VSEBINA – CONTENTS IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES Thermal fatigue of a Ni-based superalloy single crystal Termi~na utrujenost monokristala iz nikljeve superzlitine L. Getsov, N. Dobina, A. Rybnikov . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67 Numerical determination of the carrying capacity of rolling rotational connections Numeri~na dolo~itev nosilnosti vtrljivih kotalnih zvez R. Kunc, A. @erovnik, M. @vokelj, I. Prebil . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73 Anisotropic hardening of materials by non-shearable particles Utrditev anizotropnih materialov z delci N. Bonfoh, S. Tiem, P. Lipinski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77 Analysis of the boronized layer on K 190 PM tool steel Analiza boronizirane plasti v orodnem jeklu K 190 PM, izdelanem po postopku metalurgije prahu M. Hudáková, M. Kusý, V. Sedlická, P. Grga~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81 A co-precipitation procedure for the synthesis of LSM material Soobarjanje LSM za pripravo katodnih materialov za gorivne celice M. Marin{ek, K. Zupan, T. Razpotnik, J. Ma~ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85 The effect of silica fume additions on the durability of portland cement mortars exposed to magnesium sulfate attack Vpliv dodatka silica fume na trajnost cementnih malt portland, izpostavljenih delovanju magnezijevega sulfata J. Zeli}, I. Radovanovi}, D. Jozi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91 Activated sintering of magnesium oxide obtained from seawater Aktivirano sintranje magnezijevega oksida, dobljenega iz morske vode V. Martinac, M. Labor, M. Miro{evi}-Anzulovi}, N. Petric . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95 Flowing of the melt through ceramic filters Pretok taline skozi kerami~ne filtre J. Ba`an, K. Stránský . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99 Sinteza magnetnih nanodelcev, funkcionaliziranih s tanko plastjo silike Synthesis of magnetic nanoparticles functionalized with thin layer of silica S. ^ampelj, D. Makovec, M. Bele, M. Drofenik, J. Jamnik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 41(2)65–108(2007) MATER. TEHNOL. LETNIK VOLUME 41 [TEV. NO. 2 STR. P. 65–108 LJUBLJANA SLOVENIJA MAR.-APR. 2007 L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL TERMI^NA UTRUJENOST MONOKRISTALA IZ NIKLJEVE SUPERZLITINE 1Leonid Getsov, 2Natalia Dobina, 2Alexandr Rybnikov 1Department of Strength of Materials, St. Petersburg State Politechnical University, St. Petersburg, Russia 2 NPO ZKTI, Russia guetsovonline.ru Prejem rokopisa – received: 2006-07-01; sprejem za objavo – accepted for publication: 2006-11-29 This paper presents a thermal-mechanical fatigue (TMF) test procedure and some results obtained in vacuum for a Ni-based superalloy single crystal. This procedure allows the monitoring of slip lines, fatigue-crack initiation and propagation until final failure. Single-crystal specimens with different crystallographic orientations were tested under different thermal-cyclic loading conditions. Tests were conducted on smooth specimens as well as on specimens with a central hole of 0.5 mm diameter to analyze the effects of stress concentration. Both crystallographic and non-crystallographic cracks were observed. An increase of the maximum temperature in the cycle (with the same amplitude of temperature change) causes a severe shortening of the service life, as well as stress concentration. It was observed that the first initiated microcrack arrests and the sample fracture takes place on newly nascent microcracks. Key words: thermal fatigue, single crystal, crystallographic orientation, crack, concentrator V ~lanku je opisana procedura za dolo~itev termi~nega utrujanja (TMF) in nekateri rezultati preizkusov monokristala iz nikljeve superzlitine v vakuumu. Procedura omogo~a opazovanje drsnih ~rt ter za~etek in propagacijo razpok do kon~nega preloma. Monokristali z razli~no prostorsko orientacijo so bili preizku{ani pri razli~nih pogojih termi~no cikli~ne obremenitve. Preizkusi so bili izvr{eni na gladkih preizku{ancih in na takih s centralno izvrtino s premerom 0,5 mm zaradi analize koncentracije napetosti. Opa`ene so bile kristalografske in rekristalografske razpoke. Povi{anje najvi{je temperature v ciklu (pri enakem razponu spremembe temperature) povzro~i veliko skraj{anje uporabne dobe, enak je vpliv koncentracije napetosti. Ugotovljeno je bilo, da se prva mikrorazpoka ustavi in da prelom nastane zaradi novo nastalih mikrorazpok. Klju~ne besede: termi~na utrujenost, monokristal, kristalografska orientacija, razpoka, koncentrator 1 INTRODUCTION Thermal-fatigue crack initiation in single-crystal aerofoils is the major cause of damage during turbine operation. A detailed study of fatigue-crack initiation and growth has been made 1-5. This paper describes a new thermal-cyclic vacuum test procedure. During vacuum testing, this procedure makes it possible to monitor the crystallographic slip orientations, crack initiation and growth. The publications 8-12 present a detailed review of research on the strength of single-crystal superalloys, carried out on the basis of an analysis of the literature with most of the attention paid to the publications of Russian researchers and, in particular, the authors of this paper. The problems of the influence of the anisotropy of single-crystal superalloys on the parameters of fatigue and the thermal-fatigue strengths are considered; some basic phenomenological and physical models of the material for the calculation of the deformation mode in single crystals are described and the criteria of thermal-fatigue fractures are discussed. Different aspects of the initiation and the evolution of the creep, fatigue and thermal fatigue cracks in single-crystal materials are considered, and their dependence on the temperature parameters, crystallographic orientation and their formation development mode was surveyed. Approaches to the definition of the fracture criteria and the methods of determining the deformation mode have been analyzed. Considerable attention is paid to the examination of the process of fatigue-crack development in single-crystals superalloys in the cases of HCF and LCF. The crystallographic peculiarities of fatigue failure and the criteria for the determination of the crack-growth direction and the rate at the different stages are considered. Special attention is paid to the conditions of the transition between stages. 2 EXPERIMENTAL PROCEDURE 2.1 Test procedure. A comprehensive procedure is developed in NPO CKTI for the determination of the thermal fatigue resistance of the various materials and coatings. A flat specimen is rigidly fixed in the vacuum chamber and cyclically heated with electric current (Figure 1). The heating is carried out according to the specified program (Figure 2) and automatically maintained during the testing. With the increase of the vacuum depth, the ultimate cycle temperature increases also. The temperature is measured with the rate of oxide-film formation on the sample surface. The slip Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 67 UDC/UDK 539.42:669.24 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)67(2007) lines and surface microcracks are observed with an optical microscope 6. The technique enables to: • Determine the thermo-fatigue resistance for gas-turbine blades based on crack initiation and propagation. • Compare of different material-processing techniques, parts manufacturing, repair and coating technologies by TMF resistance. • Perform fatigue tests of small-scale specimens cut out from the gas-turbine blades after operation. • To define the minimum thermal fatigue resistance of construction elements including composite materials, single-layer and multilayer coatings and welded joints. • To study the details of the mechanism of thermal fatigue for material for various cycle parameters. • To study the effect of stress concentration on the thermal fatigue. • To determine crack orientations (ratio of ∆KI/∆KII), the crack propagation rate, and the effect of the cycle’s parameters on the crack propagation and on the thermal fatigue service life. 2.2 With testing the following parameters are determined: characteristic properties of the deformation relief defining the mechanism of thermal-fatigue damage accumulation; the number of cycles to the first appearance of the microcrack in the speciment and in the coating; the growth rate of incipient cracks and the number of cycles to sample failure. The amplitude value of conventional elastic stresses is determined according to: ∆σ α α ϕ= −( )max minE T E Tst st1 1 2 2 ϕ = −1 ∆ ∆k l/ (1) where ∆k is the measured value of the displacement during the cycle of control micro-hardness marks, applied on the sample surface along its working part edges; Est is a static module of elasticity; and l is the free movement of the control marks during the heating from Tmin to Tmax. The stress-strain relationship and the inelastic deformation evolution are calculated using a structural model 7. The specific feature of this procedure and those of Coffin are the constant stresses over the sample’s cross-section. The number of cycles to failure of flat samples in vacuum and of tube samples in air that were tested using the different test units of Coffin, under thermal cycle loading, turned out to be identical for a number of heat-resistant nickel-based alloys. In this project tests were carried out under permanent visual surface observation and periodic photographing of the slip lines’ orientation, the surface microcracks’ initiation and the changes of the slip lines’ orientation during thermal cycling with the optical magnification of 260 times. It should be noted that photographing during the course of the tests makes it possible to obtain high-quality images. TMF tests were conducted until failure or 2000 cycles, depending on which event happened first. After the tests the fracture surfaces were analyzed using optical fractography and a magnification of 160 times. Some tests were conducted on the samples with stress concentrators in the form of 0.5 mm holes. 