.

[2] D. M. Ebenstein, L. A. Pruitt, Nanoindentation of biological materials, NanoToday 1 (3), Elsevier Ltd (2006) 26-33.

[3] B. S. Bal, A. Khandkar, R. Lakshminarayanan, I. Clarke, A. A. Hoffman, M. N. Rahaman, Testing of silicon nitride ceramic bearings for total hip arthroplasty, J Biomed Mater Res B 87 (2008) 447–54.

[4] Československá státní norma, Stanovení nasákavosti, zdánlivé pórovitosti, pórovitosti, objemové hmotnosti a zdánlivé hustoty vypálených keramických směsí a výrobkú, ČSN 725010 (2003).

[5] M. Scheffler, P. Colombo, Cellular Ceramics: Manufacturing, Properties and Applications. Weinheim: Wiley – VCH, 1 (2005) 669. ISBN: 3527323106.

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CORROSION OF NATURAL AND SYNTHETIC BIOMATERIALS

IN ACIDIC MEDIA AND ITS EFFECT ON MECHANICAL PROPERTIES

Petra Gaalová

1, Dagmar Galusková

1, Anna Švančárková

1,2, Ján Balko

3,

Dušan Galusek1

1Vitrum Laugaricio – Joint Glass Center of the Institute of Inorganic Chemistry, SAS, Alexander Dubček

University of Trenčín, and Faculty of Chemical and Food Technology, Slovak University of Technology,

Študentská 2, 911 50 Trenčín, Slovakia 2Faculty of Chemical and Food Technology STU, Radlinského 9, 812 37 Bratislava, Slovakia

3Institute of Materials Research, Slovak Academy of Sciences, Watsnova 47, SK-043 53 Košice,

Slovakia


The influence of corrosion in acid beverages on micromechanical properties and

wear resistance of human teeth (white wine pH ~3.5) and leucite dental ceramics

(10 % citric acid/ pH ~1.6) were studied. Also, the effect of fluorine-containing

mouthwash Elmex (pH ~4.4) in terms of its protective action against corrosion, and

the improvement of mechanical properties and wear resistance at enamel was

studied. The influence of the corrosion on Vickers hardness and wear resistance of

dental enamel and leucite dental ceramics were monitored on the basis of results

from the tests carried out under dynamic conditions.

Keywords: human tooth enamel, leucite dental ceramic, microhardness, wear resistance, corrosion

Introduction

Teeth corrosion is multifactorial negative phenomenon, comprising mutual action of

chemical, biological and behavioral factors, resulting in removal of minerals from the

teeth surface [1,2]. Reduced surface hardness, which accompanies corrosion of the

enamel surface by acid beverages can then result in excessive wear of teeth, and

can be assessed using a physical measurement such as the hardness test and wear

resistance [3]. Various acids, such as tartaric acid, malic acid, lactic acid, ascorbic

acid, citric acid and phosphoric acid, which are a common component of beverages,

including wine, have a strong destructive effect on the surface of teeth and dental

ceramics. Biological and chemical factors in the oral environment influence the

progress of dental corrosion. Saliva provides protective effects by neutralizing and

clearing the acid. Suppression of tooth corrosion is possible also by proper mouth

hygiene, with the use of fluoride- containing toothpastes and mouthwashes. Among

the most commonly used mouthwash agents are those produced by Elmex (GABA

International AG, Switzerland). Due to its surface activity and slightly acidic pH,

amine fluoride present in the mouthwash promotes the formation of well-adhering

calcium fluoride globules. These act as a fluoride deposit, which, according to the

producer’s data, both protects the teeth against acid attacks and stimulates

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remineralisation, providing protection from dentine caries. The purpose of this study

was to evaluate the effect of the corrosion in a white wine (pH 3.5) on the hardness

and wear resistance of human teeth enamel. The protective and remineralization

capacity of the commercial mouthwash (Elmex, pH = 4.4) was also evaluated, in

terms of the change of hardness and wear resistance of the teeth treated by the

agent before and after exposure to white wine. In this work, we also studied the

effect of 10 % citric acid on the hardness and wear resistance of leucite dental

ceramic.

Experimental

Human tooth enamel specimens were prepared from extracted human molars

without caries, surgically removed due to orthodontical reasons, originating from

individuals of either gender. Before the use the teeth/ dental ceramics were rinsed in

deionized water and cleaned in ultrasonic bath. The specimens were embedded in a

synthetic resin (Dentacryl/ VersoCit- 2 Liquid) and then ground and polished

(Buehler Ecomet300 /Automet 300). After polishing, the specimens were stored in

distilled water. Immediately before the corrosion tests the specimens were rinsed

with alcohol to dry the exposed surface. The surface hardness of specimens was

measured using a Vickers diamond indenter with a microhardness tester (WIKI 200)

under a load of 200 gf (1.961 N) with 5 s dwell at the maximum load. Ten indents

were carried out for each specimen. The wear resistance of specimens was

measured using a nanotrimometer (TTX-NTR2) under a load of 900 mN with Al2O3

ball. The baseline surface microhardness and wear resistance before further

corrosion/recovery treatment was determined for each specimen. The corrosion

tests were conducted in vitro, using white wine (pH = 3.5) as the corroding agent.

The enamel surface was exposed to 100 µl of white wine for 40 minutes in a climatic

chamber (Angelantoni Discovery DY110) at 37 °C and 95 % relative humidity. After

each exposure time the microhardness/ wear resistance was determined. The

mouthwash Elmex (pH = 4.4) was tested for its capacity to influence the

microhardness of the enamel, and to protect human teeth against corrosion. For that

purpose three uncorroded specimens in parallel were individually exposed to 100 µl

of the mouthwash for 40 minutes in the climatic chamber under the same conditions

as above, and microhardness/ wear resistance was determined after every

exposure. The specimen exposed to the mouthwash for 60 minutes was then

corroded in the white wine for 10, 20, 40 and 60 minutes at 37 °C, and the

microhardness/ wear was determined. The samples of dental ceramics were

exposed to 4 ml of 10 % citric acid for 12, 24 and 48 hours in a climatic chamber at

37 °C and 95 % relative humidity. After each exposure time the microhardness/ wear

resistance was determined.


The baseline microhardness of tested tooth specimens showed only a slight

variation, which fell well into the range of the scatter of experimental data, and

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ranged between 340 and 360 HV0.2. The sample exposed to 40 minutes of Elmex

and then 40 minutes in white wine showed a decrease in microhardness from 350 to

320 HV0.2. The measurements confirmed a deterioration of microhardness from 350

HV0.2 in uncorroded specimens to 220 HV0.2 in surfaces exposed for 40 minutes in

white wine. Significant recovery of Vickers hardness was observed after 40 minutes

exposition time in white wine and then 40 minutes in the Elmex fluoridation solution,

with the increase from 220 to 340 HV0.2 (Fig. 1). The wear resistance was

determined on natural surfaces of tooth enamel and compared to the values

obtained from mechanically polished surfaces before and after dynamic corrosion

tests (Fig. 2). For both types of samples (polished / unpolished) a decrease in the

wear was observed by the action of Elmex and an increased wear after corrosion in

white wine. Healing of corrosion-induced defects by the action of the Elmex solution

was observed by SEM and confocal scanning microscopy, and was associated with

observed increase of hardness and wear resistance.

Fig. 1: Variation of enamel surface microhardness as a function of the exposure to acidic medium, (N- uncorroded, E- fluoridation Elmex, RB- corroded in white wine)

a) b)

Fig. 2: Variation of enamel surface wear resistance as a function of the exposure to acidic medium, a) human enamel/ polished , b) human enamel/ un polished

(N- uncorroded, E- fluoridation Elmex, RB- corroded in white wine)

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Samples of polished leucite dental ceramics were corroded for 12, 24, 48 h in 10 %

citric acid solution at 37 °C. The Vickers hardness decreased from 640 HV0.2 for

uncorroded specimens to 530 HV0.2 after 48 h exposure to citric acid.

Microhardness of the samples decreased linearly with corrosion time (Fig. 3).

The wear rates increased markedly after 12 h exposure to the acidic media. Further,

approximately linear increase of the wear rates was observed after extend exposure

to corrosion medium (Fig. 4).

Fig. 3: Microhardness of leucite dental ceramics at different time of corrosion/ dynamic test, (N- uncorroded, 12 h, 24 h, 48 h – time of corrosion).

Fig. 4: Wear of leucite dental ceramics as a function of time corrosion in citric acid by dynamic test.

Conclusion

Exposure to acidic media markedly, and negatively, influenced the mechanical

properties of both the natural human teeth and leucite dental ceramics. Healing

effect, in terms of recovery of the mechanical properties of human teeth, was

observed after exposure of the corroded teeth to the Elmex fluoridation solution.

Acknowledgment The financial support of this work by the grants VEGA 2/0058/11, and APVV 0218-11is

gratefully acknowledged. This publication was created in the frame of the project "Centre of

excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056, based on

the Operational Program Research and Development funded from the European Regional

Development Fund.

References

[1] T. Attin, U. Koidl, W. Buchalla, H. G. Schaller, A. M. Kielbassa, E. Hellwig, Correlation of microhardness and wear in differently eroded bovine dental enamel, Archives of Oral Biology. 42 (1997) 243–250.

[2] M. Eisenburger, R.P. Shellis, M. Addy, Comparative study of wear of enamel induced by alternating and simultaneous combinations of abrasion and erosion in vitro, Caries Research. 37 (2003) 450–455.

[3] Sherine B Y Badr, Mohamed A Ibrahim. Protective effect of three different fluoride pretreatments on artificially induced dental erosion in primary and permanent teeth. Journal of American Science. 6, 11 (2010) 442-451. (ISSN: 1545-1003).

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HIGH-GRADE SILICA PREPARED FROM SERPENTINITE TAILINGS

USING TWO-STAGE LEACHING AND PRECIPITATION

Alena Fedoročková, Pavel Raschman, Gabriel Sučik

Technical University of Košice, Faculty of Metallurgy, Letná 9, 04200 Košice, Slovak Republic


This study demonstrates an economically efficient synthesis of highly reactive silica using

two-stage (acid and alkaline) leaching of serpentinite tailings. The alkaline leaching, the

purification of sodium silicate solution and the effect of impurities on the precipitation of

amorphous silica under the acidic and alkaline conditions were examined in more details. The

proposed route is technologically advantageous because of a high purity (99.4 wt.% SiO2),

large specific area (541 m2 g

-1) and consistent quality of the product were achieved.

Keywords: serpentinite tailings, two-stage leaching, metasilicate solution, precipitation, SiO2

Introduction

While many published research papers describe the preparation of amorphous silica by

hydrolysis of alkoxides [1-4], in practice the methods based on precipitation of

amorphous SiO2 from water glass are mostly applied [5]. An alternative (and cheaper)

procedure to prepare highly pure silica powders is through the leaching of serpentinite

tailings.

Serpentinites – the rocks characterized by high contents of silicon oxide (35 45 %

SiO2) and magnesium oxide (35 40 % MgO) - are widely spread in nature and,

moreover, there are enormous reserves (hundreds of millions tons) of serpentinite

available in the form of tailings from mining of asbestos and chromite deposits [6]. The

use of serpentinite tailings is limited and monitored in several countries worldwide

because they can contain fibrous asbestos minerals that may be carcinogenic [7].

Though the acid leaching of serpentinite has been studied by previous authors [6, 8]

(reaction 1):

)(5)(2.3.6)()( 22

2

4523 lOHsSiOaqMgaqHsOHOSiMg , (1)

the quality of silica products obtained this way (including such parameters as purity,

specific surface area, particle size distribution or pore size and shape) was generally not

comparable with that of synthetic amorphous silica prepared by precipitation. Therefore,

the second stage of the process, involving the alkaline leaching of silica concentrate

(reaction 2) and the effect of impurities on the precipitation of amorphous silica (reaction

3) without using any additives is the main focus of this work.

lOHaqSiONaaqNaOHsSiO 2322 ..2 . (2)

.2.2. 4232 aqNasOHSilOHaqHaqSiONa . (3)

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It is hypothesised that the synthesis of silica using this procedure could be economically

more efficient than the production from water glass. Moreover, industrial utilisation of

serpentinite tailings at a large scale could be a way to decreasing the production cost

and avoiding the problems related to the waste disposal, as well.

Material and methods

A sample of serpentinite tailings obtained from the mining and processing of serpentinite

at Dobšiná was prepared by grinding, dry-screening (0 100 µm), and magnetic

separation. The non-magnetic part of raw serpentinite was treated in 10 M hydrochloric

acid (S : L ratio = 2 : 5) to obtain a silica concentrate (reaction 1). The results of chemical

analyses of dried raw serpentinite and silica concentrate are summarized in Table 1.

Table 1: Chemical composition of serpentinite prior to and after acid leaching

Chemical composition (wt.%)

sample MgO CaO SiO2 Fe2O3 Al2O3 NiO Cr2O3 L.O.I.

raw serpentinite 38.2 1.2 37.3 5.7 1.4 0.1 0.3 14.1

silica concentrate 1.5 0.1 87.2 0.4 0.8 0.02 0.04 9.2

Sodium metasilicate solution was prepared by alkaline leaching of the silica concentrate

in 5 M NaOH at 60 °C (reaction 2), and solid-liquid separation by filtration. Samples of

synthetic amorphous silica were synthesized in compliance with chemical reaction 3 by

means of three different precipitating agents: HCl, CO2 and (NH4)2CO3.

The synthesis of silica gel using HCl was carried out by adding a thin stream of 2.5 M

aqueous solution of sodium metasilicate into the bulk of intensively stirred (800 s-1) 10 M

aqueous solution of hydrochloric acid. In the case of silica synthesis through

the carbonization reaction, CO2 (0.13 MPa) or (NH4)2CO3 (in mole ratio = 1 : 2) were

introduced into the bottom of the reactor with a continuously—stirred Na2SiO3 solution.

Finally, the synthesised silica was washed, filtered and dried at 110 °C.

SEM and EDS was used to investigate the morphologies of the silica powders and

semiquantitative elemental analysis, respectively. The exact purity of final products was

acknowledged by quantitative elemental ICP-OES analysis. The specific surface area

was determined using the B.E.T. nitrogen adsorption technique.


In the preparation of pure SiO2 by acid leaching of serpentinite (reaction 1) Velinskii and

Gusev [6], and Pietriková et al. [8] consider the silica concentrate as a pure SiO2.

Though they declared relatively high purity of the prepared silica powders (up to 99 %

SiO2), it was found that the analytical methods used play an important role in

determining the apparent purity of synthesised SiO2. Semiquantitative elemental

microanalysis in this case reported only the presence of silicon and oxygen (Fig.1) and

the silica concentrate appeared to be seemingly pure SiO2. However, according to the

results in Table 1, the purity of the SiO2 concentrate was 87.2 % which, calculated to an

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anhydrous state, was about 96 % SiO2. The main accompanying components were

MgO, Al2O3 and Fe2O3.

Fig.1: SEM micrograph and the results of X-ray microanalysis of silica concentrate.

In order to reduce the content of impurities in the final product, leaching of silica

concentrate in NaOH (reaction 2) was applied. By alkaline leaching of the silica

concentrate a portion of accompanying elements passed into the solution (Table 2),

and a part remained undissolved in the form of insoluble residue, which was

removed by filtration.

Table 2: Chemical composition of a concentrated sodium silicate solution

sample

pH

Chemical composition (g.L-1)

Si4+ Al3+ Ca2+ Fe2+ / 3+ Mg2+ Ni2+ Cr3+ Ti4+

Sodium silicate solution

13.8 69.4 0.03 0.02 0.06 0.08 0.02 0.003 0.01

The process of silica synthesis from sodium silicate solution (reaction 3) strongly

depends upon pH [1, 2]. In highly acidic solution, the silicic acid polymerizes to form

nuclei that grow, and then aggregate into chains. At higher pH (i.e. at pH > pHZPC ~

2) the dissociation of the surface hydroxyl Si-OH groups (reaction 4) results in a

negative zeta potential with less aggregation [1, 2]. Therefore, the higher the pH, the

lower the surface area of precipitated silica.

lsaqs OHOSiOHOHSi 2 . (4)

The synthesis of amorphous silica by the neutralization of water glass is conventionally

carried out at pH = 8 [9]. Because of high repulsion between primary silica particles, preparation of ultrafine SiO2 powders with a high specific surface area requires additives

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such as a flocculating agent and/or surfactant [2, 10]. In order to stimulate the

aggregation without using additives, the synthesis of amorphous SiO2 was carried out

under the conditions of maximum sol stability (i.e. at pH ≤ 1.0 using HCl and at

pH ≥ 10.0 using CO2 or (NH4)2CO3 as precipitating agent). The characteristics of the

final products are shown in Table 3.

Table 3: Characteristics of the final SiO2 powders

Silica properties

Precipitation agents

HCl (aq.) (NH4)2CO3 (s) CO2 (g)

pH of precipitation

1.0 10.7 10.0

Mo

rph

olo

gy

SA [m2.g-1] 541 414 392

Purity [wt.%] determined

by ICP - OES

99.7 99.4 99.4

From the presented results it can be seen that the samples of synthetic amorphous

silica prepared by double-stage leaching of serpentinite tailings and precipitation were of

high purity (above 99.4 % SiO2) and did not contain residues of the original raw

serpentinite (Fig. 1).

To verify the effect of impurities on the specific surface area of SiO2 products, the use of

both natural sodium silicate solution (prepared from serpentinite tailings) and synthetic

solutions of Na2SiO3 (prepared by dissolving SiO2 powder of p.a. purity in a

concentrated solution of pure NaOH with a molar Na2O : SiO2 ratio = 1 : 1) was tested.

From the comparison illustrated by Fig. 2 it follows that the presence of impurities seems

to have a beneficial effect on the specific surface area; in all cases the values of the

specific surface area of SiO2 powders prepared from serpentine were higher than those

of SiO2 prepared from a synthetic solution of Na2SiO3.

Conclusion

The present study demonstrates a feasible solution for the economically efficient

synthesis of highly reactive silica using hydrometallurgical processing of serpentinite

tailings with the emphasis on solving the problems related to waste disposal.

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Fig. 2: Influence of impurities on the specific surface area of synthetized SiO2 samples.

The main results of the present study can be summarised as follows:

a) Amorphous silica with high purity (above 99.4 %) was synthesized from

serpentinite by a two stage leaching and precipitation.

b) Highly reactive product with surface area from 392 m2 g

-1 to 541 m

2 g

-1,

depending on the precipitating (acidifying) agent used, was obtained.

c) The presence of impurities in the sodium metasilicate solution had a beneficial

effect on the specific surface area.

d) The synthesis of amorphous SiO2 by precipitation at pH ≤ 1 is more advantageous

from the technological point of view, because the characteristics of synthesized

SiO2 powders are less sensitive to the presence of impurities in the sodium

metasilicate solution. Products of consistent quality can be thus prepared from

serpentine-based raw-materials of variable chemical/phase composition.

e) Total yield of silicon in the overall process was 90 – 91 %.

Acknowledgement

This publication was supported by the Slovak Grant Agency for Science (Grant 1/0378/14 and

1/0840/13).