2.3 Materials. The plates of the monocrystal alloy underwent the standard thermal treatment: heating in vacuum up to 1320 °C, holding at this temperature for 6 h, and cooling in argon. This was followed by double-stage aging in the mode: first stage, heating in air up to 1030 °C, holding at this temperature for 4 h; second stage, 24 h at 870 °C. The as-received crystal orientation is listed in Table 1. Table1: Crystallographic orientation of single-crystal plates Tabela 1: Kristalografska orientacija monokristalnih plo{~ Number of the sample series The crystallo- graphic orien- tation of the samples along the longer axis Deviation from exact orientation, degree Azimuthal orientation, degree 1 111 5.64 8.26 2 011 4.51 11.27 3 011 8.33 14.43 4 011 9.67 7.86 5 001 5.47 41.97 3 RESULTS AND DISCUSSION Despite the limited number of tests, the obtained experimental data allow us to formulate some L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL 68 Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 Figure 2: Program of the temperature-cycle variation Slika 2: Program temperaturnega cikla Figure 1: Specimen shape and size Slika 1: Oblika in mere preizku{anca relationships which, of course, require further experi- mental verification. 3.1 Directions of the slip lines. The analysis of the test data shows that for different crystallographic orientations the slip lines form under different angles, including those oriented approximately across and along the loading vector, and also at an angle close to 45°. We obtained a reasonably good convergence between the experimentally measured and predicted slip lines. For a small number of cycles the nascent slip lines are oriented in the directions corresponding to the greatest Schmid factor. It should be noted that in the samples of series 3, 4 and 5 (Table 2) the slip lines corresponding to both, the octahedral and cubic slip systems, were observed (Figure 3), and in the samples of series 1 and 2, only those corresponding to the octahedral system were seen (Figure 4). Multiple slip and even patchy appears with an increase in the number of cycles. Microcracks initiate mainly on slip lines, but the direction of their growth does not depend on the direction of the slip lines, therefore, it has a non-crystallographic nature. Thus, initially, the microcrack is parallel to the slip lines, and then the crack turns out to be oriented along the sample axis (maximum compression). It is worth noting that no slip lines were found in this direction. Also, non- crystallographic directions of microcracks were observed in the micrographs of the fracture zones of all the samples (see Figure 5). 3.2 Orientation of the fracture surface The character of the sample fracture can be seen in the scheme of Figure 6 and the picture in Figure 7. The directions of the distribution of cracks were compared to the directions of the sliding lines. The samples 3-0, 3-1, 3-3 of the series 3 011 broke mainly in a zigzag line in the lattice plane, the samples 3-2, 3-4, 3-5 ruptured in the plane. At the upper edge of the section, the samples 3-0 and 3-5 were damaged by the rupture and the samples 3-1, 3-2, 3-3, 3-4 were damaged below the edge. L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 69 Table 2: Comparison of the calculated and experimental directions of the slip lines Tabela 2: Primerjava izra~unanih in eksperimentalnih smeri drsnih ~rt Sample number Exp. slip line slope deg. Theor. slip line slope octahedron cubic 2-1 10 80; 82.5; 76.8; 78; 78.2; 82; 83; 84.2 93 6.1 81 95.8 - - - 2-2 75.8; 80.5;88.8; 86 99 114;126 81 - 95.8 - - - - 132.8 2-4 0; 1.3 84.5 0.2 81 - - 2-6 4.2; 3.8 79.8; 81.2; 82; 85.8 93.5; 95.5 6.1 81 95.8 - - - 5-1 93.2 90.2 95.5 5-2 3.5 92.6 107.1 - 90.2 100.2 5.3; 5.6 95.5 - 5-3 39.4 90.6 149 40.8 90.2 150.2 - - - 1-6 17.8 131.2 16.5 130 - - 1-2 22.4; 20.8 130.5 16.5 130 - - 1-1 15.8 83.2; 82.5 128.5; 130.8 16.5 87.5 130 - - - 3-0 15 85 170.174 - 93.6 175.7 14.8 - 3-1 70.5 89 74.8 93.6 3-2 35 71.5 93.2 - 74.8 93.6 38.6 3-3 91; 95.5 169.2 93.6 167.5 3-4 21; 22.5 65.5 93.2; 92.5; 92 - 74.8 93.6 14.8 4-1 50 78.5 155; 154; 162.4; 158.5 48.8 163.2 - 83.8; 84.9 - 4-2 48.1; 51.2; 49.2; 48.5 86.5; 81.5 48.8 - - 83.8; 84.9 Figure 3: Microstructure of the surface of specimen 3-2 after 63 cycles (250–1000 °C) Slika 3: Mikrostruktura povr{ine preizku{anca 3-2 po 63 ciklih (250–1000 °C) The samples of series 5 001 broke perpendicularly to the loading vector along the lattice plane and at the angle toward the specimen plane (shear). The samples 2-4 and 2-6, with concentrator of series 2 011, broke perpendicularly to the loading vector along the plane and on the specimen plane. The sample 2-1 without a concentrator broke by shear in the plane and on the section. The sample 4-1 of series 4 (with mixed orientation) with the concentrator was fractured perpendicularly to the loading vector, and the smooth sample 4-2 cracked at an angle toward the specimen axis and plane (shear). The samples of series 1 111 with and without a concentrator cracked by shear with a change of direction in the surface and in the section. L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL 70 Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 Figure 6: Appearance of the fractured specimen 3-0 Slika 6: Prelomljen preizku{anec 3-0 Figure 5: Microstructure in the vicinity of the fracture of specimen 3-2 Slika 5: Mikrostruktura v bli`ini preloma vzorca 3-2 Figure 4: Microstructure of the fractured specimen 5-3 Slika 4: Mikrostruktura prelomljenega preizku{anca 5-3 Table 3: Influence of the maximal temperature of a cycle on the number of total cycles to failure Tabela 3: Vpliv najvi{je temperature v ciklu na {tevilo ciklov do preloma Sample number Orientation ∆T Tmin/°C Tmax/°C N 3-0 011 750 150 900 951 3-1 200 950 450 3-2 250 1000 63 3-5 500 450 950 2535 3-3 500 1000 1220 3-4 550 1050 356 5-1 001 750 150 900 560 5-2 250 1000 95 2-4 with stress concentrator 011 500 350 850 2952 2-6 with stress concentrator 500 1000 187 Figure 7: Course of the fracture lines in the specimens Slika 7: Potek razpok v preizku{ancih 3.3 Influence of the maximum temperature in the cycle The data in Table 3 show that the increase of the maximum temperature within the ranges 850–1000 °C and 900–1000 °C with the same total temperature amplitude decreases the service life to fracture by ≈15 times. If the maximal temperature is increased from 1000 °C up to 1050 °C with the same total temperature amplitude the time-to-fracture is decreased by approximately a factor of 4. 3.4 Influence of the total temperature amplitude A decrease of the total temperature amplitude with a constant value of TMAX = 1000 °C increases considerably the service life (Table 4). Table 4: Influence of the temperature range on the number of cycles to fracture Tabela 4: Vpliv razpona temperature v ciklu na {tevilo ciklov do preloma Sample number Orientation Tmin/°C Tmax/°C N 3-2 011 250 1000 63 3-3 011 500 1000 1220 5-2 001 250 1000 95 5-3 001 500 1000 1460 3.5. Influence of the stress concentrator Concentrators in samples in the form of holes with a diameter of 0.5 mm decrease the service life by ≈15 times (Table 5). At the same time, during the test of the sample with the concentrator and Tmax = 850 °C a relatively large number of cycles before fracture (2952) was recorded. The sample without concentrator was not tested in such conditions, since the expected number of cycles before the fracture occurred was more than 10000. Table 5: Influence of the stress concentrator on the number of cycles before destruction. Tabela 5: Vpliv koncentratorja napetosti na {tevilo ciklov do preloma Sample number 4-2 4-1 with stress concentr ator 1-2 1-1 with stress concentr ator 2-2 2-6 with stress concentr ator Orientation 011 011 011 011 011 011 Tmin/°C 150 150 150 150 500 500 Tmax/°C 900 900 900 900 1000 1000 N 308 25 823 50 472 187 3.6 Influence of the crystallographic orientation of material The maximal thermal fatigue life in tests under the same thermal cycle modes was observed in the case of smooth samples with a 111 orientation and the minimal strength was obtained for the 011 direction (Table 6). Samples with a mixed orientation and 001 broke after a number of cycles decreasing in the series N111 > N001 > Nmix > N011. At the same time, the notched samples with orientation 011 had a longer life than the samples with the orientation 111. The reasons