References

[1] Pakhomov N.A., Buyanov R.A., Kinet. Catal. 46 (5), 2005, 669-683. [2] Bergna H.E., Roberts W.O. (Eds.), Colloidal Silica. Fundamentals and applications, Boca Raton, FL, CRC

Press Taylor & Francis Group, 2006, 37–782. [3] Marquez-Linares F., Roque-Malherbe R.M.A., J. Nanosci. Nanotechno. 6, 2006, 1114-1118. [4] Liu J., Zhang L., Yang Q., Li C., Micropor. Mesopor. Mat. 116, 2008, 330-338. [5] Unger, K.K.; Tanaka, N.; Machtejevas, E. (Eds.), Monolithic Silicas in Separation Science: Concepts,

Syntheses, Characterization, Modeling and Application, Weinheim, Germany, Willey-WCH, Verlag GmbH & Co. KGaA, 2011, p. 21.

[6] Velinskii V.V., Gusev G.M.: J. Min. Sci., 38, 2002, 402-404. [7] Bloise, A.et al., Environmental Earth Sciences. 71, 2014, 3773-3786. [8] Pietriková A., Búgel M., Golja M.: Metalurgija .43, 2004, 299-304. [9] Bogoevski S.: Periodica Polytechnica: Chem. Eng. 44, 2000, 133-140. [10] Kobayashi, M.; Juillerat, F.; Galletto, P. et al.: Langmuir 21, 2005, 5761-5769

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IMPACT OF WAY OF PREPARATION ON THE QUALITY

OF CERAMIC PIGMENTS

Žaneta Dohnalová, Petra Šulcová, Nataliia Gorodylova

Department of Inorganic Technology, Faculty of Chemical Technology, University of Pardubice,

Studentská 95, 532 10 Pardubice, Czech Republic,


Present work contains results related to the synthesis of yellow ceramic pigments based on

strontium stannate. Impact of way of preparation on the phase composition and pigmentary

application were studied.

Keywords: solid state reaction, perovskite, yellow pigments, colour

Introduction

Research of authors group is focused on the synthesis and characterization of

inorganic powder materials with high thermal and chemical stability which could find

the industrial application for colouring of ceramics, organic matrices, plastics or

building materials. These materials consist of a stable crystal host lattice doped with

transition metal cations which act as chromophore cation. The chromophores are

the source of colour properties of the materials. In this study, SrSnO3 was the crystal

host lattice doped by terbium cations as chromophores. Compound SrSnO3 belongs

to the family of analogous alkaline-earth stannates, MSnO3 (where M = Ca, Sr and

Ba). In the case of SrSnO3 and CaSnO3 the cubic perovskite structure is disordered

by an octahedral tilting distortion to orthorhombic [1,2]. In response to the octahedral

tilting distortion which is caused by the decreasing size of the alkaline-earth cations,

the conduction bandwidth of these alkaline-earth stannate perovskites (CaSnO3,

SrSnO3 and BaSnO3) strongly decreases. This, in turn, leads to a corresponding

increase in the band gap from 3.1 eV in BaSnO3 to 4.1 eV in SrSnO3 and to 4.4 eV

in CaSnO3 [1].

Taking into account the technological and theoretical interest in SrSnO3 doped with

terbium cations, we prepared new solid solutions of SrSn1-xTbxO3, where x = 0.1 – 1.

The aim of our research was to describe the effectiveness of synthesis route on

chemical substitution in the structure and pigment-application qualities of the

samples.

Experimental part

Four different approaches were used for the synthesis of perovskite strontium

stannates doped by terbium ions. The first method was based on classical ceramic

route, i.e. solid state reaction (SSR). Reaction mixtures containing SnO2, Tb4O7 and

SrCO3 were thoroughly ground in an agate mortar with a pestle, placed in corundum

crucibles and heated. The second variation of solid state reaction was upgraded by

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wet mechanochemical activation (Wet-MA) before heating. The high energy milling

process was carried out in a planetary mill Pulverisette 5 (Fritsch GmbH, Germany)

for 6 hrs. The reaction mixtures were ground in ethanol with agate balls (Ø 10 mm)

in a ball-to-powder weight ratio of 20 : 1. The third synthesis route was based on dry

mechanochemical activation (Dry-MA) before heating at the same conditions. The

last method of preparation was based on suspension mixing of the initial reagents

(SM). The initial reagents and foaming agents (10 wt.%) were converted into an

aqueous suspension (70 wt.%), in a porcelain mortar. The semi-products were

obtained by thermal treatment of the aqueous suspensions on a steel plate at the

temperature of 400 °C. The powdered semiproducts were placed into corundum

crucibles and subjected to the heating process. All mixtures and semi-products were

heated in air atmosphere in two stages. The first heating stage was carried out at

the temperature 1000 °C for 2 h; obtained intermediates were manually ground in an

agate mortar with a pestle. Final products were obtained after the next heating at the

temperatures 1300 °C, 1400 °C or 1500 °C for 3 h. After each heating stage, the

samples were gradually cooled to the room temperature and reground.

The phase composition of the pigments was studied by X-ray diffraction analysis.

Diffractograms of the samples were obtained using a difractometer D8 (Bruker, GB)

with a goniometer of 17 cm in the range 2Θ of 10 – 80°. Cu Kα1 (λ = 0.15418 nm)

radiation was used for the angular range of 2Θ < 35° and Cu Kα2 (λ = 0.15405 nm)

for the range of 2Θ 35°.

The colour properties of all prepared pigments were objectively evaluated by

measuring of spectral reflectance using a spectrophotometer ColourQuerst XE

(HunterLab, USA). The measurement conditions were the following: an illuminant

D65, 10° complementary observer and measuring geometry d/8°. For description of

colour the CIE L*a*b* colour space was used [3].

The particle size distribution of the samples was measured using a Mastersizer

2000/MU (Malvern Instruments, UK).

The assessment of resistance to light of coloured pigments was tested according to

the general methods of the test for pigments described in ISO standard 787–

15:1986. The samples were exposed to simulated sunlight in a Q-Sun Xenon Test

Chamber, model Xe 1 (Lab. Products, USA) for 618 h.


Incorporation of the ions of terbium into the lattice of stannates was studied by X-ray

diffraction analysis. For synthesis routes (SSR, Wet-MA and SM), the calcination of

the reaction mixtures at the temperature 1300 °C did not result in single phase

product. XRD diffractograms contain next to the lines of orthorhombic SrSnO3 with

primitive Pbnm (62) space group (JCPDS card 01-072-7523) and tetragonal

Sr2SnO4 with body centred I4/ mmm (139) space group (JCPDS card 00-24-1241)

also the lines of unreacted terbium oxide TbO2 (JCPDS cars 01-075-0209) [3]

(Figure 1).

XRD analysis of the powders prepared by dry mechanochemical activation (Dry-MA)

and by calcination at the temperatures 1300 °C revealed more interesting results. In

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the whole range of substitution, i.e. x = 0.1 – 1. XRD, single phase solid solution has

been identified. The doping of strontium stannate by terbium cations with bigger

ionic radii resulted in an expansion of the volume of the elementary cell, and

therefore caused a slight shift of the diffraction lines. Single phase solid solutions

were obtained also by Wet-MA and SSR methods with calcination at 1400 °C and

1500 °C. The samples prepared by SM route were no single phase at all

temperatures. The lines of TbO2 next to the lines of SrSnO3 and Sr2SnO4 were

identified at the diffraction patterns of the powders heated at the temperature 1300

and 1400 °C. The calcining temperature 1500 °C resulted in the formation of two

phases: orthorhombic SrSnO3 and Sr2SnO4.

Fig. 1: X-ray diffraction patterns of samples SrSn0.6Tb0.4O3 (T = 1300 °C): a) Wet-MA 1300; b) SSR; c) Dry-MA; d) SM

The comparison of the effect of the synthesis route on the colour properties of

SrSn0.5Tb0.5O3 samples (Table 1) showed that the yellow powders prepared by Dry-

MA and SSR route (1400 °C) are less interesting than the samples prepared by

Wet-MA method and SM method. Colour of these samples contains a less amount

of yellow hue (b*) and a less amount of red hue (a*) in the case of SSR method.

Colour coordinate a* of sample prepared by Dry-MA is even in the green direction (-

a*). Both samples are also darker and less saturated than the samples prepared by

Wet-MA and SM method. From the colouristic point of view, SrSn0.5Tb0.5O3 is the

most interesting yellow sample, especially, prepared by Wet-MA method.

The results of the measuring of particle size of the pigments are also summarized in

Table 1. The values of mean size value (d50) of the pigments vary between 5.3 and

12.5 μm. Results showed that intensive mechanochemical activation of initial

reagents before calcination (both, wet and dry) reduces the final granulometric

composition of samples and forms fine particles (5 – 6 m). The other methods

(SSR and SM) formed the powders with coarser particles.

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Table 1: Impact of way of preparation on the colour properties and particle size of powders SrSn0.5Tb0.5O3 (1400 °C)

Preparation L* a* b* C H° d50 m

SSR 75.78 3.69 37.88 38.06 95.56 12.43

Wet-MA 80.27 5.43 47.41 47.72 99.40 6.22

Dry-MA 70.58 -6.16 39.27 39.75 98.91 5.32

SM 77.86 7.46 46.29 46.89 99.15 12.04

Light stability is a material property. It is defined as the resistance of materials to

discolouration caused by the action of global radiation (daylight phase: standard

illuminant D65). The colour stability at daylight of SrSn0.5Tb0.5O3 powders, which

were prepared by calcination at 1400 °C, was tested. The change of the colour

properties of the powders is expressed by values of total colour difference after

exposition to a radiation load of 973.7 kJ/m2. After irradiation, the samples became

lighter, greener and also contained a higher amount of yellow hue. Sample Wet-MA

exhibits the best resistance to daylight, and its value of the total colour difference is

0.31. It signalizes that the changes of the colour properties are recognizable only by

measurement of spectral reflectance. Slightly worse result brought irradiation of the

sample prepared by SSR method (ΔE*CIE = 0.44), which total difference was also

less than 0.5 and it is imperceptible to the human eye. The significant changes in the

colouration were observed in the sample prepared by SM route. Its value of ΔE*CIE is

1.5 and it is the difference visible by the human eye. The results of testing the light

stability of samples prepared by Dry-MA have not been finished yet.

Conclusion

Present work contains results related to the synthesis of yellow ceramic pigments

based on strontium stannate in which a part of tin ions was substituted by terbium

ions. The pigments SrSn0.5Tb0.5O3 were synthesized by three methods based on

solid state reaction and by suspension mixing of initial reagents. The samples were

characterized by X-ray diffraction analysis, colouristic analysis, lightfastness and

particle size analysis. Single phase powders were prepared by dry

mechanochemical activation and calcination at the temperature of 1300 °C. An

increasing content of terbium causes a shift of the diffraction lines and an expansion

of the unit cell. Wet-MA method can be suggested as the preferable method for

obtaining the pigment with the most interesting colour properties, optimal particle

size and with the highest light stability.

References

[1] H. Mizoguchi, H.W. Eng, P.M. Woodward, Inorg. Chem. 43 (2004) 1667-1680. [2] A.K. Prodjosantoso, Q. Zhou, B.J. Kennedy, J. Solid State Chem. 200 (2013) 241–245. [3] Ž. Dohnalová, N. Gorodylova, P. Šulcová, M. Vlček, Ceram. Int. 40 (2014) 12637-12645.

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THE REMOVAL OF HEAVY METALS FROM WATER BY NATURAL LIMESTONE AND SECOND CONCRETE

Beatrice Plešingerová, Nicolette Jádi, Agnesa Doráková

Department of Ceramics, Faculty of Metallurgy, Technical University of Košice, Letná 9, Slovakia,


The groundwater contaminated by ions of heavy metals can be purified by the permeable

reactive barriers (PRBs). The limestone and concrete were studied as a promising source of

alkalinity and adopted as the reactive media for the permeable reactive barriers. The results

of batch tests at L/S ratio = 300 ml of contaminated water with Ni2+

, As5+

, Zn2+

, Cu2+

and/or

Fe2+

to 0.5 g a limestone or a concrete of two size particles (80 – 100 m and 355 - 400 m)

showed, that the efficiency of heavy metal ions removal with concrete was better than with

limestone. Neutralization of water initiated the precipitation of metal compounds. Although the

fine particles dissolved faster than coarser particles in the batch test were their efficiencies

comparable. The Fe ions in the water were precipitated in preference and initiated the

removal of As ions.

Keywords: concrete, calcite, neutralization, groundwater, metal ions

Introduction

Devastated lands by mining and industrial activity are the primary sources of

permanent contamination of groundwater. The remediation technologies, as the

permeable reactive barrier (PRB), are developed in order to prevent a spread of the

pollutants by groundwater. The materials in barriers react through the chemical and

physical and/or biological processes including dissolution and precipitation, sorption

and oxidation/reduction 1-5.

The toxicity and reactivity of metals depend on conditions, e.g. pH, Eh, temperature

of water etc. The neutralization and oxido-reduction processes are typical for

removal of heavy metal ions from flow acid groundwater 1. Many different

materials, e.g. red mud, peat, zeolites, iron sorbents and slag, alkaline complexation

agents – lime, calcite, caustic magnesia etc., are using to remove the heavy metals

from water, the mobility of metal ions by converting them to inactive states 5.

The purpose of this study was to compare the activity of limestone and concrete in

model soft acid waters with Ni2+

, As5+

, Zn2+

and Cu2+

metals, their effect on

increasing the pH value and precipitation at conditions of the batch test with liquid-

to-solid ratio: 300 ml /0.5 g limestone/concrete. Tests were carried out with model

solutions with and without the presence of Fe ions.

Experimental part

Crushed limestone (calcite) from the locality of Včeláre (Slovakia) and concrete

panel were used for laboratory batch leaching test. Two particle size fractions: 80 –

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100 µm and 355 400 µm of samples were used for tests. The chemical

composition of samples is shown in Table 1. The concentration of Mex+

in model

solutions which were prepared from sulphates is listed in Table 2.

Table 1: Chemical analysis of limestone and concrete

Sam

ple

Chemical composition wt.% Loss on ignition

wt.%

Particle size µm

Al2O3 SiO2 CaO MgO Fe2O3 K2O Na2O 80 –100 350 –400

Surface area m3/g

Limestone 0.5 0.7 51.7 0.35 0.45 0.02 0.38 42.86 1.2 0.5

Concrete 5.1 42.8 23.3 4.7 3.54 0.54 0.59 21.02 12 – 14 3.5 - 5

Table 2: Concentration of Mex+ in model water prepared from sulphate salts

Ion in solutions Single Cu2+

Mixed Ni2++As5++Zn2++Cu2+ + (Fe2+) Other

observed ions in solution

Ca2+ Al3+ Si4+ Concentration mg.l-1 20 1 + 2 + 10 + 20 + (25)

Initial pH of solution 5.2 2.3 – 3.0

Liquid to solid ratio (L/S) in batch test was 300 ml model solution to 0.5 g concrete or

limestone. Experiments were carried out at laboratory temperature in 500 ml glass

flasks, solutions were stirred at 600 rpm. Sampling in amount of 15 ml was carried in

time 10, 20, 30, 60 and 120 min; at the same time the pH value was measured. The

concentration of Ni2+

, Zn2+

, Cu2+

, Fe2+,

As5+

, Ca2+

, Mg2+

, Al3+

and Si4+

ions was

measured by the inductively coupled plasma-atomic emission spectroscopy (ICP-

AES/ iCAP 6000 Series).

The efficiency of limestone and concrete on remove the Ni2+

, Zn2+

, Cu2+

, Fe2+,

As5+

ions (E%) from model solution was evaluated from concentration changes of Mex+

:

%100.o

o

C

CCE , (1)

oC mg.l-1 - the initial Me

x+ ion

concentration in solution and

C mg.l-1 - the

concentration of Mex+

in time .


The removal efficiency of reactive materials (particle size of 80 100 µm and 355

400 µm) was observed in depending on leaching time and pH. The series of Figs. 1

show simultaneously increase of the pH value in model single solutions with

limestone and concrete; and increase of Ca2+

ion concentration in solution with

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exposure time; and the effect of pH value on precipitation of Cu2+

compounds from

solution. The initial pH values of solutions were about 5.3 – 5.4. Significant

differences were observed between results of lime and concrete while the effect of

particle size of materials was negligible. It could be caused by the fundamentally

higher surface area of the concrete unlike limestone and even their phase

composition (limestone – calcite; concrete limestone – portlandite, calcium/ alumina-

silicates, quartz) 6. Higher removal efficiency of Cu2+

of concrete from solution was

due to higher pH value which was created by concrete, unlike limestone.

a I b

c

Fig.1: The change of pH with time of limestone/ concrete in single solution (Cu2+) – a); the change of Ca2+ concentration on exposure time – b); the Cu2+ removal depending on pH – c).

Fig. 2 shows the simultaneous increase of the pH value and Ca2+

ion concentration in

the mixed solution with exposure time. The initial pH value of solutions was 2.3 and

3.1. After 60 minutes of testing the results confirmed that the concentration of Ca2+

ions and the pH in solution with concrete were comparable or higher than solution

with limestone. Lower Ca2+

ion concentration in the solutions with iron in comparison

to the solutions without iron indicates that the Fe-precipitates inhibited the releasing

of Ca2+

from concrete and limestone into solution. The pH value had even then

gently increasing tendency. After the long-time experiments with limestone, the

solutions were milk coloured because of the gypsum precipitates.

The results in Fig. 3 presented the influence of pH on removal efficiency of Cu2+

and

As5+

from mixed solution. The removal of Cu2+

and As5+

ions increases sharp at

pH 4.5 and at pH 6.5 is 100 %. In the case of these ions was the remove effect

greater when the Fe ions were in the mixed solutions present, too. Once again, the

L18 THE REMOVAL OF HEAVY METALS FROM WATER BY NATURAL LIMESTONE

AND SECOND CONCRETE B.Plešingerová, N. Jádi, A. Doráková


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concrete reacted more perfectly than limestone. The Ni2+

and Zn2+

ions were from

the mixed solutions removed at pH 6.5 with efficiency about 5 – 10 and 15 – 20 %,

respectively.

without Fe with Fe

Fig. 2: (up) The change of pH value with exposure time; (down) time-dependence of concentration of Ca2+ in mixed solutions with and without iron.

Conclusion

At the condition of batch test, concrete (L : S = 300 ml : 0.5 g, 120 min), and

limestone, were able to neutralize the acid water. The increase of pH by concrete

was more significant (pH = 9 - 10) than by limestone (pH = 6 – 7) therefore the

removal efficiency of heavy metal ions from contaminated water with concrete was

higher. This effect was due to the different composition and surface area of

limestone and concrete.

The dissolving rate decrease with the alkalinity of the solution and the particle size of

reactants. For the removal efficiency of Ni2+

, As5+

, Zn2+

, Cu2+

the effect of particle

size was irrelevant. The pH 6.5 was a bottom limit of Cu-hydroxide precipitation

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90

and limit for Ni2+

and Zn2+

higher. The Fe-precipitates inhibited the dissolution of

surface area.

without Fe with Fe

Fig. 3: Influence of pH on removal efficiency of Cu2+; influence of pH on removal efficiency of As5+ from mixed solutions with and without iron.

The systems did not achieve the equilibrium state after 120 min. Long-time column

tests are necessary to appreciate and confirm the effect of concrete and limestone in

permeable reactive barriers.

Acknowledgements This work was supported by the project APVV-0351-12.

References

[1] Golab A.N., Peterson M.A., Indraratna B.: Quarterly Journal of Engineering Geology and Hydrogeology. 39 (2006) 206–223.

[2] Indraratna B., Pathirage P. U., Rowe R.K., Banasiak L.J.:Computers and Geotechnics, 55 (2014) 429–439. [3] Oliva J., De Pablo J., Cortina J.L., Cama J., Ayora C.: Journal of Hazardous Materials. 194 (2011) 312–323. [4] Noiriel,C., Luquot L., Madé B., Raimbault L., Gouze Pb, Van der Lee J.: Chemical Geology. 265 (2009) 160–

170. [5] Hashim M.A., Mukhopadhyay S., Sahu J. N., Sengupa B.: Journal of Environmental Management. 92 (2011)

2355–2388. [6] Allahverdi A., Škvára F.: Ceramics − Silikáty 44 (2000) 114–120.

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CASTABLES IN U. S. STEEL KOŠICE STEEL LADLES

Jana Bounziová, Iveta Nagyová, Martin Černík

Unit of GM USSE Research and Development, U. S. Steel Košice, s.r.o.,

Vstupný areál U. S. Steel, 044 54 Košice, Slovakia


In the U. S. Steel Košice steel ladles (SL), unshaped refractories are implemented

on the bottom and lip ring of ladles. High requirements are placed on these

important parts of ladle, from the perspective of lifetime, wear range, operational

security and economic indicators. The ladle lip ring is lined by high-alumina

castables and the ladle bottom is lined using high alumina castables based on

corundum with the addition of synthetic MA spinel (MgO.Al2O3) or with forming

MA spinel.

Keywords: SL– steel ladle, MA spinel – MgO.Al2O3

Introduction

Unshaped refractory materials are progressive types of products, which often

replace the traditional lining made of piece shaped materials. The advantages of

their use lie in the quick and easy way of forming a monolithic lining without joints in

good resistance to thermal cycling and in the increased thermal insulation capability

compared to the linings made of shaped material [1].

Castables are tested in the ceramics laboratory in terms of material data sheets from

different vendors and pursuant to internal corporate standards. Results of laboratory

tests are a prerequisite for smooth operation, but only the real conditions of the

technological process of secondary metallurgy of steelmaking will verify their

properties in practice.

Lining of SL bottom

The greatest wear of SL lining occurs in the bottom impact area. Wear in this area is

caused by: the impact of steel stream, thermal shock, erosion, and corrosion. The

wear mechanism depends on the process conditions and the quality of refractory

material used.

Process conditions may include [2]:

weight of steel incident per unit of time,

incident stream velocity ,

steel temperature and temperature distribution in the refractory material

before tapping,

grade of steel produced,

steel slag amount left over from the previous tap.

CASTABLES IN U. S. STEEL KOŠICE STEEL LADLES L19

J. Bounziová, I. Nagyová, M. Černík

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MA spinel containing castables with the addition of synthetic MA spinel exhibit good

thermal shock resistance, slag resistance, they are less resistant to corrosion by

molten material. MA spinel forming castables addition of free MgO reacts with

corundum at higher temperatures and creates so called insitu MA spinel. The

reaction takes place in a castable matrix. The spinel formation compacts the

structure of castable and increases the resistance to molten material corrosion

comparing to the castable with MA spinel. Its disadvantage is the low strength

parameters and a large volume changes (about +8 %) associated with the formation

of spinel and linear changes (about +2 %). Selection of the appropriate type of

castable depends upon the specific conditions of secondary metallurgy, which needs

to be verified directly in practice [2, 3].

When choosing a method of lining, the important factor is technical process and

complexity of steel grades produced.

There are several design options for SL bottom lining. The selected design

determines the repair method and the material and time demands [4].

In conditions of U. S. Steel Košice, working lining of SL bottom is made using the

concept of a monolithic spherical bottom made of castable with low content of mixing

water (water demand), followed by casting and vibrating. At the end of 2014, U. S.

Steel Košice tested the concept of inclined bottom to increase yields of steel.

SL lip ring

Closing (edge) ring of SL lining, so

called. „lip ring“ is located just below

the flange of steel shell of SL, Fig. 1.

It performs a safety function, and

eliminates the effect of internal

stresses due to thermo-mechanical

stress of the lining.

After laying the last row of MgOC

bricks in the slag zone, the

monolithic lip ring is formed by

casting into the mould reaching to

the edge of the steel flange. The

castable is under the effect of

thermomechanical stress caused by

temperature deviations.

Fig.1: SL lining diagram [5]

Testing of refractory materials

Besides the testing of thermomechanical parameters at different firing temperatures,

the castables intended for the SL lip ring are tested for thermal cycling resistance

according to the methodology of Research and Technology Center in Munhall, USA.

Testing for thermal cycling resistance is carried out on five blocks (cubes) with size

50 x 50 mm, by heating to 1093 °C/30 min., followed by cooling in water at the

ambient temperature for 5 minutes. After drying the test samples, this process is

Edge ring - lip ring

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93

repeated 10 times, or less, until the disruption of blocks. If the test specimen

remains intact (unaffected) even after 10 cycles, the degree of cracks occurrence is

evaluated by marks 0 4, as is shown in Fig. 2 [5].

Fig.2: Evaluation of resistance to thermal cycling [5]

Fig.3: Corrosion test with steel [5] Fig.4: Corrosion test with slag [5]

Castables intended for SL bottom are tested by static pot corrosion tests

with operational slags at temperature 1600 °C/5 h.

Department of Metallography and Failure Analyses made Xray diffraction phase

analysis on samples of castables. As an example, phase analysis of MA spinel

forming castable in dry state (sample 1) and after firing at 1600 °C/5 h (sample 2)

CASTABLES IN U. S. STEEL KOŠICE STEEL LADLES L19

J. Bounziová, I. Nagyová, M. Černík

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94

can be mentioned. Sample firing formed about 40 % spinel, and also about 8.8 %

hibonite, which originates in the corundum impurities diaoyudaoite (Table 1).

Table 1: X-ray diffraction phase analysis of castable [5]

Identified phase composition Sample 1 Sample 2

Chemical formula Mineralogical name Space group Content [wt%]

Content [wt%]

Al2O3 Corundum (167) R-3c 87.3 47.1

MgAl2O4 Spinel (227) Fd-3m - 39.5

NaAl11O17 Diaoyudaoite (194) P63mc 3.6 4.5

(Ca,Mg)Al12O19 Hibonite (194) P63mc - 8.8

MgO Periclas (225) F m-3m 9.1 -

Fig.5 Detail of more precise X-ray diffraction record of sample 2 [5]

Conclusion

In conclusion, it can be stated that the technology castable processing and

production of monolithic parts of SL lip ring and bottom is mastered and widely used.

Selection of the appropriate type of castable, castable processing, drying process

and, last but not least, technological impact of secondary metallurgy have a decisive

impact on the life and economic evaluation of the cost of steel ladle.

References

[1] Staroň,J., Tomšů,F., Žiaruvzdorné materiály výroba, vlastnosti a použitie, Slovmag, a.s. Lubeník, SMZ, a.s. Jelšava, Keramika, a.s. Košice, (2000)

[2] Hwang,K.H., K.D. Oh and R.C. Bradt: In Situ Spinel Bound Formation (Expansion/Contraction) During Firing, 1575-1580 in v. III. of Proc. UNITECR 97, New Orleans, USA (1997)

[3] Wöhrmayer,C., Parr,C., Chassing,P.: The Role of High Alumina Cement in Spinel Forming and Spinel Containing Castables. 13th International Conference on Refractories, Prague, 28-29 March, 1999

[4] Tatič,M. et al.: Optimization of the ladle bottom lining. In: 16th International Conference on Refractories Proceedings: Czech Silicate Society, 14.-15.5. 2008, Praha, s. 141- 145 (2008). ISBN 978-80-08-02021-9

[5] Technická dokumentácia spoločnosti U. S. Steel Košice, s.r.o., 2000-2015


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LECTURES

– industrial and education activities –

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96

Technical University of Kosice

FACULTY OF METALLURGY

Address: Letná 9, 042 00 Košice,

Slovak Republic Phone: ++421 55 6022023

Fax: ++421 55 6337048 E-mail: [email protected]

www.tuke.sk/hf

Research

Frontier and applied research focused on:

pig iron and steel production

ladle metallurgy of steel and continuous casting

metallurgy of non-ferrous metals

modern foundry technologies

metal forming

metal finishing and protection

refractories - production and application

energy balances of the thermal processes and aggregates

energy resources efficient utilisation

environmental aspects of the metallurgical production

waste recycling and utilisation

research and development of new materials and technologies

development of new methods for the material properties evaluation

physical and mathematical modelling of technological processes

quality control and environment protection in the industrial enterprises

Staff of the Faculty is involved in solving 26 international projects and 39 projects funded by the SR government.

The Faculty publishes a specialised journal Acta Metallurgica Slovaca, which is distributed to 25 countries.

The most significant industrial partners of the Faculty of Metallurgy are:

U. S. Steel Košice (Slovakia)

Silicon, a.s. (Slovakia)

Železiarne Podbrezová, a.s.(Slovakia)

ArcelorMittal Ostrava, s.r.o. (Czech Republic)

Třinecké železárny, a.s. Třinec (Czech Republic)

Outokumpu Research, OY (Finland)

SE (ENEL)- Elektrárne Vojany (Slovakia)

Slovalco, a.s. (Slovakia)

Slovenský plynárenský priemysel (Slovakia)

Slovenské magnezitové závody (Slovakia)

Education

The Faculty of Metallurgy offers the three-level

university education in the following fields:

Metallurgy

Foundry

Metal Forming

Materials Engineering

Ceramics / Refractories

Energetic Engineering

Gas Technology

Management Systems

Environmental Engineering

Waste Recycling

All forms of life-long education (post-diploma and

requalification courses, expert seminars and

training).


http://www.tuke.sk/hf

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APPLICATION OF CASTABLES PRODUCED BY RMS, a.s. KOŠICE

IN THE PUSHER FURNACE

Rastislav Hirjak1, Jana Bounziová

2, Dáša Chudíková

1, Jaroslav Parnahaj

1,

Michal Mingyár2

1 RMS, a.s. Košice, Vstupný areál U. S. Steel, 044 54 Košice, Slovakia 2 U. S. Steel Košice, s.r.o., Vstupný areál U. S. Steel, 044 54 Košice, Slovakia


Considering the energy-intensive production of burnt refractory materials and implementation

of new monolithic lining manufacturing technologies, the share of the unshaped materials

production is growing at the expense of building materials. The first monolithic hearth of low-

cement castable KOFOND LC 18 was applied in pusher furnace #3 at U. S. Steel Košice,

s.r.o. in 2010. This pusher furnace hearth working lining reached a record amount of pushed

slabs. The RMS, a.s. Košice Company’s production program includes a wide range of

insulating and dense castables based on various raw materials with different cement contents

(MCC, LCC, ULCC, and NCC castables). The paper deals with practical applications of

castables produced at RMS, a.s. Košice for pusher furnaces.

Keywords: castable, pusher furnace, lining

Introduction

One of main steel industry trends is the use of high quality refractory materials

resistant to the external corrosive environment. There are ever-increasing quality

requirements placed on refractory materials in accordance with economic impact.

First of all it is the thermal unit's service life increasing with respect to reducing the

refractory material consumption, increasing productivity, saving material, saving

energies, etc. One possible way to meet such requirements is the use of monolithic

linings. Since 2010 there has been a change in the pusher furnace (PF) hearth

working lining at U. S. Steel Košice, s.r.o. Original melted corundum blocks were

replaced with monolithic working linings by applying high-alumina castables. The

first pusher furnace hearth monolithic lining of low-cement castable KOFOND LC 18

was implemented in 2010. As of today, its service life has experienced a record

amount of pushed slabs. Second monolithic lining of this castable is applied in

pusher furnace #4. The purpose of this paper is to present a number of types of

refractory materials produced by RMS, a.s. and used for pusher furnaces.

PF hearth lining

At U. S. Steel Košice, s.r.o. pusher furnaces are an important production node in the

metallurgical process. Their performance and slab heating quality have a major

impact on the quality of rolled material. The pusher furnace hearth lining is one of

the factors affecting the pusher furnace continuous operation and the finished

product quality.

APPLICATION OF CASTABLES PRODUCED BY RMS, a.s. KOŠICE IN THE PUSHER FURNACE L20 R. Hirjak, J. Bounziová, D. Chudíková, J. Parnahaj, M. Mingyár

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98

The permanent hearth shell lining consists of a number of layers of insulating

and refractory materials. The following high-level requirements are placed on the

hearth working lining in demanding high-temperature conditions:

- No impact on thermal field and surface quality of the heated charge

- Resistance to high temperatures and sudden changes in temperature

- Minimum linear and volume changes

- Resistance to chemical action of scale and furnace atmosphere

- Abrasion resistance and resistance to mechanical shocks

- Sufficient service life at a minimum maintenance

Original hearth working linings consisted of melted corundum blocks (Fig. 1), which

were characterized by high strength, abrasion resistance, and low porosity (2 – 8

%). Their disadvantage was high thermal conductivity (5 – 7 W.m.K-1

at 1300 oC)

which resulted in increasing heat transfer from the furnace hearth to the environment

(heat loss) and deterioration of soaking in the slab profile. Due to dilatation these

blocks had lesser stability resulting in displacement of individual rows of blocks

mainly along the length of the furnace. Due to frequent temperature changes in the

pusher furnace the lining working layer was gradually worn out (Fig. 2), thermal field

homogeneity deteriorated and increased scale accumulation took place on the surface.

Fig. 1: New corundum blocks Fig.2: Worn corundum blocks

Wear of blocks and formation of scale buildups gave rise to more frequent repairs

and thus extended downtimes of the pusher furnace. The hearth repairs were

carried out mechanically or using a strong stream of water. Furnace temperature

changes generated stress in the corundum blocks due to their thermal expansion.

Generated stress disturbed the structure of the hearth lining surface (Fig. 3). The

disturbed surface of corundum blocks caused mechanical damage, grooves on the

bottom parts of steel slabs, thus increasing nonconforming production. [1]

New technical solution of pusher furnace #3 hearth lining was implemented in May

2010 by applying LCC castable KOFOND LC 18. United States Steel Corporation

has developed standards for evaluation of refractory materials designed for the

pusher furnace hearth. In order to evaluate the castable suitability, Research and

Development U. S. Steel Europe carried out a series of tests regarding the

resistance to sudden changes in temperature, thermal expansion, thermogravimetry,

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differential thermal analysis and corrosion tests with different steel types. Physical

and chemical properties of castable KOFOND LC 18 are shown in Table 1.

Table 1: Technical parameters of KOFOND LC 18

BD

110°C (kg.m-3)

BD

1650°C (kg.m-3)

CCS

110°C (MPa)

CCS

1650°C (MPa)

CMOR

110°C (MPa)

CMOR

1650°C

(MPa)

PLC

1650°C

(%)

Water Demand

(%)

Content of Al2O3

(%)

Content of CaO

(%)

typ.

3,110

typ.

3,070

min.

70

min.

90

typ.

7

typ.

25

typ.

- 0.1 4.5 - 5

min.

94

typ.

1.8

Fig. 3: Hearth of corundum blocks after scale removal

Fig. 4: Even scale surface of the monolithic hearth

Castable KOFOND LC 18 based on corundum and additives with hydraulic bond is

intended for casting and vibration processing. Maximum operating temperature is

1750 °C. It is characterized by high resistance to sudden changes in temperature,

resistance to action of molten metal, slag, and by high thermo-mechanical

properties. The castable has an excellent ability to resist to scale buildup formation -

and thus it ensures even surface of the hearth (Fig. 4) and improvement of the

heated slab quality. Due to a lower thermal conductivity (Fig. 5) of the castable

compared to corundum blocks, heat loss through the furnace shell are reduced to

3 W/mK at 1200 °C. The reduction of temperature difference between the top and

bottom surface of slabs improves the heated charge thermal profile homogeneity.

The advantage of the application of the unshaped refractory materials consists in a

rapid and simple method of the hearth monolithic lining manufacturing minimizing

the gaps. Reduced physical load during the application and repairs reduces the risk

of injuries of employees. Reduced work volume reduces costs of refractory material

and refractory work. The castable thermal expansion (1.046 % at the temperature of

1200 °C – Fig. 6) that depends on the dilatation of individual phases, characterizes

reversible dimensional changes in the material. It is important as regards the design

of the size, number, and thickness of individual panels when designing the FP

monolithic hearth, mainly in terms of the dilatation joints sizing. [2]


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100

Fig. 5: Thermal conductivity of KOFOND LC 18 Fig. 6: Thermal expansion of KOFOND LC 18

PF #3 hearth medium repair

In the conditions of U. S. Steel Košice, s.r.o. the hearth working lining is applied in

situ by casting castables into panels. The size of the whole hearth surface without

the surface of extractor slots is ca 120 m2.

One layer of hard fire clay bricks in the thickness of 65 mm (KOSAM TB

manufactured by RMS, a.s. Košice) is stacked in panels, directly on even permanent

hearth lining (Fig. 8). Prepared castable is installed in a checkerboard manner in

individual 190 mm thick panels (Fig. 9). After the castable hardening the mold is

removed and dilatation joints are filled with mineral water. The drying control and

service temperature achievement is carried out by thermocouples located directly in

the castable bottom layer above the permanent lining and on the castable surface.

Fig. 8: Substrate fire clay layer in the mold of panels

Fig. 9: Panel casting finalization

The planned repair of Pusher Furnace #3 was carried out in January 2015.

Monolithic hearth of Castable KOFOND LC 18 reached a record service life of more

than 4.7 million tons of pushed slabs. After the PF inspection a medium hearth

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repair was carried out with the repair of the leading edge in the length of ca 1.5 m

and replacement of extractor slots. The main part of the hearth was without

significant damage. The castable thickness in the middle of the hearth is 160mm,

compared to original 185 mm. The castable residual thicknesses in the middle of the

leading edge were 210 mm compared to original 250 mm (Fig. 10). A minimum scale

penetration in the castable was observed in the hearth lining cross-section. The

buildup surface layer varies in the range of maximum 10 – 20 mm on the leading

edge (Fig. 11). At the exit from the NP (extractor slots) the buildup is growing and

scale reached the level of 60 mm. What is important is a compact and even scale

surface layer with no need for deburring and special maintenance. A more

significant hearth working lining wear was observed at the furnace edges. Lateral

edges of slabs created by mechanical action a groove deep ca 140mm from the

original lining, in the distance of ca 300 mm from the wall (Fig. 12, 13). The medium

repair of the PF hearth is carried out approximately once every two years.

Fig. 10: Castable residual thickness Fig. 11: Scale buildup thickness

Fig. 12: Mechanical wear of the hearth Fig. 13: Wear by the slab edge 140 mm

Unshaped materials for PFs

In addition to the castable for the PF hearth RMS, a.s. Košice also supplies other

castables which help to achieve continuous operation of the pusher furnace.