for such behavior require further research. The influence of the azimuthal orientation of single-crystal alloys 011 on the number of cycles before destruction is illustrated by the data in Table 7. Table 6: Influence of orientation on the number of cycles before destruction. Tabela 6: Vpliv orietacije na {tevilo ciklov do preloma Sample number Orientation Tmin /°C Tmax/°C N 2-1 011 150 900 100 4-2 011 308 1-2 111 823 5-1 001 560 3-2 011 250 1000 63 5-2 001 95 2-2 011 500 1000 472 3-3 011 1220 5-3 001 1460 2-4 with stress concentrator 011 350 850 2952 1-6 with stress concentrator 111 320 Table 7: Influence of the azimuthal orientation of single-crystal alloys 011 on the number of cycles before destruction (Tmin = 150 °C, Tmax = 900 °C ) Tabela 7:Vpliv azimutne orientacije monokristala 011 na {tevilo ciklov do preloma (Tmin = 150 °C, Tmax = 900 °C) Sample number A deviation from exact orientation, degree Azimuthal orientation, degree N 2-1 4.51 11.27 100 3-0 8.33 14.43 951 4-2 9.67 7.86 308 L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 71 Figure 8: Microstructure of the surface of specimen 3-0 after 100 cycles (150–900 °C) Slika 8. Mikrostruktura povr{ine preizku{anca 3-0 po 100 ciklih (150–900 °C) On of the samples the initiation of several micro- cracks, was observed often (but not always) oriented along the slip lines; some of the cracks originated near the pores. However, the fracture, as a rule, occurred from cracks located at some distance from the first micro- crack. 3.7 Influence of the recrystallization The recrystallization on sample 3-0 with the 011 crystallographic orientation hindered the development of slip lines. The microcrack initiation was observed on the sample surface beside the recrystallization areas (Figure 8). Nevertheless, these microcracks did not evolve into macrocracks and the sample fracture had a complex trajectory after a large number of cycles. 4 CONCLUSIONS Within the framework of this project ingot blanks of single-crystal superalloys with five different crystallo- graphic orientations were investigated. A new testing technique was developed, allowing obtain data on the role of the maximal and minimal cycle temperature, the stress concentration and the crystallographic orientation on the formation of slip lines, of microcracks and their propagation. On the basis of 18 tests with a maximum number of cycles equal to 2952 in the temperature range Tmin = 150–500 °C, Tmax = 850–1050 °C it was found that the thermal fatigue-crack initiation and growth are strongly influenced by the crystallographic orientation. Nevertheless, the fracture proceeds in a non-crystallo- graphic mode and in some cases the cracks change their direction of propagation. The following conclusions are proposed: 1. It was observed that the sample fracture takes place, depending on the orientation and the test mode, both in crystallographic and non-crystallographic directions independently of the crystallographic slipping. 2. Different directions of slip in different sections of the sample surface were discovered, including multiple slip. 3. A sharp reduction in fatigue life was observed with an increase in the maximum temperature in the cycle (with the same amplitude of temperature change). 4. The stress concentrators decrease significantly the fatigue life. 5. It was observed that the initial microcrack decelerates and stops and the sample fracture takes place on a newly nascent microcrack. In the authors’ opinion it would be useful to carry out further research, focusing on the detailed study of the fractured samples and thermal fatigue tests, keeping the samples at the maximum temperature for the same time interval in each cycle. Further study of the problem of the thermal fatigue strengths of the blades should concentrate on the development of computational methods for the deformation mode of single-crystal blades by creep and cyclic change of temperature, on the estimation of the blades’ strength for thermal-cycle loading and on the development of a method for calculating the crack-growth rate, taking into account creep, stress concentration and fracture-mode changes. Acknowledgement The project was funded by the Pratt & whitney UTC Company. Special thanks are also due to Dr. Alexander Staroselsky, for problem formulation and program management. 5 REFERENCES 1 P. A. S. Reed, X. D. Wu, Sinclair Fatigue crack path prediction in UDIMET 720 nickel-based alloy single crystals. Metallurgical and Materials. 31A (2000), 109–120 2 J. Telesman , L. J. Ghosn. Fatigue crack growth behavior of a PWA 1484 single crystal superalloy at elevated temperatures. ASME Paper 95-GT-452, 1995 3 K. S. Chan, J. Feiger, Y.-D. Lee, R. John, S. J. Hudak. Fatigue crack growth thresholds of deflected mixed-mode cracks in PWA1484, Journal of Engineering Materials and Technology. 127 (2005), 2–7 4 S. X. Li, D. J. Smith, An overview of combined fatigue and creep response of single crystal nickel base superalloys. Proc. 5-th Liege Conf. on Materials for Advanced Power Engineering, Part II, October, 1974, Belgium, 1175–1184 5 N. Marchal, S. Forest, L. Remy, S. Duvinage. Simulation of fatigue crack growth in single crystal superalloy using local approach to fracture. Euromech- mecamat 2006, Local approach to fracture, 9-12 May 2006 6 A. I. Rybnikov, L. B. Getsov. New technique and results of ther- mal fatigue tests of superalloys and coatings. Proceedings of the sixth International Congress on Thermal Stresses. Vienna, Austria, 1 (2005), 305–309 7 L. G. Padva, L. B. Getsov, O. S. Sadakov, A. I. Rybnikov Design method of strength estimation for gas turbine blades with coatings, Journal of Machinery Manufacture Under Reliability Problem, 4 (1992), 57–63 8 Golubovskiy E., Svetlov I., Nozhnitsky Yu. Relationship of stress rupture and crystallographic orientation for Ni-base superalloys single crystal. EUCASS. European Conference for Aerospace Sciences. July 4-7, 2005, Moskow, Russia, CD 9 Rtishchev V. V. Anisotropic alloys with columnar and single crystal structures used for blades of stationary gas turbine plants. Proc. 5th Liege Conf. on Materials for Advanced Power Engineering, Part II, October, 1994, Belgium, 1135–1144 10 R. E. Shalin, I. L. Svetlov, E. B. Kachanov et al.: Single crystal nickel base superalloys, Mashinbuilding, (1997) 333 (In Russian) 11 Erickson G. L., Harris K. Ds and sx superalloys for industrial gas turbines, Proc. 5th Liege Conf. on Materials for Advanced Power Engineering, Part II, October, 1994, Belgium, 1055–1074 12 Jo C.-Y., Kim H.-M. Effect of recrystallisation on microstructural evolution and mechanical properties of single crystal nickel based superalloy CMSX-2 Part 2 – Creep behaviour of surface recry- stallised single crystal., 19 (2003) 12, 1671–1676 L. GETSOV ET AL.: THERMAL FATIGUE OF A Ni-BASED SUPERALLOY SINGLE CRYSTAL 72 Materiali in tehnologije / Materials and technology 41 (2007) 2, 67–72 R. KUNC ET AL.: NUMERICAL DETERMINATION OF THE CARRYING CAPACITY ... NUMERICAL DETERMINATION OF THE CARRYING CAPACITY OF ROLLING ROTATIONAL CONNECTIONS NUMERI^NA DOLO^ITEV NOSILNOSTI VTRLJIVIH KOTALNIH ZVEZ Robert Kunc, Andrej @erovnik, Matej @vokelj, Ivan Prebil Faculty of Mechanical Engineering, University of Ljubljana, A{ker~eva ul. 6 1000 Ljubljana, Slovenia robert.kuncfs.uni-lj.si Prejem rokopisa – received: 2006-05-17; sprejem za objavo – accepted for publication: 2007-02-16 The exploitation of large axial bearings includes load peaks, which cause permanent deformation of the rolling contact. The plastic strain of the base material under the hardened rolling layer starts to grow and micro-cracks on the edge of the hardened layer and the peeling of this layer may occur. In our computation we have used an elasto-plastic model that combines isotropic and kinematic hardening and the growth of material damage. The damage model allows us to follow the variation of the elastic and plastic stress and deformation as a function of the number of cycles. In the article, an experimental verification of the described numerical model is shown, which can be used to determine the actual carrying capacity of the rolling contact for low-speed axial bearings. Key words: rotational connections, low cycle fatigue, cyclic plasticity, damage Pri uporabi velikih aksialnih le`ajev se pojavljajo preobremenitve, ki povzro~ijo trajno deformacijo le`ajnega kontakta. Plasti~na deformacija osnovnega materiala pod utrjeno le`ajno povr{ino za~ne rasti in lahko nastanejo mikrorazpoke na robu kaljene plasti ter njeno lu{~enje. Pri na{em izra~unu smo uporabili elastoplasti~ni model, ki kombinira izotropno in kinematsko utrjevanje ter rast po{kodovanosti materiala. Ta model omogo~a sledenje spremembe elastoplasti~ne napetosti in deformacije kot funkcijo {tevila ciklov. V ~lanku je opisana eksperimentalna verifikacija opisanega numeri~nega modela, ki ga lahko uporabimo za dolo~itev prave nosilnosti le`ajnega spoja za aksialne le`aje, ki delujejo pri majhni hitrosti. Klju~ne besede: vrtljiva zveza, malocikli~na utrujenost, cikli~na plasti~nost, po{kodbe 1 INTRODUCTION Bearings are among the most frequently used elements in machine engineering. Owing to their wide- spread use, the requirements that a bearing must meet are highly diverse. Bearing use ranges from applications in which a bearing collapse does not constitute a major problem, to applications where the collapse of a bearing could lead to enormous economic losses and potentially disastrous consequences for people. The large bearings used in rolling rotational connections are an example of such an application. The basic element of a rotational connection is a large rolling bearing, with attachment holes in the bearing rings, and a gear wheel (Figure 1). Alloyed steels such as 42 CrMo 4 and C 45 – ISO 683/1 are normally used in the production of bearing rings since the process requires thermal and mechanical treatments. The raceway is surface-hardened to attain the minimum hardness necessary to prevent the pressing-in of the rolling element. The actual external load and the elasticity of the upper and support structures cause an uneven distribution of the external loads over the rolling bearing’s diameter 1-3. This leads to local plastic deformations of the bearing raceway and to the initiation of damage, largely conditioned by the load size and material fatigue 4-5. The existing models for determining the carrying capacity of large rolling-bearing conne- ctions with surface-hardened raceways do not fully consider or even greatly simplify the actual state of the bearing ring 3-12. In analyzing the actual carrying capacity of rolling contacts in large rolling bearings with surface-hardened raceways we have decided to use a combined elasto- plastic constitutive model, which links the material Materiali in tehnologije / Materials and technology 41 (2007) 2, 73–76 73 UDC/UDK 621.8 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)73(2007) Figure 1: Elements of rotational connections Sllika 1: Elementi rotacijskih povezav damage mechanics with the isotropic and kinematic hardening/softening 13-15. The material model considers 17 material parameters, selected from a series of monotone and cyclic experimental tests. The model, which was built in the finite-element code, was used to investigate the development of elasto-plastic deformations, stress areas and damage occurrence for low-cycle loads that occur in the raceways of large rolling bearings at a low rotational speed. In the article an experimental verification of the described numerical model is shown, which can be used for the determination of the actual carrying capacity of the rolling contact in low-speed axial bearings. 2 NUMERICAL MODEL A developed and built-in finite element is used in the application of a numerical model focusing on the rolling contact between a raceway model and a rolling element, and is carried out in the ELFEN© commercial software environment. The discretisation of a bearing raceway’s geometric model applies a developed finite element that covers isotropic hardening/softening, kinematic hardening and material damage 13-15. This makes it possible to monitor the actual low-cycle development of the strain deformation or the hardening/softening and material damage growth. The end code of a developed element is written in the FORTRAN programming language. The geometry of the contact between the rolling element and the raceway is determined with square surface-contact elements (Figure 2). The numerical model of a rolling contact considers the actual geometrical and material characteristics of bearing raceway models. The characteristics of a bearing ring model change in line with the raceway depth. After having measured the changes in hardness along the cross-section of a bearing raceway model, we entered various material parameters into the damage model. These parameters were derived from the results obtained from single-axis monotone and cyclic tear tests 13. In describing the material characteristics of raceway models, we also took into account the remaining material stresses that occurred during the production of the test raceways 14. 3 NUMERICAL RESULTS The numerical rolling-contact model was used to simulate a contact of the rolling element and a surface-hardened bearing raceway at a varying contact force from Fmax = 35.2 kN to Fmin = 1 kN. The model enables every element to be assigned its deformation strain state and growth of damage relative to the number of contact load cycles. Figure 3 shows the evolution of the equivalent von Mises stresses σVM and strains εVM at the location of maximum-damage accumulation for the first 40 loading cycles, and Figure 4 shows a distribution R. KUNC ET AL.: NUMERICAL DETERMINATION OF THE CARRYING CAPACITY ... 74 Materiali in tehnologije / Materials and technology 41 (2007) 2, 73–76 Figure 4: Damage distribution in 3D model after 100 load cycles Slika 4: Porazdelitev po{kodbe v 3D-modelu po 100 nihajih Figure 2: Numerical model Slika 2: Numeri~ni model N σ V M 2 ε V M Figure 3: Development of the stress σVM and strain εVM paths on the edge of the raceway-model hardened layer Slika 3: Razvoj napetosti σVM in deformacije εVM na robu utrjene plasti le`ajne te~ine of the damage after 100 contact force cycles. In contrast, Figure 5 shows the damage distribution along the ring depth that determines the damage accumulations in the base material on the edge of a hardened layer. The carrying capacity of the raceway model for rolling elements of different diameters was determined on the basis of a diagram showing the growth of material damage relative to the number of load cycles (Figure 5)14. 4 VERIFICATION OF THE MODEL A comparison of the results of the experimental work and the numerical calculations of the contact between the raceway model and the rolling element can only be made by comparing the directions of the maximum total and permanent displacements of the bearing-raceway model. The comparison shows an even development of the raceway-model displacements and thus supports the use of the proposed mechanical and mathematical models to determine the carrying capacity of large rolling bearings. There is a close match for the initial values of the measurements and the numerical calculations of the total and permanent displacement of the bearing raceway (Figure 6). The matching is also evident from the comparison of the results obtained from measurements and numerical calculations of the first displacement-contact force hysteresis loop (Figure 7). Even when they are subjected to different loads and different geometric contact models, the measured and the calculated results of the total and permanent raceway displacements, which depend on the number of cycles, match closely their initial value and the deformation growth gradient 14. 5 CONCLUSION The determination of the low-cycle carrying capacity of a large bearing raceway with the numerical model material behavior, which is applied with the finite-element method, is very effective. The model takes into account the geometric and construction characte- ristics of the bearing rings, the raceways, and the rolling elements as well as the material characteristics of the surface-hardened raceways, which link the continuum damage dynamics with the isotropic and kinematic hardening/softening. With regard to the expected loads and the geometric limitations of the rotational connection, both the geometry of the contact between the raceway and the rolling element as well as the heat treatment of the bearing rings can be optimized. We can also assess the life expectancy of a rotational connection, i.e., the time and location of the initiation of the macro-cracks on the surface-hardened bearing raceway. Further investigations into the carrying capacity of a surface-hardened bearing raceway are focused on the R. KUNC ET AL.: NUMERICAL DETERMINATION OF THE CARRYING CAPACITY ... Materiali in tehnologije / Materials and technology 41 (2007) 2, 73–76 75 0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0 5 10 15 20 25 30 35 C o n ta ct L o a d F kN Contact Displacement, δ / / mm exp. simul. cycle 1 and 2 cycle 100 Figure 7: Comparison of the development of the measured and calculated displacement-contact force hysteresis of the raceway-model damage distribution Sllika 7: Primerjava razvoja izmerjene in izra~unane histereze v kontaktni sili za porazdelitev po{kodbe le`ajne te~ine 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 0.000 0.002 0.004 0.006 0.008 0.010 0.012 Depth, /mmD cycle 1 cycle 10 cycle 50 cycle 100 D a m a g e D Figure 5: Damage distribution in ring depth Slika 5: Porazdelitev po{kodbe v globino le`ajnega obro~a δ t δ δ δ δ Figure 6: Comparison of the development of measured and calculated total and permanent displacements of the raceway-model damage distribution Sllika 6: Primerjava razvoja izmerjenega in izra~unanega skupnega in permanentnega pomika za porazdelitev po{kodbe le`ajne te~ine possibility of upgrading the material model with a rheological model for a description of the natural limits of the material yield, which occurs in the normalized and optimized bearing-ring base material, and with a model to describe the growth of a crack and the life expectancy of heat-treated steels used for bearing raceways. 