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Gunning Castable KOTOR TO7A-1 or KOTOR TO7A-3 is used for hot as well as

cold repairs of the pusher furnace walls. These are used for hot spraying of extractor

slots. These materials are used in local repairs of flue pipes during the PF shutdowns.

At U. S. Steel Košice, s.r.o. the PF exit door refractory lining was applied using heat-

insulating castables KOPORO 011, KOPORO 012, KOPORO 0165. It was necessary

for the door service life to harmonize strength characteristics of the insulating castable

with operating conditions where it is most exposed to thermal shocks, vibration, and

mechanical shocks. [2] Insulating castables are used for local repairs of the PF roof.

The monolithic prefabricated shapes produced of high-alumina Castable

KOFOND F6A-1 are used for the skid support heat protection in pusher furnaces at

U. S. Steel Košice, s.r.o. as well as at other metallurgical companies.

The KOFOND FA 5 Castable monolithic blocks were designed for the leading edge

permanent lining supporting structure under the monolithic hearth.

The charge heating in pusher furnaces is ensured by the combined heating of roof

radiating burners and bottom head turbulent burners. Currently the mixture gas

heating medium energy-saving project is being undertaken at pusher furnaces. The

sloped wall design was implemented in the bottom preheating zone which is

intended to direct flue outgoing to the flue pipe as close as possible to the bottom

part of slabs and this way use the secondary heat for the charge heating. Project

design, refractory work, and supplied refractory materials (KOSAM TB, KOFOND F4S,

KOLAS 130, and KOLAS 50) were carried out in cooperation with RMS, a.s. Košice.

Summary

Implementation of the new technical solution related to the hearth lining repair at

pusher furnaces of U. S. Steel Košice, s.r.o. is an effective castable monolithic lining

application method. The use of this method is accompanied by material, energy and

work savings. Obtained results of the PF monolithic lining indicate that these have

their application and are beneficial both in operational and economic terms.

RMS, a.s. Košice appreciates its participation in the implementation of the project,

the aim of which is to reach the limit of 5 million tons of pushed slabs in PF #3. This

year the Company celebrates the 50th anniversary of its unshaped and shaped

refractory products manufacturing tradition. The Company constantly innovates its

products and seeks to offer new solutions to its consumers. The Company is not only

the producer of castables but also a complete refractory work performer. It offers to

its customers a complete service from the lining draft design to the author and

service supervision during bricklaying and during the metallurgical unit operation.

References

[1] J. Bounziová, M. Mingyár: Pusher furnace hearth monolithic linings in the conditions of U. S. Steel Košice, s.r.o., Proceedings of the International Scientific Conference Refractory Materials, Furnaces and Thermal Insulations, April 8-10, 2014.

[2] D. Chudíková, J. Parnahaj, V. Chudý, J. Bounziová: Possibilities of using castables manufactured by Refrako in the metallurgical or other industrial sectors, Proceedings of the International Scientific Conference Refractory Materials, Furnaces and Thermal Insulations, April 11-19, 2012.

[3] M. Tatič: Industrial furnaces and thermal units, Textbooks for postgraduate study Refractory Materials, 2006.


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THE IMPACT OF INDUSTRY ON CLIMATE CHANGES FROM HOLISTIC

ENVIRONMENTALISM POINT OF VIEW

Damir Hršak, Anita Štrkalj, Ljerka Slokar, Zoran Glavaš

University of Zagreb Faculty of Metallurgy, Sisak, Croatia


Holistic environmentalism is a holistic consideration of the environment in its material and

spiritual nature. Essence of the environmentalism goes far beyond what we constrainedly call

environmental protection. Many indicators of environment crisis which are clear visible in

various forms worldwide warn that current development of our civilization is not sustainable.

We must be ready for significant changes in our way of living. Use of technologies which do

not minimise sufficiently the negative impact of manufacturing process on the environment is

no more acceptable if humankind wants to survive.

Key words: climate changes, industry, holistic environmentalism

Introduction

There is no longer any doubt that climate change is due to the behaviour of

mankind. The owners of profits in industrial production generally did not care for the

consequences for the environment. The main culprits, on world level, for climate

change consistently for many year intensely pressured, sometime blackmailed,

conscientious scientists and politicians who recognized the threat. It was not easy to

be belittled and mocked. Now we have a significant problem.

Many indicators of crisis of the environment which are visible in various forms

worldwide clearly warn that current development of our civilization is not sustainable.

We must be ready for significant changes. Such generally negative trend of social

impact on the environment follows from the growth of Earth population which in

recent years experienced manifold increase and continues at such a trend;

"consumeristic" civilisation, which under the influence of marketing and the overall

economic system, experiences and covets the standard and quality of life primarily

in the form of possession and ever more intense consumption of new material

assets; the use of technologies which do not minimise sufficiently the negative

impact of manufacturing process on the environment, all that resulting from

endeavour that product price be formed so as to make uncompetitive on the market,

and thus economically unsustainable, those manufacturing methods which are

sufficiently environment-friendly.

One of the positive outcomes of escalation of mentioned crisis is the fact that it is all

the more difficult to ignore various symptoms of negative human impact on the

environment. Consequently, more and more people are becoming aware that

healthy environment is not something that is granted regardless of their activities,

and they are prepared to change the environment-unfriendly habits, to invest the

L21 THE IMPACT OF INDUSTRY ON CLIMATE CHANGES

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effort and support the activities for conservation and sustainable use of the

environment. [1]

Climate changes

Periodical changes of Earth climate brought about as a result of changes in

the atmosphere and interactions between the atmosphere and other geologic,

minerological, chemical, biological, and geographic factors within the Earth system.

Our atmosphere is a dynamic fluid which is continually in locomotion. Atmosphere

physical properties and its velocity and direction of motion are on influence

of atmospheric chemistry, position of continents, ocean currents, solar radiation as

well as vegetation growing on the land surface. These impacts are changeable

through time. Climate, which results from the physical properties and locomotion of

the atmosphere, varies at every achievable time scale.

For example, ocean water is cool but the chemical reaction is exothermic and

generates its own heat. In the past it was thought serpentinite was quite rare, and

occurred only above subduction zones, or within small ultramafic intrusions. Now it

is known that serpentinites pretty much underlay the entire ocean floor, forming part

of the ophiolite sequence, as olivine-rich magmas oozing up through the mid-ocean

rift are serpentinized by ocean water.

Impact of Industry

Radioactivity, petrochemical, and synthetic organic chemicals were largely

developed and surfaced in the environment in the middle of 20th century. During this

period major environmental problems appeared with rapid and serious

consequences.

Admit of Earth climate change as an environmental issue has drawn attention to the

climatic impact of mankind actions. Most of this attention has focused on carbon

dioxide emission via fossil-fuel combustion and deforestation. Mankind actions yield

releases of other greenhouse gases, such as and chlorofluorocarbons. There is no

doubt among serious climatologists that greenhouse gases influenced the radiation

budget of Earth. The nature and magnitude of the climatic response are a subject of

intense research work.

Paleoclimate records from tree coral, and ice cores indicate a clear warming trend

spanning the entire 20th century and the first decade of the 21st century. We knows

that the 20th century was the warmest of the past 10 centuries, and the first decade

21th century was the warmest decade since the beginning of modern instrumental

record keeping.

It is becoming total clear that human influence on vegetation cover can have global

effects on climate, due to changes in the sensible and latent heat flux to the

atmosphere and the distribution of energy within the climate system. The extent to

which these factors contribute to recent and ongoing climate change is an emerging

area of research work.

In figure 1 there is projections of surface temperature changes until 2099.

http://www.britannica.com/EBchecked/topic/175962/Earth

http://www.britannica.com/EBchecked/topic/121560/climate

http://www.britannica.com/EBchecked/topic/41364/atmosphere

http://www.britannica.com/EBchecked/topic/41364/atmosphere

http://www.britannica.com/EBchecked/topic/175962/Earth

http://www.britannica.com/EBchecked/topic/211221/fluid

http://www.britannica.com/EBchecked/topic/108987/chemistry

http://www.britannica.com/EBchecked/topic/424354/ocean-current

http://www.britannica.com/EBchecked/topic/552984/solar-radiation

http://www.britannica.com/EBchecked/topic/596034/time

http://www.britannica.com/EBchecked/topic/121560/climate

http://www.britannica.com/EBchecked/topic/155854/deforestation

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Holistic Environmentalism

Holistic environmentalism is a holistic consideration of the material and spiritual

environment. Essence of the environmentalism goes far beyond what we

constrainedly call environmental protection. To understand the nature means do not

destroying life, do not lower quality of air, water and soil, do not make any

irreversible destruction, do not perform what we are instantly able to perform.

Fig. 1: Projections of Surface Temperature Changes [2].

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The ones who define ecology as removal of eligible amount of industrial leftovers

exhibit an alarming misunderstanding and dangerous primitivism. Responsible

intellectuals are obliged to confront such an opinion from their viewpoints

respectively, taking the stand of holistic environmentalism as well. It does not matter

which viewpoint we present – city, country or the whole world as it is necessary to

act and confront the destruction of dignified future. It is ethical duty of each person to

take care of everybody and act on everything around him.

Conclusion

The holistic environmentalism point of view is not against of industrial production

and technical development of modern society. On the contrary we want developed

but solidary and justly society with environmental protection for our and future

generation.

But there is urgent requirement of contemporary times that a mankind as reasonable

actor in nature should find the way of wellbeing for himself and his immediate

environment as a part of world ecosystem. As long as we use the environment,

mankind is forced to recognize the fact that the environment is not the object of his

highhandedness.

Global warming affecting the whole mankind and all living creatures on the planet is

the consequence of sick greed for wealth conforming to the egoistic goals of a small

number of powerful people who are mainly high but wrong educated with no respect

for Earth as whole.

Holistic environmentalism point of view is always reasonable and conscientious, too.

It is mankind future friendly point of view.

References

[1] I. Gudelj, L. Runko Luttenberger, A. Senta Marić, M. Šiljeg, The Need for Integrated Approach to Environmental Protection in the Republic of Croatia, The Holistic Approach to Environment, 3 (2013)1, 41-51.

[2] Forth Assessment Report of Intergovernmental Panel on Climate Change

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COMMENTS ON THE CURRENT RELATIONSHIP BETWEEN GEOLOGY

AND CHEMISTRY AT SLOVAK UNIVERSITIES

Karol Jesenák

Faculty of Natural Sciences, Comenius University, Mlynská dolina, Ilkovičova 6,

842 15 Bratislava, Slovak Republic.


Keywords: education, geology, ceramics

Introduction

A growing divergence of geology and inorganic chemistry has been the main trend

during the last two decades in Slovakia. The main reason is the significant decrease

in the exploitation of primary inorganic materials in Slovakia. The reduction in the

total number of workers in mining industry has had a tremendous impact on the

number of students interested in the study of geology. The overall trend in relations

between geology and chemistry was buffered until recently by handling of the

environmental problems associated with termination of the ore mining at the end of

the last century.

Unlike in the past, when the major area of cooperation between geologists and

chemists in Slovakia was a mining and processing industry, currently it is almost

exclusively basic research. Universal method how to achieve success in research is

to focus on a very narrow scientific problem. This approach usually results in mutual

separation of these different disciplines.

Impact on Education

The long-term trend in distancing geology and inorganic chemistry is probably

already irreversible. A serious problem, however, has negative impact on science

and technical education. This is particularly true in case of general chemical

education. A significant reduction of geological education in most primary and high

schools has resulted in unacceptable state for students not knowing the most

important natural inorganic substances. For example, students often deal with

various modern or trendy issues at conferences, although they have no idea of the

extraction and use of the limestone. University students often claim that iron is

produced from iron oxide and glass is made of silicon. George Bernard Shaw's put it

well: "No man can be a pure specialist without being in the strict sense an idiot."

The question is who is responsible for this state.

Conclusion

This paper highlights the most important changes regarding the relationship

between geology and chemistry over the last 20 years. It also points out the

consequences of these changes for the general chemical education. Improvement of

the current situation could be achieved if chemistry teachers would be more

L22 COMMENTS ON THE CURRENT RELATIONSHIP BETWEEN GEOLOGY AND CHEMISTRY

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interested in some geological topics. One possibility how to solve this problem is

free use of two publications listed at the end of the paper [1, 2].

Fig. 1: This image can be found in publication [2] in the part "The history of ceramic production in Slovakia"

Fig. 2: This image can be found in publication [2] in the part "The history of ceramic production in Slovakia"

Fig. 3: This image can be found in publication [1]. Pyrargyrite in calcite (Author: D. Ozdín)

Acknowledgements

The work has been supported by grant

VEGA 1/0276/15 of the Ministry

of Education of the Slovak Republic.

References

[1] Jesenák, K.: Exkurzia po miestach ťažby a spracovania rudných surovín na Slovensku. 1. vyd., Bratislava: Univerzita Komenského, 2011, 957 strán, ISBN 978-80-223-3127-2

[2] Jesenák, K.: Exkurzia po miestach ťažby a spracovania anorganických nerudných surovín na Slovensku. 1. vyd., Bratislava: Univerzita Komenského, 2011, 948 strán, ISBN 978-80-223-3128-9


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POSTERS – ABSTRACTS

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PREPARATION AND CHARACTERIZATION BIOCERAMICS PREPARED

FROM TETRACALCIUM PHOSPHATE-NANOMONETITE CEMENT

Radoslava Štulajterová, Ľubomír Medvecký, Mária Giretová, Tibor Sopčák

Institute of Materials Research of SAS, Watsonova 47, 040 01 Košice, Slovakia


Keywords: calcium phosphate cement, hydroxyapatite, bioceramics, in-vitro tests

The calcium phosphate ceramics (CPCs) have been widely used for bone tissue

repair and augmentation. The proportion of HAP and TCP in the biphasic calcium

phosphate mixtures after decomposition of the initial CDHA may also greatly

influence sintering and the final microstructure of ceramics [1]. The crystallite size of

HA may play an important role in governing the expression of osteoblast activities

[2]. The possibilities of the preparation of porous calcium phosphate ceramic

(hydroxyapatite ceramic) from the tetracalcium phosphate-nanomonetite biocement

were studied. This preparation method is advantageous from the point of formation

of the biphasic calcium phosphate ceramics composed of β- or α-TCP + HA. The

powder cement mixture was synthesized by in situ reaction between TTCP and

diluted solution of the orthophosphoric acid in ethanol (1:4). Cement ceramics was

prepared by sintering of hardened cement samples at 1050, 1150 and 1300 °C for

1 h. The hardened cement was composed of calcium deficient nanocrystalline

hydroxyapatite, which was partially decomposed to crystalline hydroxyapatite and

β- TCP after sintering at 1050 °C and the morphology of origin needle-like HAP

particles was fully changed and hydroxyapatite particles obtained the spherical

shape. No changes in the phase composition and content of β- TCP were observed

after heating at temperature of 1150 °C what is in accordance with the CaO-P2O5

phase diagram. In the case of conventional HA ceramics, both the β- TCP and

α- TCP minor phases were found in HA ceramics after sintering at 1150 °C. After

annealing at 1300 °C, irregularly shaped and spherical micropores were observed.

The influence of sintering temperature on the compressive strength (CS) and the

phase composition of calcium phosphate ceramics after sintering process are shown

in Table 1.

Table 1: Sintering temperatures, phase composition, compressive strength and porosity of substrates.

Sintering temperature [°C]

XRD analysis Compressive strength

[MPa]

cement Nano-HAP 24±3.79

1050 HAP, β-TCP 20±2.08

1150 HAP, β-TCP 22±4.36

1300 HAP, β-TCP 36±1.53

Ceramics 1150 HAP, α,β-TCP 340±50


PREPARATION AND CHARACTERIZATION BIOCERAMICS PREPARED FROM TETRACALCIUM PHOSPHATE-NANOMONETITE CEMENT

R. Štulajterová, Ľ. Medvecký, M. Giretová, T. Sopčák


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From the comparison of the proliferation activity of osteoblasts cultured on cement

ceramics and conventional ceramics resulted that the growth of osteoblast

population was about 30 % higher on cement ceramic substrate (Fig.1).

Fig. 1: Osteoblast proliferation on substrates measured by conversion of Alamar blue relatively to negative control (wells with cells cultured for 8 days).

Biphasic porous calcium phosphate ceramics with hydroxyapatite as the major and

α- or β-TCP as the minor components was prepared by sintering of transformed

calcium phosphate cement. The compressive strength of bioceramics sintered up to

1150 °C was reduced in comparison with hardened cement. The cement ceramics

had improved proliferation activity in comparison with conventional ceramics.

Acknowledgements

This work was supported within the framework of the project „Advanced implants seeded with

stem cells for hard tissues regeneration and reconstruction“, which is supported by the

Operational Program “Research and Development” financed through the European Regional

Development Fund.

References

[1] S. Raynaud, E. Champion, D. Bernache-Assolant, Calcium phosphate apatites with variable Ca/P atomic ratio II. Calcination and sintering. Biomaterials. 23 (2002) 1073-80.

[2] JL Ong, CA Hoppe, HL Cardenas, R Cavin, DL Carnes, A Sogal, GN Raikar. Osteoblast precursor cell activity on HA surfaces of different treatments. J Biomed Mat Res Part A. 39 (1998) 176-83.

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INFLUENCE OF THE pH VALUE ON PHASE COMPOSITION

AND MORPHOLOGY OF CaO–SiO2–P2O5 BIOACTIVE GLASSES

SYNTHESIZED BY SOL – GEL PRECIPITATION

Tibor Sopčák, Ľubomír Medvecký, Radoslava Štulajterová, Juraj Ďurišin

Institute of Materials Research of SAS, Watsonova 47, 040 01 Košice, Slovakia


Keywords: sol-gel precipitation, bioactive glass, hydroxyapatite

For decades, the sol-gel processes have been widely used in various fields of

materials science especially in the production of ceramics and glasses for

biomedical application. This reaction process can be influenced by several factors

including hydrolysis and condensation rate of the precursors, method of mixing, pH,

nature and concentration of reactants, aging time etc. One of the most important

factors that may affect the material phase composition, particle size and morphology

is the solution pH value. Herein, different pH levels were maintained during the

preparation of ternary CaO-SiO2-P2O5 (CaSiP) glass systems in order to study this

parameter influence on final materials properties.

Two CaSiP systems with starting composition 50.0 wt.% CaO – 36.7 wt.% SiO2 –

13.3 wt.% P2O5 were synthesized by means of sol-gel precipitation using Na2SiO3,

H3PO4 and Ca(NO3)2 solutions. In both systems, the initial pH of the Na2SiO3 and

H3PO4 mixtures were first adjusted to (pH = 10.0) by HNO3 (1+1) addition. In CaSiP

1, the pH of the reaction mixture was kept at 10.0 during Ca(NO3)2 solution addition.