6 REFERENCES 1 J. Brandlein, Lastubertragung durch Grosswalzlager bei elastichen Ringtragern als Unter- und Oberkonstruktion, f+h – Fordern und Heben, 30, (1980) 2 I. Prebil, S. Zupan, P. Lu~i~, Lastverteilung auf Wälzkörper der Drehverbindungen. Konstruktion, 47 (1995), 339–345 3 E. V. Zaretsky, STLE Life Factors for Rolling Bearings, Chapter 4 – Bearing Load, STLE Publications, Illinois, 1992 4 T. A. Harris, Rolling bearing analysis – 3rd edition, John Willey & Sons, New York, 1991 5 R. A. Pallini, J. E. Sague, Computing core-yeild limits for case- haredened rolling bearings, ASLE Trans., 28 (1985) 1, 91–96 6 J. E. Sague, The special way big bearings can fail, Machine Design, 50 (1978) 21, 113–120 7 P. C. Bastias, Analysis of cyclic crack growth under rolling contact loading condition, Vanderbilt University, Dissertation, 1989 8 ISO 281, Rolling bearings – Dynamic load ratings and rating life, 1990 9 C. P. Jones, W. R. Tyfour, J.H. Beynon, A. Kapoor, The effect of strain hardening on shakedown limits of pearlitic steel, Proc. Instr. Mech. Engrs., 211 (1997), 131–140 10 G. Lundberg, H. Sjovall, Stresses and deformation in elastic contacts, Chalmers University of Technology, Gothenburg, 1958 11 H. R. Thomas, V. A. Hoersch, Stresses due to the pressure of one elastic solid upon another, University of Illinois, 212 (1930) 12 G. Rydholm, On inequalities and shakedown in contact problems, Department of Mechanical Engineering, Linkoping University, 1981 13 R. Kunc, I. Prebil, Low-cycle fatigue properties of steel 42CrMo4, Mat. Sci. Eng. A, 345 (2003), 278–285 14 R. Kunc, Low cycle carrying capacity for bearing raceway with hardened rolling surface, Ph.D. Thesis, University of Ljubljana, 2002 15 R. Kunc, I. Prebil, Numerical determination of carrying capacity of large rolling bearings. J. Mater. Process. Technol., 155/156 (2004), 1696–1703 R. KUNC ET AL.: NUMERICAL DETERMINATION OF THE CARRYING CAPACITY ... 76 Materiali in tehnologije / Materials and technology 41 (2007) 2, 73–76 N. BONFOH, P. LIPINSKI: ANISOTROPIC HARDENING OF MATERIALS BY NON-SHEARABLE PARTICLES ANISOTROPIC HARDENING OF MATERIALS BY NON-SHEARABLE PARTICLES UTRDITEV ANIZOTROPNIH MATERIALOV Z DELCI 1Napo Bonfoh, 2Sonnou Tiem, 1Paul Lipinski 1Laboratoire de Fiabilité Mécanique (L.F.M.), Ecole Nationale d’Ingénieurs de Metz (E.N.I.M.), Ile du Saulcy, F 57045 Metz Cedex 01, France 2Ecole Nationale Supérieure d'Ingénieurs (E.N.S.I.), Université de Lomé, B.P. 1515 Lomé, Togo bonfohenim.fr; sonnou.tiemtogotel.net.tg; lipinskienim.fr Prejem rokopisa – received: 2006-05-17; sprejem za objavo – accepted for publication: 2006-11-07 A new description of the hardening of polycrystalline materials is proposed. This hardening is mainly induced by non-shearable particles. Anisotropic and kinematic hardening of this kind of material has been mentioned in a previous paper (Bonfoh et al., 2003). In contrast to this initial investigation, the proposed modified formulation introduces a non-linear hardening for the single crystal without particles and better renders the material behaviour at finite deformations. The micromechanical modelling is based on a two-level homogenization approach: a micro-mesoscopic transition is firstly performed to derive the equivalent behaviour of a single crystal containing a volume fraction of intra-crystalline particles. Then, a second transition to the macroscopic scale leads to the elastic-plastic behaviour of the polycrystalline material. The numerical results of the model are analysed in terms of the internal stress developed during plastic loading. The material behaviour, when subjected to a multiaxial loading, is also studied through the concept of the yield surface. Some other aspects, such as the influence of particles size, shape and volume fraction, are also investigated. Key words: crystal plasticity, inhomogeneous material, residual stress, yield surface, kinematic hardening Predlo`en je nov opis utrditve polikristalnih materialov. To utrditev povzro~ajo nestri`ni delci. V prej{njem delu je bila opisana anizotropi~na in kinemati~na utrditev materiala te vrste (Bonfoh in dr., 2003). Nasprotno od te za~etne raziskave predlo`ena modificirana formulacija uvaja nelinearno mikromehansko modeliranje za nanokristal brez delcev in bolje opisuje vedenje materiala pri kon~ni deformaciji. Mikromehansko modeliranje temelji na pribli`ku dvostopenjske homogenizacije: mikro-mezoskopska tranzicija je uporabljena za opis ekvivalentnega vedenja monokristala, ki vsebuje neki prostorninski dele` intrakristalnih delcev, nato pa tranzicija na makroskopsko merilo vodi v elasto- plasti~no vedenje polikristalnega materiala. Numeri~ni rezultati modela so analizirani s stali{~a notranjih napetosti, ki nastanejo pri plasti~ni obremenitvi. Vedenje materiala pri ve~osni obremenitvi je analizirano s stali{~a plasti~nega popu{~anja povr{ine. Analizirani so tudi drugi vidiki, npr. vpliv velikosti oblike in prostorninskega dele`a delcev. Klju~ne besede: plasti~nost kristala, nehomogen material, zaostale napetosti, plasti~no popu{~anje povr{ine, kinemati~na utrditev 1 INTRODUCTION Most polycrystalline materials exhibit, at the microscopic level, some heterogeneities, such as foreign atoms or particles. These heterogeneities, which may appear in the form of a solid solution or second-phase particles, may be located either between the crystals or inside the single crystals of the polycrystalline materials. In the present study, only the latter heterogeneities are considered. These intra-crystalline particles, such as precipitates, interact with moving dislocations during the material’s plastic flow and lead to a modification of the single-crystal hardening. Some papers have been devoted to a description of the influence of intra-granular heterogeneities on the behaviour of polycrystalline materials: Schmitt et al (1996), Bochet et al (2001), Barbe et al. (2001) and Han et al. (2004). Recently, Reza et al (in press) has proposed a description of crystalline materials containing non- interacting elastic particles. Performed within a small strain formulation and confined to uniform elasticity, the authors propose a calculation of the additional work of deformation (strain energy) arising from plastic strain-field incompatibilities between the elastic-plastic matrix and the elastic particles. In this work, a micromechanical description of a single crystal containing particles is developed through a micro-meso transition. First, a modified Schmidt’s law and a new hardening matrix, taking into account the usual dislocation-dislocation interactions and also the interactions between dislocations and particles, are proposed. Next, a meso-macro transition using the self-consistent method developed by Lipinski et al. (1993, 1998) is applied to deduce the global response of the polycrystalline material. A comparison between the classical self-consistent predictions and the new approach is made. Special attention has been focused on the prediction of the yield surfaces and the residual stresses. 2 DESCRIPTION OF THE SINGLE CRYSTAL WITH PARTICLES In this paper, single crystals are assumed to contain a certain volume fraction of ellipsoidal particles. Moreover, at the microscopic level, the single crystal is Materiali in tehnologije / Materials and technology 41 (2007) 2, 77–80 77 UDC/UDK 621.7.016.3:539.2 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)77(2007) considered as a two-phase material: the elastic-plastic crystalline matrix with elastic particles. The plastic straining of the crystal results from the movements of dislocations on geometrically well-defined slip planes. The classical theory of single-crystal plasticity is adopted and Schmidt’s law is valid. At the crystalline level, the presence of non-shearable particles results in some incompatibilities of the strain field between the crystalline matrix and the particles. These incompatibilities are mainly due to the anisotropy of elastic behaviour and the heterogeneity of plastic straining between the matrix and the particles. The stress rate inside the crystalline matrix is then given by: :σ σ εm m= +1 P (1) The last term P:εm , represents internal stresses arising from these incompatibilities, where the fourth- rank tensor P is the polarisation tensor describing the evident interactions between the two phases evolving inside the heterogeneous grain. The expression of this tensor P can be derived using classical models of the interactions problems of heterogeneous materials (Kröner, Mori-Tanaka, Self-consistent, etc.). In this study, the last of these is used, and this leads to: P I b I= − − −f fãI( ): ( )3 3 1K (2) where ã 3 and b3 are, respectively, the strain- and stress- rate localisation tensors inside the particles, KI is the tangent elastic-plastic moduli of the equivalent single crystal with particles. Moreover, assuming the following relationship for the hardening of the single crystal without particles: τ γcr g gh h N hH= = ∑ 1 (3) It is also possible to derive a kind of Schmidt law for the equivalent heterogeneous grain: [ ]R f fg m I: ( ) : ( ) : 1 3− + −I P I bû σ = = ( : : ):( )H R R fgh g h h N h− − = ∑ P 1 1 γ (4) The first term appears as the resolved shear-stress rate for the equivalent grain I subjected to the σ I stress rate, where the second one reveals a new form of hardening matrix: ℵ = −gh gh g hH R R: :P (5) taking into account the initial dislocation-dislocation interactions through the matrix Hgh but also the inte- ractions between the particles and the dislocations during their movements ( : : )R Rg hP . The description of the equivalent behaviour of the single crystal with particles is therefore completed by the expression of its tangent elastic-plastic properties: K I I I g gh h I h N g N R K R= − == ∑∑7 7 7: : :U 11 (6) U I P I b= − + −( ) : :( )1 3f fmû and K R Rgh gh g I h= ℵ + −( : : : )U 7 1 (7) The superscripts m, 3 and I are, respectively, related to the matrix (single crystal without particle), particles and the considered grain I. After this first transition from local microscopic properties to the mesoscopic level, leading to the grain behaviour, a second transition, to the macroscopic level, is then performed to derive the equivalent properties of the polycrystalline material consisting of a large number of single crystals with particles. 