However, in the case of CaSiP 2, the reaction pH was left uncontrolled and

decreased to final (pH = 7.4) after the slightly acidic Ca2+

solution addition. The

resulting precipitates were additionally stirred for 1 hour, washed with de-ionized

water (500 ml), ethanol, filtered over the membrane filter (Millipore, 0.2 µm pore

size) and dried at 120 °C for 3 h. As shown in Fig.1, the CaSiP 1 pattern revealed

amorphous character, while in the case of CaSiP 2, obtained at neutral final pH,

both the formation of nanocrystalline hydroxyapatite (JCPDS 24-0033) and a new

amorphous phase (amorph II) can be also observed. From the TEM analysis (Fig. 2)

clearly resulted that when the pH was kept at 10.0, during the precipitation process,

formation of agglomerates composed of 10-50 nm spherical individual particles were

found (Fig 2a, CaSiP 1). The selected area electron diffraction (SAED) revealed

amorphous character of powder agglomerates. On the other hand, as the solution

pH decreased to final 7.4 value (Fig. 2b, CaSiP 2), agglomerates of finer 5 25 nm

globular particles were observed. The results of SAED verified higher crystallinity

and presence of the nanocrystalline apatite phase in the system with lower solution

pH value.

INFLUENCE OF THE pH VALUE ON PHASE COMPOSITION AND MORPHOLOGY OF CaO-SiO2-P2O5 BIOACTIVE GLASSES SYNTHESIZED BY SOL – GEL PRECIPITATION

T.Sopčák, Ľ.Medvecký, R.Štulajterová, J Ďurišin


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Fig. 1: XRD patterns of the CaSiP 1 (pH = 10.0 cont) and CaSiP 2 (pH = 10.0 7.4) samples

Fig. 2a: TEM record of CaSiP 1 Fig. 2b: TEM record of CaSiP 2 (pH = 10.0 cont) sample (pH = 10.0 - 7.4) sample

Acknowledgements

This work was supported by the Slovak Grant Agency of the Ministry of Education of the

Slovak Republic and the Slovak Academy of Sciences, Project No. 2/0047/14.

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NEW TYPES OF TiSi ALLOYS WITH TITANIA AND SILVER SOL-GEL

COATINGS

Diana Horkavcová1, Pavel Novák

2, Martin Černý

1, Iva Fialová

1,

Eva Jablonská3, Zuzana Zlámalová Cílová

1, Aleš Helebrant

1

1Department of Glass and Ceramics, Faculty of Chemical Technology, University of Chemistry and

Technology Prague, Technická 5, 166 28 Prague, Czech Republic 2Department of Metals and Corrosion Engineering, Faculty of Chemical Technology, University of

Chemistry and Technology Prague, Technická 5, 166 28 Prague, Czech Republic 3Department of Biochemistry and Microbiology, Faculty of Food and Biochemical Technology,

University of Chemistry and Technology Prague, Technická 3, 166 28 Prague, Czech Republic


Keywords: titanium alloy, sol-gel coatings, silver, dip-coating technique, antibacterial properties

Titanium and its alloys are presently used as biomaterials for orthopaedic and dental

implantology. They have excellent mechanical properties and exhibit good

biocompatibility. Today´s research is focused on preparation of a porous material

that would ensure the interconnection with bone tissue [1-3] and also on modification

of the surface of the material to be bioactive. Another important property of

biomaterials is resistance to microorganisms causing postoperative complications.

One option to achieve the bioactivity or antibacterial effect of implants is the

preparation of functional coatings [4]. At present there are many methods for coating

and each has its advantages and disadvantages. Sol-gel dip-coating method is

technically simple, cheap and allows coating complex shaped substrates. The

principle consists in the preparation of sol (which contains the functional component)

into which the substrate is dipped and after subsequent with drawing fired.

Conditions for the production of the sol, coating and firing are specific to a particular

type of substrate and coating.

The aim of the study was to characterize both newly developed porous (prepared by

sintering) and non-porous (prepared by melting) TiSi5, TiSi10 titanium alloys and

create the titania sol-gel coatings with the nitrate and phosphate silver on their

surfaces by dip-coating technique. Coating process was carried out under

continuous stirring. The coated porous and non-porous substrates were fired at

400 °C for 2 hours at air. The coatings on the porous titanium alloys TiSi5 and

TiSi10 partly cracked (Fig. 1, Fig. 2) and the silver nanoparticles were distributed

homogeneously in the coating after firing. The silver particles were predominantly

placed in cavities of the porous titanium alloys. The cracks in the coatings on the

non-porous titanium alloys TiSi5 and TiSi10 (Fig. 3, Fig. 4) spread from the silver

particles Ag3PO4. The coatings with silver in form of AgNO3 had the minimum

cracks.


NEW TYPES OF TiSi ALLOYS WITH TITANIA AND SILVER SOL-GEL COATINGS D. Horkavcová, P. Novák, M. Čer1, I. Fialová, E. Jablonská, Z. Zlámalová Cílová, A. Helebrant


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Fig. 1: Titania coating with Ag on porous TiSi5 Fig. 2: Titania coating with Ag on porous TiSi10

Fig. 3: Titania coating with Ag on non-porous

TiSi5

Fig. 4: Titania coating with Ag on non-porous

TiSi10

Antibacterial effect against Escherichia coli and Staphylococcus epidermidis was

tested on all types of substrates and coatings. The coatings had the good

antibacterial effect against E. coli after 24 hours of interaction.

Acknowledgement

This work was supported by the Technology Agency for the Czech Republic within the project

TE01020390 Center for development of modern metallic biomaterials for medicinal implants.

References

[1] Novák P., Vojtěch D., Šerák J., Kubásek J., Průša F., Knotek V., Michalcová A., Novák M.: Chemické Listy 103 (2009) 1022-1026

[2] Novák P., Michalcová A., Šerák J., Vojtěch D., Fabián T., Randáková S., Průša F., Knotek V., Novák M.: Journal of Alloys and Compounds 470 (2009) 123-126

[3] Novák P., Popela T., Kubásek J., Šerák J., Vojtěch D., Michalcová A.: Powder Metallurgy. 54. 1 (2011) 50-55 [4] Horkavcová D., Běloubková T., Mizerová Z., Šanda L., Cílová Z., Častorálová M., Helebrant A.: Ceramic-

Silikáty. 56. 4 (2012) 314-322.

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MALAYAITE PIGMENTS DOPED BY Fe

Jana Luxová, Petra Šulcová

Department of Inorganic Technology, Faculty of Chemical-Technology, University of Pardubice,

Studentská 573, Pardubice 532 10, Czech Republic,


Keywords: Ceramic pigments, malayaite, optical properties, structure

Malayaites are calcium tin silicate compounds derived from natural occurring

mineral. The Malayaite has the chemical formula CaSnSiO5 or precisely

CaSnOSiO4. Malayaites together with titanites (CaTiOSiO4) belong to a large family

of the sphene pigments. Nowadays chromium doped malayaite-Ca(Sn,Cr)SiO5 is

the most important pink chromium pigment used in the ceramic industry for coloring

glazes and it is catalogued under number 12-25-5 in the CPMA classifications [1].

Compounds based on malayaite structure have a lot of good properties, but for

ceramic industry it is possible to emphasize their high temperature stability due

which malayaites are predestined as high temperature ceramic pigments [2].

Malayaite pigment with chemical formula CaSn0,9Fe0,1SiO5-δ has been investigated

in this work. Studied malayaite compound was prepared by solid state reaction. For

reaction CaCO3, SnO2, SiO2 and Fe2O3 were used. Homogenized initial mixtures

were calcined in the furnace at temperatures 1200-1500 °C with step 50 °C and

maintained for 4 hours on maximum temperature. The heating rate was 10 °C/min.

This study was focused on effect of firing temperature of colour properties of

prepared samples. The colour properties were detected for powdered samples and

for pigments after application in two binding systems, i.e. in transparent leadless

glaze P 07410 in 10 wt.% of pigment and in organic acrylate matrix in mass tone.

The colour properties were measured in visible region of the light (400 700 nm)

with using of the spectrophotometer ColorQuest XE (HunterLab, USA). Figure 1

shows the dependence of colour properties of studied pigment on calcination

temperature for all applications. From figure it is evident, that the colour coordinates

a* and b* increase with growing calcining temperature and the Fe-doped malayaite

compound is characterized by more red respectively yellow hue. Temperature 1350

°C as the most suitable firing temperature can be assessed.

The thermal stability of studied samples was verified by heating microscope EM201-

12 (Hesse-Instruments, Germany) which uses automatic image analysis. This

thermal analysis confirmed that in temperature region 30 1500 °C Fe-doped

malayaites are stable and therefore they can be used as ceramic high-temperature

pigments.

Sample prepared at 1350 °C was chosen for phase analysis. Two compounds were

identified on XRD pattern, namely malayaite as the major phase and Fe2O3 as

minor.

MALAYAITE PIGMENTS DOPED BY Fe J. Luxová , P. Šulcová


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Fig. 1: The influence of calcination temperatures on the colour properties of CaSn0,9Fe0,1SiO5-δ pigment

Particle size distributions were also characterised with using Mastersizer 2000/MU

(Malvern Inst., Ltd., GB). The mean particle size (d50) moved from 2.34 to 12.27 m

and generally it increased with calcination temperature up to 1350 °C (12.27 m),

then it decreased (around 8.77 m). This parameter has influence on optical

properties.

References

[1] Classification and chemical descriptions of the complex inorganic color pigments. 4th ed. Alexandria, Virginia:

Color Pigments Manufacturers Association, Inc. 2010; pp. 24.

[2] Harisanov V., Pavlov R.S., Marinova I.T, Kozhukharov V., Carda J.B.: Influnce of crystallinity on chromatic

parameters of enamels coloured with malayaite pink pigments, J. Eur. Ceram. Soc., 23 (2003) 429-435.

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SYNTHESIS AND STUDY OF COMPOUNDS BASED ON Bi-Zn-Ce-Nb

Kateřina Těšitelová, Petra Šulcová


Studentská 95, 532 10, Pardubice, Czech Republic


Keywords: ceramic pigments, pyrochlore, BZN system, thermal analysis

In the field of powder materials there are the current most discussed topic in the

environmental impact of the synthesized inorganic compounds. Commercial yellow

pigments (PbCrO4, PbMnO4, Pb2Sb2O7 and CdS) used in industry contain heavy

metals and are therefore in terms of environmental protection and health of

individuals inconvenient. Elimination of these pigments results in a significant is

narrowing the existing product range of colours. Increasing consumption of pigments

and developed pressure on manufacturing of economically acceptable pigments

leads to the study of new quality and economically interesting compounds. For this

reason the main attention has been directed to the synthesis of compounds based

on Bi-Zn-Ce-Nb.

Compounds based on Bi2O3ZnONb2O5 exhibit high dielectric constant and low

sintering temperatures. They belong to the Class-I dielectric group. The most

important compound in the BZN system is the Bi1.5Zn0.92Nb1.5O6.92 (BZN) cubic

pyrochlore [1,2]. BZN pyrochlore ceramics are now in the focus of researchers due

to their astonishing electronic applications. On the other hand, this presented thesis

is focused on the study of mixed oxide compounds based on Bi-Zn-Ce-Nb which

would be found useful as colour pigments. The main objective of this work was to

evaluate the prepared pigments from the point of their colour options for colouring

organic matrix, respectively ceramic glazes.

The compounds with formula Bi2ZnNbO7, Bi1.5Zn0.5Ce2O7, Bi2ZnCeO7,

Bi1.5Zn0.5CeNbO7, Bi1.5Zn0.5Ce0.5Nb1.5O7, Bi1.5Zn0.5Ce1.5Nb0.5O7, Bi1.5Zn0.5Nb2O7,

BiNbO4, BiNbO4 + 5 mol.% Bi2O3 and BiNbO4 + 10 mol.% Bi2O3 were prepared by

calcination of the powder raw materials, namely Bi2O3, ZnO, CeO2 and Nb2O5.

Precursors for the traditional solid state reaction were initially manually

homogenised in a porcelan mortar with the pestle. The homogenous mixture was

calcined in corundum crucibles in an electric furnace at the temperature 900 ºC and

with the heating rate 10 ºC/min. The calcination was maintained in isothermal

conditions for 2 hours. After spontaneous cooling, obtained pigments were applied

into organic matrix (dispersive acrylic paint Parketol, Balakom, a.s. Opava, CZ) in

mass tone and borate-silicate (transparent leadless) glaze G 070 91 (Glazura, s.r.o.,

Roudnice nad Labem, CZ).

The information about the temperature region for the formation of Bi1.5Zn0.5CeNbO7

was followed by thermal analysis using STA 449C Jupiter (Netzsch, Germany),

SYNTHESIS AND STUDY OF COMPOUNDS BASED ON Bi-Zn-Ce-Nb K. Těšitelová, P. Šulcová


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which allows the simultaneous registration of the thermoanalytical curves TG and

DTA. Colourimetric study of final coating films and coloured glazes were carried out

by a ColorQuest XE (HunterLab, USA) spectrophotometer in the 400 700 nm

wavelength range. The CIE L*a*b* system (1976) colour scales were used for the

assessment of the colour properties of the pigment powder. Phase composition of

pigment Bi1.5Zn0.5CeNbO7 was determined by diffractometer D8 Advance (Bruker

AXS, Ltd. Coventry, UK).

From thermoanalytical measurements, it follows that the starting mixture for the

synthesis of pigment Bi1.5Zn0.5CeNbO7 must be subjected to calcination

in a temperature range of 900 1,000 ºC. The colour properties of the pigments

applied to ceramic glaze provide yellow and yellow colour with little greenish tint.

The best results were obtained for the pigment Bi1.5Zn0.5CeNbO7 which indicates

high values of colour coordinate b* and chroma C in ceramic glaze (b* = 47.93 and

C = 47.97) and also in an organic matrix (b* = 46.84 and C = 47.16). Generally, the

pigments in an organic binder system in mass tone have worse colour properties

than the same pigments in ceramic glaze. Pigment Bi1.5Zn0.5CeNbO7 obtained at

900 ºC has crystalline structure but with low peak intensity and contain two phases –

monoclinic Zn0.67Nb1.33Bi2O7 and cubic CeO2.

Acknowledgements The authors would like to thank for the financial support IGA University of Pardubice

(SGSFChT_2015005).

References

[1] A. F. Qasrawi, B. H. Kmail, A. Mergen: Synthesis and characterization of Bi1.5Zn0.92Nb1.5-xSnxO6.92-x/2

pyrochlore ceramics, Ceramics International, 38 (2012) 4181. [2] L. Li, Y. Jin, P. Zhang: Improved dielectric properties of (Bi1.5Zn0.5)(Zn0.5Nb1.5)O7 by the substitution of Zn2+ on

the A site based on the structure characteristics, Ceramics International, 40 (2014) 10601.

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NONTRADITIONAL METHODS OF SYNTHESIS OF SPINEL

COMPOUNDS

Lenka Pluhařová, Ladislav Svoboda, Petr Bělina


Studentska 95, 53210 Pardubice


Keywords: spinel, ceramic pigments, thermal analysis

Nanocrystalline spinel compounds have been studied for decades, due to its wide

applicability as ceramic pigments, magnetic devices, semiconductors, refractories

and others. In order to obtain a powder of ceramics with high quality the factors of

the preparation methods are very important. Synthesis method and synthesis

conditions affect the chemical and physical properties of the final spinel compounds.

Spinel compounds have general formula represented by M2+

M3+

2O4 in which oxygen

ions form a face-centered cubic lattice with cations occupy interstitial positions

(tetrahedral M2+

-sites and octahedral M3+

-sites). The cation M2+

can be represented

by various ions of bivalent metals: Zn, Mg, Cu, Fe, Mn, Ni, Co, Ba, Sr, and Cd; and

cation M3+

by various ions of trivalent metals: Al, Fe, Mn, Cr, Zn, Ga, etc. [1, 2]

Research is focused on the synthesis CoMn2O4, MnCo2O4, MnFe2O4, CoFe2O4

compounds. Mentioned spinel compounds were prepared using three different

laboratory methods. The first method was a classic ceramic technique; it means

solid state synthesis at higher temperature. Starting materials (oxides, carbonates.

etc.) were mixed in a proper molar ratio using an agate mortar with pestle. Finally,

the samples were calcined at a various temperatures between 700 and 950 °C for

3 hours in open ceramic crucible in an electric furnace (heating rate 10 °C/min). The

second preparation method was based on decomposition of mixed carbonate

intermediates. As starting materials were metal sulphates, sodium carbonate used

and as surfactant was used PEG (50 wt.%). Starting materials were mixed in the

proper molar ratio using mortar grinder for 35 minutes. The prepared mixed

carbonates were subsequently calcined at 500 °C for 3 hours in open ceramic

crucible in an electric furnace (rate 10 °C/min). For the third method sodium oxalate

and metal sulphates were used as starting materials. Raw material with surfactant

PEG (50 wt.%) were mixed in the proper molar ratio using mortar grinder for

35 minutes. Prepared oxalates were subsequently calcined at 500 °C for 3 hours in

open ceramic crucible in an electric furnace (rate 10 °C/min).

The mechanism of spinel lattice formation was studied using the methods of thermal

analysis – differential thermal analysis and thermogravimetry (Jupiter STA, Netzsch,

DE). The mixture of starting materials or the intermediates powders were weighted

into open ceramic crucible and heated in temperature range from 30 °C to 700 °C or

1100 °C (heating rate 10 °C/min),


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Color properties, thermal stability, specific surface area, phase composition, particle

size distribution are only some of the studied properties of final powder materials.

On Figures are listed thermal analysis curves of the initial mixture for spinel

MnCo2O4 preparation. Figure 1 illustrates the decomposition of manganese

carbonate (endothermic peak 478 °C), Figure 2 illustrates the reduction of Co2O3 to

CoO (endothermic peak 923 °C). Figure 3 illustrates the analysis after

homogenization of the starting oxides, where it may be expected at a temperature of

754 °C the formation of spinel lattice means formation of solid solution.

Fig. 1: DTA-TG curves MnCO3

Fig. 2: DTA-TG curves Co2O3

Fig. 3 DTA-TG curves in air of prepared precursors using MnCO3 + Co2O3

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Reactions during preparation of the compounds CoMn2O4, MnCo2O4, MnFe2O4,

CoFe2O4 was studied using TG/DTA methods in order to determine the minimum

required calcination temperature. The results of the measurements are given in Table

1. In all cases the spinel lattice was formed, it was confirmed by XRD analysis. Using

XRD analysis in this case it is not possible to detect directly the particular

compound, only spinel structure was confirmed. The results of this work are

calcination temperatures listed in the text above.

Table 1: The minimum temperature required for formation of the spinel lattice depending on the method of preparation (determined by DTA)

Method/Compound MnFe2O4 CoFe2O4 MnCo2O4 CoMn2O4

Solid state reaction 754 812 627 891

Carbonate intermediates 385 317 387 396

Oxalate intermediates 419 391 413 429

Acknowledgements

The authors would like to thank for the financial support IGA University of Pardubice (SGFChT04).

References

[1] W.A.A. Bayoumy, Journal of Molecular Structure, 1056-1057 (2014), 285-291 [2] S. Naghibzadeh, M. A. Faghihi-Sani, S. Baghshahi, International Journal of Engineering and Advanced

Technology, 2 (2013), 606-609.