3 NUMERICAL SIMULATIONS The proposed model is applied to the simulation of the material’s behaviour. The selected material is a SiC particulate-reinforced 5456 aluminium alloy matrix com- posite: • Matrix material (aluminium): Young’s modulus Em = 73 GPa, Poisson’s ratio m = 0.33, and initial uniaxial yield stress y = 230 MPa. • Elastic properties of SiC particles: Young’s modulus Ep = 485 GPa and Poisson’s ratio p = 0.2. 3.1 Material behaviour under monotonous loading The selected material is subjected to a uniaxial tension in direction 1, and its macroscopic response in terms of stress–strain is depicted in Figure 1. The plotted curves exhibit an evident hardening of the macroscopic behaviour with the volume fraction of particles, since the particles considered here are harder than the crystalline matrix. Moreover, this hardening is characterized by an increasing of the equivalent elastic threshold and a non-linear evolution in its elastic-plastic part. Since the initial hardening of the single crystal is assumed linear, this non-linear aspect results from the original amendments introduced in the proposed hardening model. N. BONFOH, P. LIPINSKI: ANISOTROPIC HARDENING OF MATERIALS BY NON-SHEARABLE PARTICLES 78 Materiali in tehnologije / Materials and technology 41 (2007) 2, 77–80 Figure 1: Macroscopic response for different volume fractions of particles Slika 1: Makroskopski odgovor za razli~en volumenski dele` delcev In the present investigation, particles are non-shear- able and during the plastic flow, dislocations bypass these obstacles with the Orowan-looping mechanism. The initial critical shear stress on slip systems is influenced by particle interactions at the beginning of the plastic flow. 3.2 Yield surface The material behaviour during multiaxial and non-proportional loads is analysed through the concept of a yield surface depicted in a stress plane. The yield surfaces proposed in this study are determined for a plastic offset of E E Ep ij p ij p eq = 2 3/ = 0.2 %. The results in Figure 2 correspond to the same material identified in the previous paragraph, and represent the macroscopic yield surface after two amounts of macroscopic equivalent strain (Eeq = 5 %, Eeq = 10 %) of a tensile pre-straining in direction 1. These yield surfaces reveal: • A vertex in the pre-straining direction, due to great inter and intra-granular residual stress being deve- loped. These back stresses arising from incompati- bilities of strain fields pointed out previously, increase the yield stress in this pre-loading direction. • A flattening of the yield surface in the opposite direction to the pre-straining is also observed. This relative softening results from the decrease of the yield stress due to superimposed back stresses in this opposite direction. • A hardening in directions transverse to the pre- loading is also visible. Whatever the amount of pre-straining, the obtained equivalent behaviour exhibits a mixed anisotropic and kinematic hardening due to the intra-crystalline particles. A translation in the pre-straining direction combined with a non-homogeneous expansion of the yield surfaces is observed in Figure 3. The presence of intra-crystalline-heterogeneities- induced internal stresses relates to the evident interactions between the latter and the dislocations during their movement. These internal stresses are developed both in the matrix and the particle, where a strong polarization is observed. After a pre-loading, leading to a significant permanent strain, followed by an unloading, there remains internal stresses called residual stresses. Figure 4 depicts these residual stresses, developed inside particles for three selected grains of the polycrystalline material. Due to the elastic properties of the selected particles, these residual stresses seem to be relatively important for the macroscopic stress level. In contrast to previous investigation, their evolution is non-linear, because of the introduced new hardening matrix and differs from grain to grain. N. BONFOH, P. LIPINSKI: ANISOTROPIC HARDENING OF MATERIALS BY NON-SHEARABLE PARTICLES Materiali in tehnologije / Materials and technology 41 (2007) 2, 77–80 79 Figure 4: Residual stress inside particle after a tensile pre-straining (in direction 1) Slika 4: Zaostale napetosti v notranjosti delca po natezni predobre- menitvi (v smeri 1) Figure 2. Macroscopic yield surfaces after a tensile pre-straining (in direction 1) Slika 2: Makroskopsko plasti~no popu{~anje povr{ine po natezni predobremenitvi (v smeri 1) Figure 3: Material yield surfaces after a tensile pre-straining (in direction 1, Eeq = 10 %) Slika 3: Povr{ine plasti~nega popu{~anja materiala po natezni pred- obremenitvi (v smeri 1, Eeq = 10 %) 4 CONCLUSION The paper proposes a new formulation for the hardening of polycrystalline materials due to intra-crystalline heterogeneities such as precipitates. The suggested model takes into account internal stresses arising from the incompatibilities of the strain field between the crystalline elastic-plastic matrix and the purely elastic particles. The observed hardening has been elucidated with numerical simulations of the behaviour response under monotonous, uni- and multi-axial loadings. The results, in terms of the yield surfaces and the internal stresses evolution with the volume fraction of the particle and also with the amplitude of specific pre-straining, are provided. These internal stresses, characterised by a strong polarisation in the particles, leads to important interfacial stresses and the subsequent damage initiation by the particle’s interface debonding. For this specific problem, a combined stress and energetic criterion is introduced, to take into account the particles’ size effect, pointed out for the damage initiation in such materials. 5 REFERENCES 1 Ashby, M. F., The deformation of plastically non-homogeneous materials. Phil. Mag. 21 (1970), 399–424 2 Berveiller M, Zaoui A., An extension of the self-consistent scheme to plastically-flowing polycrystals. J. Mech. Phys. Solids. 26 (1984), 325–344 3 Barbe F., Decker L., Jeulin D., Cailletaud G., Intergranular and intragranular behavior of polycrystalline aggregates. Part 1: F. E. model. Int. J. Plasticity 17 (2001), 513–536 4 Barbe F., Forest S., Cailletaud G., Intergranular and intragranular behavior of polycrystalline aggregates. Part2: Results. Int. J. Plasticity 17 (2001), 537–563 5 Bochet, L., Delobelle, P., Robinet, P., Feaugas, X., Mechanical and microstructural investigations of an austenitic stainless steel under non-proportional loadings in tension-torsion-internal and external pressure. Int. J. Plasticity 17 (2001), 1491–1530 6 Bonfoh N., Modélisation micromécanique de l’écrouissage des matériaux polycristallins contenant des hétérogénéités intra- granulaires. Introduction à l’endommagement ductile. PhD thesis, Université de Metz, 2001 7 Bonfoh N., Lipinski P., Carmasol A., Modeling of intra-crystalline hardening of materials with particles. Int. J. Plasticity,19 (2003), 1167–1193 8 Bonfoh N., Lipinski P., Carmasol A., Tiem S., Micromechanical modeling of ductile damage of polycrystalline materials with heterogeneous particles. Int. J. Plasticity, 20 (2003), 85–106 9 Han, C. S, Wagoner, R. H., Barlat, F., On precipitate induced hardening in crystal plasticity: theory. Int. J. of Plasticity 20 (2004), 477–494 10 Reza S. Yassar, Sinisa Dj. Mesarovic, David P. Field, Micro- mechanics of hardening of elastic-plastic crystals with elastic inclusions. I-Dilute concentration (submitted) Int. J. Plast. 11 Schmitt C., Lipinski P., Berveiller M., 1997. Micromechanical Modelling of the elastoplastic behavior of polycrystals containing precipitates-Application to hypo- and hyper-eutectoid steels. Int. J. Plast. 