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WILLEMITE BASED FLUORESCENT MATERIALS

Peter Švančárek, Robert Klement, Lívia Dvorská, Dušan Galusek



Keywords: Zn2SiO4, willemite, Mn2+

, Eu3+

, green/orange-red emission, effect of charge

compensation, luminescence

Zinc silicate (Zn2SiO4) has been identified as a suitable phosphor host matrix with

excellent luminescent properties in the blue, green and red spectral regions and

excellent chemical stability. While being used as luminescent activators in phosphor

materials, rare-earth elements exhibit excellent luminescent characteristics, with

high emission efficiency and high colour purity, due to their specific electronic

configuration. Interesting candidate for red-emitting phosphors such as Eu3+

ion is

widely adopted as a red-emitting activator via the 5D0

7F2 transition at about

617 nm. In the present paper we report the preparation and the effect of Li+ ions

addition (co-doping) on microstructure and luminescence properties of Mn2+

and

Eu3+

phosphors with willemite host matrix.

The luminescent materials were prepared by mixing and annealing the powders of

ZnO, SiO2, Li2CO3 and/or MnO2/Eu2O3 at the compositions corresponding to

Zn0.97Eu0.03SiO3 Zn0.94Eu0.03Li0.03SiO3, Zn0.97Mn0.03SiO3 and Zn0.94Mn0.03Li0.03SiO3,

respectively. All chemicals used were of analytical grade or higher purity. The SiO2,

ZnO, Li2CO3 and MnO2/Eu2O3 fine powders were milled and homogenized together

in vibratory mill for 30 minutes. Obtained fine mixture of the powders was calcined

for 2 h at 1300 °C in air. For the SEM/EDX analysis (JEOL JSM-7600 Thermal FE

SEM) the powder was cast into phenolic conductive resin and polished by diamond

polishing disc. Both emission and excitation fluorescence spectra were measured by

Fluorolog 3 (FL3-21, Horiba) fluorescence spectrometer in front-face mode. The Xe-

lamp (450 W) was used as an excitation source. The phase composition was

determined using powder X-ray diffraction (PANalytical Empyrean Series 2 X-ray

diffractometer).

SEM shoved marked differences in morphology of the calcined powder product. The

particles of lithium free samples showed darker core with lighter shell. Addition of

lithium changed situation. Not only the shell expanded in volume, but lighter crystals

were present also in the core. This suggests the Li+ ions may promote the diffusion

of Zn2+

and other ions into the SiO2 particles. SEM-EDX mapping and point analysis

revealed that for Mn2+

-containing phosphors, the distribution of Mn2+

ions in the

samples is more homogeneous that in the case of Eu3+

doped samples; the

aggregates containing relatively high concentration of Eu3+

(up to 15 at.%) were

observed in both samples. It is well known that Mn2+

ions effectively substitute Zn2+

ions in the willemite structure due to the similarity of their ionic radii. The X-ray

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powder diffraction patterns of all four studied systems show the presence of

willemite (Zn2SiO4) as a major crystalline phase, together with traces of unreacted

SiO2 (cristobalite). The XRD patterns corresponding to Li+ undoped and co-doped

analogues are almost identical and no lithium containing phases were identified.

The photoluminescence emission spectra are shown in Fig. 1. The Mn2+

doped zinc

silicate (Fig. 1A) exhibits strong green emission centred at 523 nm when excited at

277 nm, corresponding to the d-level spin-forbidden transition of Mn2+

ion

(4T1g

6A1g). When co-doped with Li

+ ions, the intensity of green emission decreases

more than two times compared to the Li+-free analogue, Zn0.97Mn0.03SiO3. The

excitation spectra (not shown, emission monitored at 613 nm) of Eu3+

doped

systems Zn0.97Eu0.03SiO3 and Zn0.94Eu0.03Li0.03SiO3 exhibit, instead of broad charge

transfer band (CTB), some narrow excitation peaks originated from the f-f transitions

within the Eu3+

4f6 configuration. The most efficient excitation peak is the one

originated from 7F0

5L6 transition (∼394 nm), which means that this phosphor can

be effectively excited by UV LED chip.

425 450 475 500 525 550 575 600 625 650 675

exc

= 277 nm

Inte

ns

ity

(a

.u.)

Wavelength (nm)

Zn0.97

Mn0.03

SiO3

Zn0.94

Mn0.03

Li0.03

SiO3

(A)

4T

1g

6A

1g

550 575 600 625 650 675 700 725 750

(B)

Zn0.97

Eu0.03

SiO3

Inte

ns

ity

(a

.u.)

Wavelength (nm)

Zn0.94

Eu0.03

Li0.03

SiO3

exc

= 394 nm

5D

0

7F

0

5D

0

7F

3

5D

0

7F

2

5D

0

7F

1

5D

0

7F

4

Fig.1: The emission spectra of Mn2+ (A) and Eu3+ (B) doped phosphors.

The PL spectra of the studied Eu3+

doped samples, recorded under excitation at

394 nm, (Fig. 1B) show five major emission lines centred at app. 578, 590, 617, 650

and 700 nm corresponding to the 5D0

7FJ (J = 0, 1, 2, 3, 4) transitions,

respectively. It is common knowledge that 5D0 →

7F0,1 transition is directed by

selection rules for intermediate magnetic-dipole coupling J = 0, ±1, and the 5D0 →

7F2,4,6 are allowed electronic-dipole transitions. Generally, when the Eu

3+ ion

occupies the crystallographic site with inversion symmetry, its magnetic-dipole

transition 5D0 →

7F1 orange emission dominates in the emission spectrum, while the

electric-dipole transitions 5D0 →

7F2,4 red emission is dominant if the Eu

3+ ion is

located at an non-inversion center. Based on such general presumptions, Eu3+

ions

most likely occupy both inversion and non-inversion lattice sites in the studied

Zn0.97Eu0.03SiO3 and Zn0.94Eu0.03Li0.03SiO3 systems. Moreover, the exceptionally well

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resolved splitting of 5D0 →

7F1,2 transitions to three and five lines is observed due to

the Stark splitting in the strong crystal field. This also indicates that the symmetry at

Eu3+

ion site is low. The co-doping of Zn0.97Eu0.03SiO3 with Li+ ions leads to

significant increase of emission intensities by the factor of 2.5.

The opposite effect of Li+ co-doping on PL emission intensity in Mn

2+ and Eu

3+

doped systems can be explained by the formation of vacancies when substituting

Zn2+

ions by the Mn2+

/Eu3+

/Li+ ions. The formation of vacancies is not favourable for

the emission PL activator (Mn2+

/Eu3+

) because of the energy transfer from activator

to vacancy is more efficient, and hence a certain amount of vacancies will affect the

photoluminescence intensity. While in the case of Eu3+

doped systems the Li+ ion

decreases the number of Zn vacancies (charge compensation effect), in the Mn2+

doped systems most likely the Li+ ions increases the number of vacant sites.

Additionally, the strong polarisation effect of Li+ ion may also affect the close

coordination environment around the Mn2+

and structure of the host matrix.

Acknowledgements

The financial support of this work by the projects SAS-NSC JRP 2012/14 and VEGA 1/0631/14, is gratefully acknowledged. This publication was created in the frame of the project "Centre of excellence for ceramics, glass, and silicate materials" ITMS code 262 201 20056, based on the Operational Program Research and Development funded from the European Regional Development Fund.

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THE POSSIBILITIES OF HYDROMETALURGICAL USING

OF MANGANESE ORE FROM BIHAĆ REGION

Damir Hršak, Vladimir Bermanec1, Anita Štrkalj, Ljerka Slokar,

Zoran Glavaš, Dario Mašinović

University of Zagreb Faculty of Metallurgy, Sisak, Croatia 1University of Zagreb Faculty of Scince, Zagreb, Croatia


Keywords: manganese, hydrometallurgy, Bihać region.

Manganese reserves are predominantly in South Africa and former USSR [1].

Metallic manganese is predominantly producing by hydrometallurgical route [2].

Elemental manganese is hard and shiny and, like vanadium and chromium, is used

mostly to make steel alloys. A small amount of manganese, lower than 1 percent,

makes steel easier to roll, forge, and weld. Steel made with 12 percent manganese

is tough enough to be used for bulldozer buckets, and other extremely hard steel

objects. Small amounts of manganese are added to aluminium beverage cans and

bronze alloys to make them stiffer and tougher as well.

The chemistry of manganese resembles that of chromium in some respects. The

free metal is quite reactive and readily to reduce H+

from acids, forming the pale-pink

Mn2+

ion. Like chromium, manganese can use all its valence electrons in its

compound, exhibiting every possible positive oxidation state with the +2, +4 and +7

states most common. As the oxidation state of manganese rises, its valence states

electronegativity increases and its oxides change from basic to acidic. [3]

Hydrometallurgy of manganese

Manganese ores are used for making ferromanganese, pure MnO2 for batteries, or

to metallic manganese. Pyrolusite (manganese dioxide) is the most important

manganese mineral. Manganese nodules found at the bottom of the ocean are not

only an important source of manganese but also of copper, nickel, and cobalt. They

are composed mainly of manganese dioxide and iron hydroxide. The fact that the

nodules when collected contain appreciable amounts of water and high porosity

suggest that their treatment by pyrometallurgical methods would not be practical

because of the cost of drying.

While the processing of manganese ores is always for the recovery of its

manganese values, that for nodules may be carried out only for the recovery of

copper, nickel, and cobalt. The nodules may therefore be treated in practically the

same way as copper oxide ores or nickel laterites. Any leaching agent that

solubilized manganese dioxide will also solubilize copper, nickel and cobalt.

Manganese dioxide is insoluble in dilute sulphuric acid but when reduced to

manganese oxide, it readily dissolves to give a solution of MnSO4. Since

manganese is recovered from this solution by electrolysis the spent electrolyte can

THE POSSIBILITIES OF HYDROMETALURGICAL USING OF MANGANESE ORE FROM BIHAĆ REGION D.Hršak, V.Bermanec, A.Štrkalj, L.Slokar, Z.Glavaš, D.Mašinović


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be used for leaching. Manganese dioxide can be solubilized in dilute sulphuric acid

in presence of reducing agent such as ferrous sulphate, sulphur dioxide, coal or

oxalic acid which reacting by oxidation-reduction steps.

Manganese dioxide dissolves in hydrochloric acid liberating chlorine. The large

proportion of acid used per mole manganese dioxide necessitates its regeneration

from chlorine as well as from aqueous solutions. Although this technology may be

established for the treatment of manganese nodules yet it becomes costly because

of the large number of circuits involved.

Manganese Oxide Ore from Bihać Region

First preliminary investigation from manganese oxide ore from Bihać region was

indicated that this ore is suitable for exploitation in production of metallic

manganese.

Chemical analysis according to ISO/R319 standard indicated that content of

manganese in ore is 12.67 % of manganese. After thermic treatment content of

moisture was 2.45 % (heating 4 hour on 105 °C) and content of crystal water was

7.49 % (heating 3 hour on 670 °C).

In first experiments, after hydrometallurgical treatment with concentrated sulphuric

acid in period of 5 hours on temperature 80 °C, content of manganese in leaching

solution was 174.3 mg. Measurement of manganese concentration in leaching

solution was done by Inductively Coupled Plasma (ICP-OES) equipment. Sample of

ore has 2 g. In 2 g of ore there is 253.4 mg of manganese. Yield of the leaching

process is 68.8 % and there is lot of possibilities for optimization of the leaching

process.

A residual after leaching was analyzed by X-ray powder diffraction (XRPD). XRPD

measurement was performed on sample using a Philips X’pert Pro powder

diffractometer, with CuKα radiation filtered with a graphite monocrystal

monochromator, running at 40 kV and 40 mA.

On the basis of its XRPD pattern, the unleached residual contains predominantly

quartz (Fig. 1), both significant peaks. There is no evidence of significant content of

any manganese mineral.

Position [°2Theta]

10 20 30 40 50 60

Counts

0

200

400

600

800

Mn_BiH poslije izluzivanja

Fig. 1: X-ray diffraction pattern of residual after leaching.

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First preliminary results of hydrometalurgical treatment of manganese oxide ore from

Bihać region indicate that there is possibility to use researched ore for metallic

manganese production.

Hydrometallurgical treatment of manganese oxide ore by sulphuric acid as leaching

agent is just first step in process of metallic manganese production. It is reasonable

to continue researches and improve extraction of manganese from ore by finding the

optimal experimental conditions.

References

[1] F. Habashi, Metals from Ores, Metallurgie Extractive Quebec, Sainte-Foy, 2003.

[2] F. Habashi, A Textbook of Hydrometallurgy, Metallurgie Extractive Quebec, Sainte-Foy, 1999.

[3] M. S. Silberberg, Chemistry: The Molecular Nature of Matter and Change, McGraw-Hill, New York, 2003.

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MICROSTRUCTURE AND BASIC MECHANICAL PROPERTIES OF Si3N4

+ GRAPHENE PLATELETS COMPOSITES

Richard Sedlák1,2

, Alexandra Kovalčíková1, Zuzana Pramuková

1,2,

Pawel Rutkowski3, Ján Dusza

1

1Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, Košice, Slovakia 2Technical University of Košice, Faculty of Metallurgy, Letná 9, Košice, Slovakia

3AGH University of Science and Technology Krakow, Faculty of Material Science and Ceramics, Department of Ceramics and Refractories, al. Mickiewicza 30, Krakow, Poland


Keywords: microstructure, mechanical properties, silicon nitride, graphene platelets

The microstructure and basic mechanical properties of the graphene reinforced

composites with silicon nitride matrix with addition of 1 wt.% and 2 wt.% of graphene

platelets (GPLs) have been investigated and compared. Si3N4 material is commonly

used for producing of cutting tools, bearing balls, and other elements of devices

working in different conditions. Ceramic machinery working parts should have good

mechanical properties, such as bending strength, fracture toughness and abrasive

wear [1]. In the last few years, graphene has emerged as a promising reinforcement

material for an improvement of fracture toughness and wear resistance in composite

materials [2]. Dusza et al. [3] reported significantly improved fracture toughness of

Si3N4 ceramics reinforced with various GPLs. The aim of presented contribution is to

investigate the influence of the addition of GPLs on the microstructure development

and fracture toughness of Si3N4–GPLs composites.

Si3N4 – GPLs composites were prepared using commercial powders: Si3N4 (0.5 –

0.8 µm) and graphene (average flakes thickness 8 nm, 20–30 monolayers with

average particles size 550 nm). To activate the sintering the AlN (0.8 – 1.8 µm) and

Y2O3 (0.5–0.8 µm) powders, with the 2.5 wt.% and 4 wt.% were used. The powder

mixtures were homogenized for 6 h in propanol using a rotary-vibratory mill and

Si3N4 grinding media. Dried and granulated powders were hot-pressed at 1750 °C

for 1 h under 25 MPa in nitrogen flow. Sintered bodies with a diameter of 50 mm

were obtained. Apparent densities of the sintered samples were calculated by the

Archimedes method. Mechanical characterization was performed in terms of

measurement of basic mechanical properties of bulk materials such as hardness

and fracture toughness using indentation method, Young's modulus by resonance

method and bending strength from the four-point bending test. Indentation fracture

toughness (KIC) was calculated from the Shetty equation.

The SEM observations showed activation of toughening mechanisms in the form of

crack deflection, crack branching, crack bridging and graphene sheet pull-out. Fig. 1

and Fig. 2 show the microstructure of Si3N4–GPLs composites. As it can be observed

at fracture surfaces, microstructure consists of α-Si3N4 and β-Si3N4 grains with


MICROSTRUCTURE AND BASIC MECHANICAL PROPERTIES OF Si3N4 + GRAPHENE PLATELETS COMPOSITES R. Sedlák, A. Kovalčíková, Z. Pramuková, P. Rutkowski, J. Dusza

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132

0 50 100 150 200

5,0

5,5

6,0

6,5

7,0

KIC

[M

Pa.m

1/2]

Load [N]

Si3N

4 + 1% GPLs

Si3N

4 + 2% GPLs

0 50 100 150 200

14

15

16

17

18

19

20

21

22

HV

[G

Pa

]

Load [N]

Si3N

4 + 1% GPLs

Si3N

4 + 2% GPLs

majority predominance of β-Si3N4 grains. Well distribution of GPLs additives was

observed, but also cluster formation mostly in Si3N4 + 2 wt.% GPLs composite was

presented. Investigated materials were very well sintered, both of them achieved

more than 98 % of their theoretical density.

Fig. 1: Si3N4 + 1 wt.% GPLs Fig. 2: Si3N4 + 2 wt.% GPLs

The Vickers hardness (HV) of the composites ranged from 14.71 GPa at HV20 to

19.46 GPa at HV1 while KIC was in the range from 5.22 MPa.m1/2

at HV1 to

6.78 MPa.m1/2

at HV20. R-curves have been prepared for an evaluation of results.

Increase of load in Fig. 3. causes decrease of HV from 19.46 GPa to 15.83 GPa at

1 wt.% GPLs addition and from 19.44 GPa to 14.71 GPa at 2 wt.% GPLs addition.

Fig. 4. shows influence of the applied load on fracture toughness with increasing

trend from 5.27 MPa.m1/2

to 6.78 MPa.m1/2

at 1 wt.% GPLs addition and from

5.22 MPa.m1/2

to 6.45 MPa.m1/2

at 2 wt.% GPLs addition.

Fig. 3: Influence of the applied load on hardness Fig. 4: Influence of the applied load on fracture toughness

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GPLs addition leads to the hardness decrease while the fracture toughness

increase. Table 1 shows no significant effect of GPLs addition on Young's modulus

and bending strength was observed.

Table 1: Basic mechanical properties of Si3N4–GPLs composites

Acknowledgements

The authors gratefully acknowledge the financial support of projects VEGA 2/0043/14, VEGA

2/0189/15 and APVV-0108-12.

References

[1] Rutkowski P, Stobierski L, Zientara D, Jaworska L, Klimczyk P, Urbanik M. The influence of the graphene additive on mechanical properties and wear of hot-pressed Si3N4 matrix composites. Journal of the European Ceramic Society. 35 (2015) 87–94.

[2] Gutierrez-Gonzalez CF, Smirnov A, Centeno A, Fernández A, Alonso B, Rocha VG, Torrecillas R, Zurutuza A, Bartolome JF. Wear behavior of graphene/alumina composite. Ceramics International. 41 (2015) 7434–7438.

[3] Dusza J, Morgiel J, Duszová A, Kvetková L, Nosko M, Kun P, Balázsi C. Microstructure and fracture toughness of Si3N4 + graphene platelet composites. Journal of the European Ceramic Society. 32 (2012) 3389–3397.