13 (1997) 3, 183–199 N. BONFOH, P. LIPINSKI: ANISOTROPIC HARDENING OF MATERIALS BY NON-SHEARABLE PARTICLES 80 Materiali in tehnologije / Materials and technology 41 (2007) 2, 77–80 M. HUDÁKOVÁ ET AL.: ANALYSIS OF THE BORONIZED LAYER ON K 190 PM TOOL STEEL ANALYSIS OF THE BORONIZED LAYER ON K 190 PM TOOL STEEL ANALIZA BORONIZIRANE PLASTI V ORODNEM JEKLU K 190 PM, IZDELANEM PO POSTOPKU METALURGIJE PRAHU Mária Hudáková, Martin Kusý, Viktória Sedlická, Peter Grga~ Department of Materials Engineering, Faculty of Materials Science and Technology, Slovak University of Technology, J. Bottu 24, 917 24 Trnava, Slovakia maria_hudakovastuba.sk Prejem rokopisa – received: 2006-09-29; sprejem za objavo – accepted for publication: 2006-11-10 This work deals with the characterization of the phases in boronized layers formed for particular processing parameters on tool steel of the ledeburite type. The high-alloyed tool steel K190 PM prepared by powder metallurgy was used as a substrate. X-ray diffraction (XRD) measurements were carried out at different depths of boronized layers, as well as of transient layers, using a Philips PW1710 diffractometer. The analysis revealed that the surface consisted of FeB (MeB) and an unidentified crystalline phase (or phases), which amount decreases with increasing distance from the surface. Also, a gradual increase in the amount of Fe2B (Me2B) at the expense of MeB was observed with increasing distance from the surface. The layer neighboring the substrate consists mainly of Me2B and CrB. The transition layer between the borides and the base material was found to contain the Me23(C,B)6 phase. Key words: tool steel, boronized, layer, constitutive phases, microhardness Predstavljena je karakterizacija faz v boronizirani plasti, pripravljeni na ledeburitnem orodnem jeklu. Kot podlaga je uporabljeno jeklo K 190 PM, pripravljeno po postopku metalurgije prahu. Difrakcija rentgenskih `arkov je bila izvr{ena pri razli~nih globinah boronizirane in prehodne plasti z difraktometrom Philips PW 1710. Analiza je pokazala na povr{ini fazo FeB (MeB) in neidentificirano kristalno fazo (ali faze), katerih vsebnost se zmanj{uje z razdaljo od povr{ine. Ugotovljena je tudi postopna rast koli~ine faze Fe2B (Me2B) na ra~un faze MeB z oddaljevanjem od povr{ine. Plast na meji s podlago je iz Me2B in CrB. V prehodni plasti med boridi in osnovnim materialom smo na{li faze Me23(C,B)6. Klju~ne besede: orodno jeklo, boronizirana plast, sestavne faze, mikrotrdota 1 INTRODUCTION In addition to the popular techniques of carburizing and nitriding, boronizing also plays an important, though less common, role in thermo-chemical treatments. Boronizing usually leads to the formation of a very hard and wear-resistant layer on the surface of the material. Due to the diffusion of the boron into the surface of the carbon steel, iron borides with very high hardnesses are formed. The boronized layers exhibit a higher hardness than the layers prepared with carburizing or nitriding. Generally, low-carbon and low-alloyed steels are boronized. This process can also be applied to austenite steels, non-ferrous alloys (nickel or cobalt), heat- resistant metals such as Ti and Ta and also to sintered carbides, such as WC-Co, with the aim of improving their wear resistance 1,2. Boronizing is currently also used for a number of applications where no surface treatment has been used so far. It is applied to various forming tools, components working in an abrasive environment, dies for the injection casting of non-ferrous alloys and others. Furthermore, it is suitable for preparing functional layers with an extended lifetime of the substrate material due to the increased hardness of the surface. During the boronizing of plain steels, FeB or Fe2B is formed in the surface layers. The thickness of the boride layers typically varies in the range from 20 to 300 µm 3,4. The aim of this paper is to present a detailed XRD characterization of the boronized layer formed on the surface of K 190 PM high-alloyed tool steel. 2 EXPERIMENTAL The boronized layer investigated was formed on K 190 PM tool steel prepared by powder metallurgy. The chemical composition of the steel is shown in Table 1. Table 1: Chemical composition of K 190 ISOMATRIX PM steel Tabela 1: Kemijska sestava jekla K 190 ISOMATRIX PM Chemical element Fe C Si Mn P S Cr V Mo Content w/% ba- lance 2.29 0.5 0.32 0.0260.019 12.3 4.06 1.06 The steel was delivered as a rod with a diameter of 10 mm. For the purpose of the boronizing experiment, the rod was cut into 10-mm-long pieces. The pieces were cleaned and degreased and boronized at at 1000 °C for 8 h in a (80 × 85 × 125) mm hermetic container filled with a DURBORID powder mixture. After boronizing the samples were cooled to room temperature in containers in the open air. The microstructure of the boronized K 190 PM tool-steel layer consisted of a boride and a transient layer. The microhardnesses of the boride layer, Materiali in tehnologije / Materials and technology 41 (2007) 2, 81–84 81 UDC/UDK 669.14.018.252:661.977 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)81(2007) the transient layer, and the substrate were measured with a Hannemann indenter fitted to a Zeiss Neophot 21 light microscope. The XRD patterns of the boride layer, the transient layer and the substrate were recorded with a Philips PW 1710 diffractometer using Fe-filtered CoKα radiation. The experimental device was equipped with a secondary graphite monochromator and a pulse-height proportional detector. Data were recorded in the range 20–144° (2θ) using a constant step of 0.05° and a 10 s exposure time and samples rotating. Qualitative and quantitative analyses of the phases were also carried out using X-ray diffraction on the as-delivered substrate of soft-annealed K 190 PM steel. Exploiting the natural penetration depth of the X-ray radiation, the analysis was performed in consecutive steps. First, the as-prepared boronized substrate without any mechanical cleaning was examined using XRD and microhardness measurements. Next, a predetermined volume of the surface layer was removed with grinding with APEX diamond discs. The new surface was then examined using the same experimental tools. The series contained a total of 10 XRD patterns, taken from the as-prepared surface and the surfaces at the depth of (2, 7, 20, 30, 40, 50, 60, 70, 80) µm. It should be emphasized that the information extracted from a single diffraction pattern taken from a boride layer represents the volume corresponding to the penetration depth of X-rays. Therefore, it cannot be considered as a characterization of the surface but rather of a layer of a certain thickness located at various depths. The identification of the peaks in the XRD patterns was based on a comparison of the calculated hkl and the dhkl list of relevant phases determined from the chemical composition of the substrate, further revised with data on the phase identification in the boride layers 1-5. For the quantitative treatment of the XRD data the Powder Cell 2.4 computer program and, in special cases, the MAUD program based on a Rietveld algorithm 6-8) were used. All the crystallographic data used for the calculation of the theoretical diffraction patterns were taken from Pearson’s handbook 9. 3 RESULTS AND DISCUSSION The microstructure of the soft-annealed substrate of K 190 PM is shown in Figure 1. It consists of a uniform distribution of small spheroidized M7C3 and MC carbides in the ferrite matrix, which is characteristic for the soft annealed K 190 PM tool steel. The results of the quantitative phase analysis of the substrate determined from the XRD (Figure 1) by the Rietveld method using the Maud computer program are shown in the table in Figure 1. Figure 2 shows the XRD data collected at different depths of the boride and transient layers. The XRD pattern taken from the as-prepared surface and from 2 µm and 7 µm into the boronized layer consist of the dominant FeB (MeB), the minor Fe2B (Me2B), CrB, and an unidentified phase (or phases). The diffraction peaks attributed to the unidentified phase (or phases) did not appear in the diffraction pattern taken from the surface located at a depth of 20 µm. The intensity of the reflections attributed to the Me2B phase increased slightly between the 7-µm and 20-µm scan. The observed trend of increasing intensity of the Me2B phase was confirmed with a series of XRD scans taken from depths of 30 µm and 40 µm. It can be concluded from the next layer in the series, taken from a depth of 50 µm, that the MeB phase becomes a minor phase. It is also possible to observe two strong reflec- tions of the Me2B phase with increasing intensity. Moreover, the reflection observed at approximately 52° (2) suggests an increasing diffraction intensity of the ferrite matrix. Both the MeB and the Me2B observed in the depth ranges from 0 µm to 60 µm were influenced by epitaxial growth of the (020) and (002) planes, respectively. When the material from the boronized layer is removed down to 70 µm or 80 µm the XRD patterns change significantly. It was found that the layer down M. HUDÁKOVÁ ET AL.: ANALYSIS OF THE BORONIZED LAYER ON K 190 PM TOOL STEEL 82 Materiali in tehnologije / Materials and technology 41 (2007) 2, 81–84 a. u. θ ϕ Figure 1: Microstructure and XRD pattern of the soft-annealed K 190 PM steel. A table summarizing the amounts of the different phases is included Slika 1: Mikrostruktura in XRD-posnetek mehko `arjenega jekla s tabelo z vsebnostjo razli~nih faz below 70 µm contains complex iron-chromium carbo- boride, with a crystal structure corresponding to Me23(C,B)6, and ferrite. Neither MeB nor Me2B were identified in this depth range. It is, therefore, suggested that the patterns taken from depths of 70 µm and 80 µm are already located in the transient zone between the boride layer and the substrate. The presence of this phase agrees well with earlier publications, where Cr was considered to be an element less soluble in the boride layer. As a result of this