Material Density [g/cm3]

Young's modulus [GPa]

Bending strength [MPa]

Si3N4 + 1 wt.% GPLs 3.191 313 534 ± 56

Si3N4 + 2 wt.% GPLs 3.171 301 560 ± 63

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THE EFFECT OF Mn2+ CONCENTRATION AND ZnO/SiO2 RATIO

ON LUMINESCENCE INTENZITY AND LUMINESCENCE DECAY

IN GREEN EMITTING PHOSPHOR Zn2SiO4:Mn2+

Robert Klement, Peter Švančárek, Dušan Galusek



Keywords: Zn2SiO4, willemite, Mn2+

, green emission, luminescence, luminescence decay

Zinc silicate (Zn2SiO4) is one of the most practical phosphors, and it has been

extensively studied for more than one century due to its high luminescence

efficiency, good color purity, and excellent chemical and thermal stability. It finds

wide applications in fluorescent lamps, cathode ray tubes, medical imagining

radiation detectors and plasma display panels. Guests-ion-doped Zn2SiO4 practical

phosphors emit from ultraviolet to visible/infrared depending on which ion is

incorporated (transition metal or lanthanide ions) and what phase it belongs to (-,

-Zn2SiO4). Commercial Zn2SiO4 : Mn phosphor exhibits intense green emission at

around 524 nm corresponding to Mn2+

spin- and parity-forbidden transition 4T1g(

4G)

6A1g(

6S). The lifetime for emission of Mn

2+ in Zn2SiO4 : Mn is known to be 25 ms,

which is too long than desired e.g. for TV applications. In the present paper we

report on the study of the effect of Mn2+

concentration and ZnO/SiO2 ratio on

luminescence intensity and luminescence decay in green emitting phosphors

Zn2SiO4 : Mn2+

.

The luminescent materials were synthesized by solid state reaction method from

precursor powders ZnO, SiO2 and MnO2; the homogenized powder mixture was

calcined for 2 h at 1300 °C in air. The prepared compositions correspond to

formulae Zn2-xSiO4 xMnO2, Zn1-xSiO3 xMnO2, and Zn1-xSi2O5 xMnO2, where x =

0.01, 0.03, 0.05; XRD revealed the presence of willemite (Zn2SiO4) as a major

phase in all studies systems. The emission and excitation fluorescence spectra were

measured by Fluorolog 3 (FL3-21, Horiba) fluorescence spectrometer in front-face

mode. The Xe-lamp (450 W) or Xe-flash lamp (25 Hz) were used as excitation

sources for steady-state and time-resolved measurements, respectively. For

comparison of emission intensities of prepared phosphors, the emission spectra

were also measured using an integration sphere (quanta-, Horiba).

The studied Mn2+

doped zinc silicates exhibit strong green emission centred at

523 nm when excited at 277 nm. The mechanism involved in the generation of a

green emission from Zn2SiO4 : Mn2+

under UV excitation is well understood, and can

be described as follows: the electron from the Mn2+

ground state 6A1g is excited to

the conduction band (CB) of willemite by UV photons, and the free electrons in the

CB can relax to the 4T1g excited state of Mn

2+ through non-radiative processes,

THE EFFECT OF Mn2+ CONCENTRATION AND ZnO/SiO2 RATIO ON LUMINESCENCE INTENZITY AND LUMINESCENCE DECAY IN GREEN EMITTING PHOSPHOR Zn2SiO4:Mn2+

R. Klement, P. Švančárek, D. Galusek


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135

which is then followed by radiative transitions from the 4T1g excited state to the

6A1g

ground state giving green emission. The integrated emission intensities for all

studied phosphors are shown on Fig. 1. The highest emission intensity was found for

the ZnSiO3 : Mn2+

phosphor. Whereas the well-known concentration quenching

effect operates in the ZnSiO3 : Mn2+

and ZnSi2O5 : Mn2+

phosphors, respectively, it

was not observed for the Zn2SiO4 : Mn2+

system in given Mn2+

concentration range.

The emission decay curves (in log scale) for the ZnSiO3:Mn2+

phosphors are shown

on Fig. 2. The similar trend in decay curves were also observed for the systems

Zn2SiO4 : Mn2+

and ZnSi2O5 : Mn2+

, respectively. The decay curves were fitted with

2-exponential function (Eqn.1) and selected data are summarized in Table 1; the

decay curves cannot be characterised using a single-exponential decay function.

The average decay time was calculated using Eqn. 2.

𝐼 = 𝐴1𝑒(−𝑡/𝜏1) + 𝐴2𝑒(−𝑡/𝜏2) (1)

𝜏𝑎𝑣𝑒 = ∫ 𝑡𝐼𝑑𝑡

∞0

∫ 𝐼𝑑𝑡∞

0

= ∫ 𝑡(𝐴1𝑒(−𝑡/𝜏1)+𝐴2𝑒(−𝑡/𝜏2))𝑑𝑡

∞0

∫ (𝐴1𝑒(−𝑡/𝜏1)+𝐴2𝑒(−𝑡/𝜏2))𝑑𝑡∞

0

= (𝐴1𝜏12 + 𝐴2𝜏2

2) (𝐴1𝜏1 + 𝐴2𝜏2 )⁄ (2)

0.30 0.35 0.40 0.45 0.50 0.55 0.60 0.65 0.70

ZnSi2O5

Zn2SiO4

1 mol.% MnO2

3 mol.% MnO2

5 mol.% MnO2

Inte

ns

ity

(a

.u.)

x(ZnO)

ZnSiO3

0 5 10 15 20 25 30 35 40 45 50

0,1

1

ZnSiO3 - 1 mol.% MnO

2


2


2

No

rma

lize

d in

ten

sit

y (

a.u

.)

Time (ms)

exc

= 277 nm

Fig. 1: Corrected integrated emission intensities

of studied phosphors at exc = 277 nm.

Fig. 2: Decay curves of ZnSiO3–X mol.% MnO2

(X = 1, 3, 5) at exc = 277 nm (Xe flash lamp)

Table 1: Decay times for selected samples obtained by decay curves fitting with 2-exp. function.

Zn:Si 1 (ms) A1 2 (ms) A2 ave (ms) 10% (ms)

ZnSiO3 – 1 mol.% MnO2 1:1 1.99 0.144 12.49 0.858 12.21 27

ZnSiO3 – 3 mol.% MnO2 1:1 1.92 0.318 10.52 0.686 9.85 20

ZnSiO3 – 5 mol.% MnO2 1:1 1.58 0.494 6.75 0.519 5.81 11

Zn2SiO4 – 5 mol.% MnO2 2:1 2.04 0.269 11.07 0.731 10.50 22

ZnSi2O5 – 5 mol.% MnO2 1:2 1.61 0.540 8.19 0.472 6.98 13

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After fitting we obtain a fast component 12 ms and a slow component 210 ms.

The longer decay time is expected to be due to the homogeneous distribution of

isolated non-interacting Mn2+

ions. The fast component most likely results from Mn2+

ions that are close enough to form pairs (Mn2+ Mn

2+) and Mn

2+ ions thus interact

with each other through dipol-dipol and/or exchange interactions. While a fast

component decreases moderately, the slow component decreases about 50 % with

increasing Mn2+

concentration. This suggests that with increasing dopant

concentration the distance between the Mn2+

ions decreases and results in the

concentration quenching of the luminescence. The decay behaviour is comparable

for ZnSiO3 : Mn2+

and ZnSi2O5 : Mn2+

systems, while for Zn2SiO4 : Mn2+

system the

decay behaviour (1 and 2) is less affected by increasing Mn2+

concentration.

Acknowledgements






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EMISSION PROPERTIES OF Eu3+/Eu2+ PHOSPHORS IN THE SYSTEM

Y2O3Al2O3

Lívia Dvorská, Katarína Haladejová, Robert Klement, Anna Prnová,

Dušan Galusek


e-mail: [email protected], [email protected]

Keywords: photoluminescence, Eu3+

/Eu2+

doped glass, yttrium aluminate glasses, LED, phosphor, YAG, Y3Al5O12

Yttrium aluminium garnet (Y3Al5O12, YAG) has been recognized as one of the best

phosphor host materials for rare earth (RE) ions because of its optical transparency

in the range of ultraviolet to infrared. YAG is a host with excellent structural

compatibility. The Y3+

and Al3+

ions can be substituted by many kinds of cations with

different sizes and valence in a certain extent. Among the RE ions, trivalent

europium (Eu3+

) has been identified as an excellent luminophore and it has been

investigated for use in polychromatic displays and Hg-free lamps. Efforts to enhance

the Eu3+

luminescence have been concentrated in incorporating this ion into host

materials with low phonon energy (e.g. YAG, Al2O3). The one of the possible method

to prepare the polycrystalline phosphor material by controlled crystallization starts

from glass. In this case we can combine the photoluminescence properties of

luminophore hosted in polycrystalline and glass matrix.

In the present work we report the results of a preliminary study on the preparation

and characterization of Eu3+

doped yttrium aluminate glasses and their

polycrystalline analogues with composition of YAG (Y3Al5O12) and eutectic

composition in the system Y2O3Al2O3 containing 76,47 mol.% (60 wt.%) Al2O3 and

23.01 mol.% (40 wt.%) Y2O3, denoted as Y40A60. The doping level was 0.5 mol.%

Eu2O3 corresponding to 1.0 at.% of Eu3+

. For comparison also the Eu3+

doped Al2O3

matrix was prepared. The glasses were prepared by flame-spraying technique from

precursor powders synthesized by the Pechini method. The Eu2+

analogues were

prepared by reduction of samples in N2H2 atmosphere. The steady-state and time-

resolved luminescence properties of the prepared phosphors were studied in details.

Acknowledgements






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138

EFFECT OF TRIS AND HEPES BUFFERS ON THE GLASS-CERAMIC

SCAFFOLD

Diana Horkavcová1, Dana Rohanová

1, Iva Březovská

1, Pavlína Bozděchová

1,

Aldo Roberto Boccaccini2

1Department of Glass and Ceramics, Faculty of Chemical Technology, University of Chemistry and

Technology Prague, Technická 5, 166 28 Prague, Czech Republic 2Department of Materials Science and Engineering, Institute of Biomaterials, University of Erlangen-

Nuremberg, 91058 Erlangen, Germany


Keywords: SBF, buffer, TRIS, HEPES, glass-ceramic, in vitro test

An in vitro test is commonly used for a prediction of a biomaterial ability to produce

mineral hydroxyapatite on its surface. Tested biomaterial is exposed into simulated

body fluid (SBF) [1]. SBF (prepared according ISO 23317 standard) mimics an

inorganic part of a blood plasma and TRIS buffer tris(hydroxymethyl)amino-

methane) is used for the pH = 7.45 keeping. Recently, we found [2] that TRIS buffer

reacts with tested glass-ceramic material. Additionally, TRIS forms a soluble

complex compound with Ca2+

ions, and thus it influences the in vitro test of glass-

ceramic material. The aim of this work was to find whether HEPES buffer

(C8H18N2O4S) could be able to maintain the pH of the simulated body fluid during the

in vitro test.

Glass-ceramic scaffold (Fig.1) was exposed to buffer´s solutions HEPES and TRIS

(diluted in demineralised water) under a dynamic arrangement of the test

(48 ml.day-1

) for two weeks. Concentration of Ca2+

ions in the leachates was

analysed by AAS and (PO4)3-

ions by spectrophotometry. The scaffold surface

before and after interaction was characterized by SEM/EDS and XRD.

Fig. 1: Original scaffold (SEM/EDS)

EFFECT OF TRIS AND HEPES BUFFERS ON THE GLASS-CERAMIC SCAFFOLD D. Horkavcová, D. Rohanová, I. Březovská, P. Bozděchová,A. R. Boccaccini


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In the case of the interaction of glass-ceramic with TRIS buffer solution, crystalline

and a glassy phase of the scaffold were partly dissolved (increasing of Ca2+

, (PO4)3-

and Si). However, HEPES buffer also significantly influenced the dissolution of the

glass-ceramic scaffold. Analysis of the both types of leachates confirmed the

dissolution of the glass-ceramic scaffold immediately from the beginning of the test.

Surprisingly, hydroxyapatite (HAp) was formed in the both types of the model

solutions, although an inorganic part of blood plasma was not present. The scaffold

surface was changed into a highly porous structure (Fig.2 and 3).

Fig. 2: Scaffold after two weeks in TRIS solution Fig. 3: Scaffold after two weeks in HEPES solution

Both buffers significantly affected the dissolution of the tested material and moreover

they were not able to keep a neutral value of pH during the test.

Acknowledgement

This work was supported by the Technology Agency for the Czech Republic within the project

TE01020390 Centre for development of modern metallic biomaterials for medicinal implants.

References

[1] International standard ISO 23317, Implant for surgery – in vitro evaluation for apatite-forming ability of implant materials, Third edition, 2014.

[2] Rohanová D., Boccaccini A.R., Yunos M.D., Horkavcová D., Březovská I., Helebrant A.: Acta Biomaterialia. 7 (2011) 2623-2630.

[3] Chen Q.Z., Thompson L.D., Boccaccini A.R.: Biomaterials 27 (2006) 2414-2425. [4] Rohanová D., Boccaccini A.R., Horkavcová D., Bozděchová P., Bezdička P., Častorálová M.: Journal of

Materials Chemistry B 2 (2014) 5068-5076 [5] Rohanová D., Horkavcová D., Helebrant A., Boccaccini A.R.: Journal of Non-Crystalline Solids (in press)

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INORGANIC-ORGANIC NANOFIBROUS MATERIALS

FOR BIOMEDICAL APPLICATIONS

Ivana Veverková1, 2

, Irena Lovětinská-Šlamborová2, Petr Exnar

3

1 Textile Faculty, Technical University in Liberec, Studentská 1402/2, 461 17 Liberec 1 2 Institute of Healthy Studies, Technical University in Liberec, Studentská 1402/2, 461 17 Liberec 1

3 Institute for Nanomaterials, Advanced Technologies and Innovation, Technical University in Liberec, Studentská 1402/2, 461 17 Liberec 1


Keywords: silica nanofibers, biomedical applications, toxicity, dissolving

Next to organic nanofibers, whose utilization is widely studied and tested, the last

few years, attention is also paid to research new features and biomedical application

possibilities of inorganic nanofibers. Our development is focused on silica nanofibers

on a silicon dioxide (SiO2) base and their biomedical applications.

Silica nanofibers are produced by sol-gel method (Fig. 1). This method is mainly

used for the preparation of inorganic oxidic materials (silicates and similar

materials), or for the preparation of organic-inorganic composite materials, which are

difficult to obtain by other methods. Sol-gel method is based on the transition phase

of the polymer network in a colloidal suspension, sol and a subsequent gelation of

the resulting sol to form a porous spatial network in the liquid phase, the gel. The

starting materials for the production of special materials prepared by sol-gel method

are alkoxides derived from alcohols. [2] Sol is subsequently spinned through

electrospinning. Electrospun silica nanofibres are formed with a diameter of 150

600 nm.

Due to the potential use of the nanofibers for biomedical applications, it is important

to consider a potential risk for human organism. The important indicator of

nanofibers´ toxicity is its ability to resist the physiological degradation in the body -

biopersistence. According to available research, nanofibers can be considered toxic

if their dissolution rate is in the order ng.cm-2

.h-1

. For the bio-soluble fibers are then

considered fibers with a dissolution rate less than 1000 ng.cm-2

.h-1

. [5]

"In vitro" biopersistence tests monitor dissolution of nanofibers in simulated body

fluids and physiological saline solution. Degradation of silica nanofibers is evaluated

depending on different process conditions (time of dissolution, temperature of

simulated body fluids and thermal stabilization of silica nanofibers).

Silica nanofibers are able to dissolve in wound in seven days. Bound functional

agents are gradually released (e.g. antibiotics, enzymes, disinfectants). Immobilized

antibiotics compared with the use of antibiotics per os do not burden the patient. The

inert substrate contains only silicon dioxide which is based on available studies

nontoxic to the organism, biodegradable. [6, 3]

Product of silica nanofibers degradation is silicon (Si). Simple forms of silicon such

as silicic acid or orthosilicic acid, which is commonly present in drinking water, beer

INORGANIC-ORGANIC NANOFIBOROUS MATERIALS FOR BIOMEDICAL APPLICATIONS I. Veverková, I. Lovětinská-Šlamborová, P. Exnar


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or other liquids, only very little interact with the mucosa in the intestine. Orthosilicic

acid is quickly absorbed in the small intestine into the blood. When silicon gets into

the blood, the body has a very effective mechanism for the elimination in the urine.

High silicon income almost does not effect its concentration in the blood, while

silicon concentration in urine is changed rapidly and considerably. However, for

majority of population with normal renal function the normal intake of dietary silicon

from foods and water has not been associated with any known toxicity. There are no

known symptoms or diseases of silicon excess or deficiency in humans. [6]

Nanofibres substrate is able to adhere to the wound; it can curl up and put into deep

wounds. As proved by practice, the material shows excellent compatibility with

patients.

Fig.1: SEM image of silica nanofibers, the initial state and after dissolving in physiological saline solution (36 °C, 168 hours), Technical University in Liberec.

Acknowledgements

The research presented was supported by the Ministry of Education, Youth and Sports in the framework of the targeted support of the “National Programme for Sustainability I” LO 1201 and the OPR&DI project “Centre for Nanomaterials, Advanced Technologies and Innovation”, CZ.1.05/2.1.00/01.0005. References

[1] Y. Dai., W. Liu, E. Formo, Y. Sun, Y. Xia: Polymers for Advanced Technologies. 22 (2011) 326-338. [2] J. Studničková, P. Exnar, J. Chaloupek: Silicone dioxide nanofibers. 13th International Conference: Structure

and structural mechanics of textiles. 2006, 13. [3] F. Schircliff. Characterization of organosilane-modified silicon/silicon dioxide systems for biological and

nanotechnology applications. Golden, Colorado. Diploma thesis. Colorado school of Mines. 2013. [4] Y. C. Xie, A.S. Hill, Z. Xiao, H. Militz, C. Mai: Composites Part A: Applied Science and Manufacturing. 41

(2010) 326-338. [5] T.W. Hesterberg, G.A. Hart.: Healt&Safety Aspects of Fiber Glass, Battery Conference on Applications and

Advances, Long Beach,CA, SA,2000. [6] R. Jugdaohsingh: J Nutr Health Aging . 11 (2007) 99-110.

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IMMOBILIZATION OF BIOMOLECULES ON SILICA NANOFIBERS

Iveta Danilová1, 2

, Irena Lovětinská-Šlamborová 2, 3

, Petr Holý 2

1 Textile Faculty, Technical University in Liberec, Studentská 1402/2, 461 17 Liberec 1

2 Institute of Healthy Studies, Technical University in Liberec, Studentská 1402/2, 461 17 Liberec 1

3 Faculty of Science, Humanities and Education, Technical University in Liberec, Studentská 1402/2,

461 17 Liberec 1, Czech Republic


Keywords: immobilization, silica nanofibers, biomolecules, proteolytic enzymes

Process of biomolecules immobilization means binding of active molecules to a

substrate – nanofibers. Nanofibers are a very promising material for immobilization

due to their size, large surface area and high porosity. This material has at least one

dimension corresponding to molecular size (from units to hundreds of nanometres).

Drugs (antibiotics, analgesics, cytostatics), proteins, enzymes and other chemical

substances (growth factors and disinfectants) are classified into biomolecules

suitable for immobilization. These molecules can be linked by weak bonds or strong,

covalent bonds on the nanofibres surface. Moreover, physical adsorption,

encapsulation or crosslinking can be used for such purposes [1].

For biomedical applications, fibres must be biocompatible, that means

biodegradable and nontoxic for living organism. Testing of inorganic nanofibres

showed their excellent biocompatible properties. We produce silica nanofibres by

synthesis of sol by sol-gel method from tetra-alkoxysilane, followed by

electrospinning and subsequent thermal stabilization [2]. Time of fiber degradation

and time of gradual release of biomolecules can be affected by changes in

conditions of stabilization (different temperature, stabilization time).

All these properties can be applied in treatment of chronic wounds, pressure sores,

burns and enzymatic debridement. Especially for some proteolytic enzymes have

been demonstrated very good catalytic properties in the healing of burns, bedsores,

large wounds or skin debridement [3, 4]. Very good therapeutic results in

dermatology are also achieved by immobilization of broad-spectrum antibiotics

(tetracycline) on silica nanofibres. The advantage is a non-invasive method (simply

attaching of nanofibre substrates to a wound) where the drug is gradually released

directly to the wound bed. In addition, this method reduces the side effects of

antibiotics on whole organisms.

For effective and prolonged biomolecules immobilisation by covalent bonds, it is

usually necessary to use reagents to functionalise and activate the fiber surface (Fig.

1). By surface functionalization, the number of reaction groups for further binding is

increased. Alkoxysilanes such as (3-aminopropyl)tri-ethoxysilane, (aminopropyl)tri-

methoxysilane, (3-mercaptopropyl)tri-methoxysilane or (3-iso-cyanatopropyl)tri-

ethoxysilane are most often used for surface functionalization [4].

IMMOBILIZATION OF BIOMOLECULES ON SILICA NANOFIBERS I. Danilová, I. Lovětinská-Šlamborová, P. Holý


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Fig. 1: The principle of activation and functionalization on the nanofibers surface [5].

Functionalized nanofiber surface is then activated by reagents that improve the

efficiency of the method. Most common activating agents are halides, bi-functional

reagents or sulfones. Their reactive groups (—NH2, —SH, —CN) react directly with

enzymes and antibiotics. Especially for enzymes is very important to avoid binding

of the nanofiber substrates to active sites of enzymes. These enzymes then lose its

catalytic function. Maintaining of activity and stability of the immobilized enzyme is

also dependent on the optimum pH and temperature; it is different for individual

enzymes.

In Figure 2, there is a diagram of esterase immobilization from porcine liver to silica

nanofibres. For silanization reaction we applied (3-aminopropyl)tri-ethoxysilane and

as bi-functional reagent we used glutaraldehyde [6].

Fig. 2: Immobilization of the enzyme esterase – scheme [6]

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Selected SEM images of the silica nanofibers samples with and without immobilized

esterase enzyme are shown in Figures 3a and 3b. In Figure 3b, there are clearly visible

structures with size of 10 40 nm, covering continuously the whole surface of silica

nanofibres.

Fig. 3: Scanning electron microscope images of silica nanofibers without immobilized enzyme (a) and with immobilized esterase (b), magnification 50000x [6].

Immobilization of biomolecules on a substrate and their application to hard healing

wounds has a large number of advantages. It certainly makes a sense to continue in

development of other immobilization methods. We should not forget to test produced

materials (reaction on contact with the skin, the solubility of fiber).

Acknowledgement The research presented was supported by the Ministry of Education, Youth and Sports in the framework of the targeted support of the “National Programme for Sustainability I” LO 1201 and the OPR&DI project “Centre for Nanomaterials, Advanced Technologies and Innovation”, CZ.1.05/2.1.00/01.0005.

References

[1] K. Drauz, H. Waldmann. Enzyme catalysis in organic synthesis: a comprehensive handbook. 2nd, 2002, vol. 37, ISBN 35-272-9949-1.

[2] P. Exnar, Book of Extended Abstracts, PANMS (Potential and Application of Nanotreatment of Medical Surfaces), 2011, p. 21-23. ISBN 978-80-7372-756-7.

[4] F. Gomes, C. De V. Spínola, H. A. Ribeiro et al. Burns, 36 (2010) 277-283. [5] P. Zucca, E. Sanjust. Molecules, 19 (2014) 14139-14194. [6] H. S. Yoo, T. G. Kim, T. G. Park. Advanced Drug Delivery Reviews, 61 (2009) 1033-1042. [7] I. Danilová., I. Lovětinská-Šlamborová, V. Zajícová et al. Vlákna a textile, 21 (2014) 3-11.

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NANOINDENTATION OF (Ti,Ta) CARBONITRIDE CERMETS

WITH VARIOUS COBALT BASED BINDERS PREPARED

BY MECHANOCHEMISTRY

Martin Fides1,2

, Pavol Hvizdoš1, Michal Novák

1,3

1Slovak Academy of Sciences, Institute of Materials Research, Watsonova 47, Košice, Slovakia 2Technical University of Košice, Faculty of Metallurgy, Letná 9, Košice, Slovakia

3Slovak University of Technology in Bratislava, Faculty of Materials Science and Technology,

Bottova 25, Trnava, Slovakia


Keywords: nanoindentation, cermets, cobalt binders, mechanochemistry

Cermets belong to the group of ceramic-metal composite materials used mainly for

the application as cutting tools [1]. Cermets, usually, consist of titanium carbide (TiC)

or titanium carbonitride (TiCN) particles distributed in Co or Co/Ni binders. Cermets

combine the advantages of ceramic and metallic materials (ceramic phase provides

high hardness and good wear resistance while the binder phase brings fracture

strength and impact resistance) [2,3]. TiTaCN based cermets are very attractive

materials which combine all of the mentioned properties.

The aim of the work deals is nanoindentation evaluation of mechanical properties

(hadness and elasticity) of various TiTaCN based cermets with different cobalt –

graphite binders.

Three types of TiTaCN based cermets embedded in cobalt based binder with

different amounts of graphite addition prepared by mechanochemistry were studied.

The prepared materials contained graphite additions of 0, 1.8, and 2.2 %.

Nanoindentation technique was used for the determination of nanohardness (HIT)

and indentation modulus (EIT). Nanoindentation tests were carried out using the

single loading-unloading regime at the maximum loads from 5 mN up to 50 mN with

constant loading rate (30 s to maximum). For the obtaining of sufficient statistics, 20

indents at each load were prepared. Additional atomic force microscopy (AFM)

observations were provided with the aim to visualize the indents presented in the

binder phase.

Fig. 1 shows the 3D representation of an indent in cobalt-graphite binder made by

5 mN load studied by atomic force microscopy (AFM). Significant pile-up around the

indented surface is visible showing ductile character of the binder phase.

Fig. 2 shows the dependence of nanohardness and indentation modulus on the

applied indentation load. No significant load size effect (LSE) can be found. It means

that measured values ranges in the approximately identical values. However, the

effect of graphite addition into the cobalt binder is significant. Average values of HIT

and EIT for the lower loads (5 mN, 10 mN, 20 mN) drop down from approx. 19 GPa


NANOINDENTATION OF (Ti,Ta) CARBONITRIDE CERMETS WITH VARIOUS COBALT BASED BINDERS PREPARED BY MECHANOCHEMISTRY M. Fides, P. Hvizdoš, M. Novák

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and 430 GPa at zero graphite addition to 13.5 GPa and 400 GPa at 2.2 % of

graphite addition.

a) b)

Fig.1: Indent (5 mN) presented in cobalt – graphite binder phase from a) AFM, b) SEM.

Fig. 2: a) Nanohardness (HIT) and b) indentation modulus (EIT) as the function of applied load.

The observed behaviour is in good agreement with the results of authors [1], but the

values measured by them are slightly different. For 5 mN loads the authors

presented hardness drop from 17.2 GPa at 0 % addition down to approx. 8.6 GPa at

2.2 % addition of graphite.

No significant effect of different load was observed. The addition of graphite into the

cobalt binder phase resulted in decrease of nanohardness and indentation modulus.

With increasing graphite addition the nanohardness decreased from ~19 GPa to

13.5 GPa and corresponding indentation modulus decreased from ~430 GPa to

400 GPa, respectively.

Acknowledgement

The work was financially supported by the projects VEGA 2/0075/13 and APVV-0108-12.

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References

[1] E. Chicardi et al: Toughening of complete solid solution cermets by graphite addition, Chemical Engineering Journal, 2015, doi: http://dx.doi.org/10.1016/j.cej.2015.01.022, Accepted for publication.

[2] H. Pastor: Titanium – carbonitride – based hard alloys for cutting tools, Materials Science and Engineering, 105 (1988) 401 – 409.

[3] J. M. Cordoba et al: Liquid – phase sintering of Ti (C, N) Cermets. The effects of binder nature and content on the solubility and wettability of hard ceramics phases, Journal of Alloys and Compounds,. 559 (2013) 34 – 38.

http://dx.doi.org/10.1016/j.cej.2015.01.022

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RHEOLOGICAL PROPERTIES OF CASTABLES

Jan Urbánek, Jan Macháček, Jaroslav Kutzendörfer, Jiří Hamáček

Department of Glass and Ceramics, University of Chemistry and Technology, Prague; Technická 5,

166 28, Prague 6, Czech Republic


Keywords: castables, rheology, solidification, viscometer.

Refractory castables are usually delivered in dry form. They consist of aggregates

and binders and they solidify when they are mixed with water [1]. According to their

rheological properties and a way of installation they are divided into three types –

self-flowing, vibrate-flowing and ramming.

Rheological properties of castable mixtures are influenced by their compositions

(aggregates, binders – including w/s ratio (water/solid), micro-additives, dispersants,

regulators of solidification, fibers). They are also influenced by size and surface of

particles, amount of refractory mixture and of course a temperature.

A liquescent test (as well solidification) belongs to classical methods of

measurement of rheological properties of castables. A truncated cone filled with

refractory castable mixture is lifted and the size of wetting out of mixture is

measured in two perpendicular ways. Another test is represented by measuring the

penetration of the cylinder with cone ending or Vicat`s needle into the material. A

use of a rotary viscometer is a modern method, which was used for measurement of

fine suspensions. Castables represent heterogeneous mixtures with a size of

particles of aggregate from micrometers to millimeters sometimes to centimeters.

Hence it is necessary to modify the construction of rotary viscometer.

A rotary viscometer Haake Mars III was used. Because the plate-plate, cone-plate,

cylinder-cylinder setting generates too big noise at signal, we used a vane rotor-

cylinder setting. We used a vane rotor with sheets 11 mm long and 16 mm high. We

constructed a cylindrical vessel with 7 cm in diameter and 6.5 cm high to avoid

generating a noise at signal again. Then we constructed reducing plate, which

allows connecting the cylindrical vessel to a base static plate of rotary viscometer.

It was necessary to modify relations used in calculating of shear rate and of shear

stress for the constructed setting. Because a vane rotor makes due to its rotation the

cylinder, so the setting was approximated to a system cylinder in infinite space.

Influence of a rail, which holds vane rotor in a material, was calculated. It was

possible to omit it because it was negligible. The constructed mechanism can be

used to measure rheological properties of castables including measurements of an

initial period of solidification for a size of aggregate up to 4 mm. Thanks to that it

creates an alternative to classical methods.

For a selected castable we have measured influence of ageing agents on ageing

speed, as you can see in Figure 1a-c.

RHEOLOGICAL PROPERTIES OF CASTABLES J. Urbánek, J. Macháček, J. Kutzendörfer, J. Hamáček


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Fig. 1: Influence of time on shear stress during the constant shear stress with different amount of ageing accelerator (a – the least one, c – the most one).

In figures 1a-c we can see the beginning of the so-called acceleration period, which is

represented by shear stress growing. When shear stress achieves 10 kPa, then the

viscometer automatically stops a rotation and the vane rotor is necessary to move

out of a mixture not to harden in it. In comparison with classical methods of

measurement rheological properties at the same setting time, the truncated cone of

castables doesn`t flow and holds its shape and the Vicat`s needle doesn`t penetrate

into a material. Therefore the device allows to substitute classical methods.

An influence of dispersant on rheological properties of material was measured as

well, as you can see in a Figure 2.

Fig. 2: Influence of amount of flow agents on shear stress at constant shear rate

It was possible to measure rheological properties of castables due to adjustment of

the Haake Mars III device. For example it has been measured influence of

solidification agents on solidification speed and influence of dispersants on

rheological properties of material.

Acknowledgement

This work was supported by the grant TA03010849 realized under financial support of the Technology Agency of the Czech Republic.

References

[1] J. Kutzendörfer, F. Tomšů, (2008). Žárovzdorné materiály Díl I. Praha. [2] F. Tomšů, Š. Palčo, (2009). Žárovzdorné materiály Díl IV. Praha. [3] Instruction Manual Software HAAKE RheoWin. Karlsruhe: Thermo Fisher Scientific. [4] R.P. Chhabra, J. R. (2008). Non-Newtonian Flow and Applied Rheology. Oxford: Elsevier Ltd.

0

50

100

150

200

250

0 0.2 0.4 0.6 0.8 1 1.2

τ [P

a]

mflow agents [g/400g dry mixture]

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NANOCOMPOSITE BASED ON CARBON

NANOTUBES AND MONTMORILLONITE

Michal Hubeňák1, Magdaléna Kadlečíková

2, Juraj Breza

2, Karol Jesenák

1,

Michal Kolmačka2

1Department of Inorganic Chemistry, Faculty of Natural Sciences, Comenius University, Mlynská

dolina, Ilkovičova 6, 842 15 Bratislava, Slovak Republic 2Department of Microelektronics, Faculty of Electrical Engineering and Information Technology, Slovak

University of Technology, Ilkovičova 3, 812 19 Bratislava, Slovak Republic


Key words: montmorillonite, carbon nanotubes, chemical vapour deposition

Carbon nanotubes (CNTs) are due to their outstanding mechanical properties

considered as a promising fibrous material for novel ceramic materials with better

mechanical performance [1,2,3]. Montmorillonite as a major part of bentonite is

widely used in ceramic production. In our work we tried to enrich montmorillonite

with CNTs. Relation between mineral morphology and carbon phase was

investigated to obtain information about prepared compound such as shape, length,

diameter and concentration of the carbon nanotubes on the mineral surface. In this

study, we demonstrate the synthesis of carbon nanotubes as a potencial

reinforcement for ceramic materials on clay mineral. The hot filament chemical

vapour deposition method (HF CVD) for synthesis of carbon nanotubes was used.

Montmorillonite with particle size under 2 m was insulated from bentonite from

Stará Kremnička, Jelšový Potok, Slovakia. Detailed information about bentonite can

be found in [4]. Prepared montmorillonite was then deposited on silicon wafers by

sedimentation. Fe(NO3)3 from water solution was used as catalyst.

HF CVD method was used to synthetize CNTs. Temperature during the experiment

was 600 °C and temperature of tungsten filament was 2020 °C. Pressure inside of

reactor was 3 kPa. Time of deposition was 25 minutes.

The morphology of the prepared composite which consist of inorganic matrix and

carbon phase was examined by scanning electron microscopy (SEM) JEOL 7500 F.

Product of the synthesis is visible on Fig. 1 and 2. Surface of montmorillonite is

covered with dense network of chaotically oriented CNTs which are in various

shape, diameter and length. The longest fibres have about 20 m. For now it is hard

to say, if the CNTs are also inside of montomorillonite.

NANOCOMPOSITE BASED ON CARBON NANOTUBES AND MONTMORILLONITE M. Hubeňák, M. Kadlečíková, J. Breza, K. Jesenák,M. Kolmačka


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Fig. 1, 2: SEM image of montmorillonite surface after synthesis

Acknowledgements

The work has been supported by grant VEGA VEGA 1/0601/13 of the Ministry of Education of

the Slovak Republic and by grant Comenius University in Bratislava UK/496/2015.

References

[1] E. Zapata-Solvas, et. al., J. Eur. Ceram. Soc., 32, 12, 2012, p. 3001 – 3020. [2] Z. Hu, et. al., Ceram. Int., 39, 2, 2013, p. 2147 – 2152. [3] J. Ping Zhou, et. al., Mater. Sci. Eng. A., 520, 2013, p. 153 – 157. [4] O. Hanzel, et. al., J. Eur. Ceram. Soc., 34, 2014, 1845 – 1851. [5] K. Jesenák, et. al., Mineralia Slovaca, 27, 1995, p. 221.

TECHNICAL UNIVERSITY of KOSICE

152

Faculty of Metallurgy

Department of Ceramics

in cooperation with Institute of Geotechnics

of Slovak Academy of Sciences, Košice Slovak Silicate Society, Slovak Glass Society

Dear Madame, Sir,

you are invited to Herľany - Spa near Košice,

on International Conference

PREPARATION OF CERAMIC MATERIALS.

SCOPES OF THE CONFERENCE:

Refractories, Fireproof Ceramics

Construction and Building Ceramics

Fine Ceramics

Surface Treatments of Ceramics

Glass, Glass-ceramic materials

Language: English

Conference Contributions will include invited

lectures (30 min) and contributed papers (15 min). Poster Section includes competitive works of young scientists and students that will be evaluated with commission of “Poster section of young

scientists and students” on this conference.

All details on application form, instruction for Authors, ect. will be given on after 10th

January 2017 on page www.tuke.sk/seminar_PKM.

All correspondence send on e-mail: [email protected],

(address: Technical University of Kosice, Faculty of Metallurgy, A.Prof. B.

Plešingerová, Letná 9, 042 00 KOŠICE, Slovak Republic), tel.: +421/ 55/ 602 2303.

Organising Committee: A. Prof. Ing. P. Vadász, CSc., TU Košice, FM DC, Slovakia

Prof. RNDr. J. Briančin, CSc., IGT SAS, Košice, Slovakia

A. Prof. Ing. B. Plešingerová, CSc., TU Košice, FM DC, Slovakia

A. Prof. Ing. M. Hnatko, PhD., IACH SAS, Bratislava, Slovakia

A. Prof. Ing. G. Sučik, PhD., TU Košice, FM DC, Slovakia

A. Prof. RNDr. A. Fedoročková, PhD., TU Košice, FM DCH, Slovakia

A. Prof. Dr. D. Hršak, PhD., MU Zagreb, Sisak, Croatia

Dr. Ing. D. Rohanová, ICT, Prague, FCT, IGC, Czech Republic

Ing. E. Grambálová, PhD. ,TU Košice, FM DC, Slovakia

Ing. Ľ. Popovič, PhD. TU Košice, FM DC, Slovakia

,

ISBN: 978-80-553-2122-6

Proceedings - XI

th International Conference

PREPARATION OF CERAMIC MATERIALS Herľany, June 9

t h – 11

t h, 2015

Editor: Organising Committee of International Conference PREPARATION OF CERAMIC MATERIALS 2015

Faculty of Metallurgy, Department of Ceramics Copyright Technical University of Kosice 2015 June 2015, The first edition , Number of Copies – 101; Number of Pages – 153.

Book reviewers/ Edited by:

Assoc Prof. RNDr František Lofaj, DrSc., Assoc. Prof. Ing. Miroslav Hnatko, Ph.D., Assoc. Prof. Ing. Beatrice Plešingerová, CSc., Assoc. Prof. Ing. Gabriel Sučik, Ph.D,, Ing. Diana Horkavcová, Ph.D., RNDr. Martin Fabián, Ph.D.,

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9th – 11th June 2015 Herľany, Slovakia 9th – 11th June 2015 Herľany, Slovakia


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