behavior, Cr is enriched in the transient zone ahead of the boride layer and forms the complex Cr-based carbide Me23(C,B)6 or carbo-borides, where carbon is partially replaced with boron 5. A simplified representation of the phase composition in the boride and the transient layers is shown in Figure 3a in the form of a color map of dominant phases as a function of layer depth. The map shows that Fe2B (Me2B) and FeB (MeB) are present in both the boride and the transition layers after 8 hours of boronizing at 1000 °C. Figure 3b shows the quantity of FeB (MeB) and Fe2B (Me2B) determined from XRD patterns and the variation of the microhardness as a function of the depth of the layer. The quantity of MeB and Me2B was determined neglecting other phases (the unidentified phase (or phases), CrB and Me23(C,B)6). The quantitative results were obtained from the Rietveld algorithm implemented in Powder Cell 2.4 6. In this graph, the microhardness profile across the analyzed depth profile is also shown. The color map (Figure 3a), in combi- nation with the quantity of phases and the microhardness M. HUDÁKOVÁ ET AL.: ANALYSIS OF THE BORONIZED LAYER ON K 190 PM TOOL STEEL Materiali in tehnologije / Materials and technology 41 (2007) 2, 81–84 83 a. u. θ θ Figure 2: XRD patterns taken from a series of depths in the boronized layer formed on K190 PM tool steel Slika 2: XRD-posnetki pri razli~nih globinah boronizirane plasti na jeklu K 190 PM θ D ep th , d /m m d ϕ (F eB )/ % , (F e B )/ % 2 ϕ a) Figure 3: Color map showing the occurrence of the identified phases as a function of the layer depth determined from XRD (3a) and the quantities of FeB (MeB) and Fe2B (Me2B) determined from the XRD experiment and the microhardness profile plotted as function of the depth of the layer (3b). Slika 3: Barvni prikaz prisotnosti identificiranih faz v odvisnosti od globine plasti, dolo~en z XRD-posnetkom (3a), vsebnost FeB (MeB) in Fe2B (Me2B) dolo~ene iz teh posnetkov in mikrotrdota v odvisnosti od globine plasti (3b) profile shown in Figure 3b, is proposed as an effective tool for designers and producers. 4 CONCLUSIONS The paper presents an estimate of the quantities of MeB and Me2B at various depths in the boronized layer. The quantitative evaluation was performed, neglecting other phases (an unidentified phase (or phases) or CrB and Me23(C,B)6) present in the boronized layer. It was found that the uppermost layers, down to the depth of 20 µm, consist of MeB and some unidentified phases. The amount of unidentified phase (or phases) decreases with the increasing depth below 7 µm and this phase is not found at the depth of above 20 µm. Most of the MeB was found in the layers from 20 µm to 40 µm. MeB was not found in the layers from 70 µm to 80 µm. The majority of the Me2B was found in the range from 50 µm to 60 µm. A minor amount of this phase was found in other layers over the entire thickness of the boronized layer. A strong epitaxy into the (020) and (002) planes was observed for the growth of the MeB and Me2B, respectively. The CrB phase was identified in small quantities from XRD patterns in the range from 0 µm to 50 µm from the surface. Diffraction patterns from depths of 70 µm to 80 µm revealed the presence of Me23(C,B)6, which identifies this region as the transient layer. The highest microhardness of the boronized layer was observed in the range 20 µm to 50 µm from the surface. In the range from 0 µm to 20 µm the measured microhardness was found lower down, probably due to the influence of the unidentified phase (or phases). Acknowledgement The research has been supported by VEGA M[ and SAV 1/3190/06 and EUREKA project E!3437. 5 REFERENCES 1 Dufek, V., Schejbal, T., Bro`ek, V.: Strojírenská výroba, 34 (1986) 3, 197–200 2 Bhushan, B., Gupta, B. K.: Handbook of Tribology, McGraw – Hill, Inc., New York, 1991 3 Kroulík, M.: Strojírenská výroba, ^VUT, Praha, 1991, 8–11 4 Grafen, W., Edenhofer, B.: New developments in thermo-chemical diffusion processes, Surface & Coatings Technology 200 (2005), 1830–1836 5 M. A. Bejar, E. Moreno, Abrasive wear resistance of boronized carbon and low-alloy steels, Journal of Materials Processing Technology 173 (2006), 352–358 6 Kraus, W., Nolze, G.: PowderCell for Windows, Version 2.4, Federal Institute for Materials Research and Testing, 2000 7 Lutterotti, L., Matthies, S., Wenk, H.-R.: MAUD: a friendly Java program for material analysis using diffraction. IUCr: Newsletter of the CPD, 21 (1999), 14–15 8 Young, R. A.: The Rietveld Method, Oxford University Press, 1995 9 Villar, P., Calvert, L. D.: Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, Vol. 1, American Society for Metals, 1989 M. HUDÁKOVÁ ET AL.: ANALYSIS OF THE BORONIZED LAYER ON K 190 PM TOOL STEEL 84 Materiali in tehnologije / Materials and technology 41 (2007) 2, 81–84 M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL SOOBARJANJE LSM ZA PRIPRAVO KATODNIH MATERIALOV ZA GORIVNE CELICE Marjan Marin{ek, Klementina Zupan, Tanja Razpotnik, Jadran Ma~ek Faculty of Chemistry and Chemical Technology, University of Ljubljana, A{ker~eva 5, 1000 Ljubljana, Slovenia marjan.marinsekfkkt.uni-lj.si Prejem rokopisa – received: 2006-10-17; sprejem za objavo – accepted for publication: 2006-11-27 Lanthanum manganite substituted with strontium is a strong candidate for use as a mixed conductor at the cathode site in solid-oxide fuel-cell systems. In the present work a co-precipitation method for the preparation of LSM is described and some material properties are reported. Since the co-precipitation involved calcination at 1000 °C and subsequent sintering, the crystalline phase development and the material homogeneity were followed by X-ray powder diffraction. It was found that no secondary crystalline phases were present in the system when the samples were thermally treated up to 1200 °C. The sintering behaviour of the calcined powders and their density evolution are also reported. Relatively dense bodies were prepared after thermal treatment at 1200 °C for 2 hours. The electrical characteristics of the material were determined from screen-printed LSM sintered layers. The electrical conductivity of the screen-printed layers increased with the sintering temperature due to an additional densification effect. Porous LSM bodies were prepared by methyl cellulose admixing as a pore-former. Key words: co-precipitation, lanthanum strontium manganite, microstructure, electrical conductivity Najpogosteje uporabljan katodni material v visokotemperaturni izvedbi gorivnih celic (SOFC) je lantanov manganit, dopiran s stroncijem (LSM). LSM s sestavo La0,85Sr0,15MnO3- smo pripravli z novo metodo soobarjanja iz raztopine karbonatov. Oborino smo po filtraciji in spiranju su{ili pri 110 °C in nato kalcinirali pri 1000 °C. Po kalcinaciji pri razli~nih temperaturah in po sintranju smo prahove analizirali z rentgensko pra{kovno analizo. Materiali so po sintranju vsebovali le perovskitno kristalno obliko lantanovega manganita. Parametre sintranja prahov LSM po kalcinaciji pri 1000 °C smo dolo~ili s segrevalnim mikroskopom. Vzorce LSM z dodatkom tvorca por in brez njega smo po sintranju analizirali z elektronsko mikroskopijo. Elektri~ne karakteristike tiskanih plasti LSM smo dolo~ili s {tirito~kovno metodo. Specifi~na elektri~na prevodnost pripravljenih plasti LSM je primerljiva z literaturnimi podatki za materiale LSM, ki so bili pripravljeni po kalcinacijskem postopku. Klju~ne besede: koprecipitacija, lantanov stroncijev manganit, katoda, mikrostruktura, elektri~na prevodnost 1 INTRODUCTION Lanthanum strontium manganite (LSM) is known as a potential cathode material for solid-oxide fuel cells (SOFCs) based on a stabilized zirconia electrolyte 1. In the present generation of SOFCs a nominal composition of La1-xSrxMnO3 (x < 0.2) is normally used. The use of LSM-based cathode materials depends not only on their chemical, structural and thermodynamic characteristics but also on their final microstructure, grain size, pore size and pore-size distribution 2-4. The porosity and pore-size distribution, in particular, play a significant role because of the double role of the LSM in the SOFC system, i.e., the LSM has to be an effective electro- catalyst and at the same time has to be permeable to cathode gases. Another aspect when LSM is used as the cathode material is its homogeneity and the possible interaction with the zirconia-based electrolyte. One potential interaction between the LSM cathode layer and the zirconia-based electrolyte in the SOFC system is diffusion (at the phase boundary) of the Y2O3 from the YSZ into the LSM phase. This diffusion leads to the formation of tetragonal and/or monoclinic ZrO2. Such a ZrO2 crystal-structure transformation is accompanied by a volume change and, as a result, possible crack formation (more intense when the SOFC system is repeatedly heated up and cooled down) 5. Another problem arises if products with a high electrical resistivity at the cathode/electrolyte phase boundary, like La2Zr2O7, are formed. The La2Zr2O7 phase is normally found in systems based on LaMnO3+d/YSZ after a prolonged treatment at high temperatures 6. To suppress the La2Zr2O7 formation, La-manganite is doped with Sr in relatively low concentrations (La1-xSrxMnO3; x < 0.3). Larger additions of Sr also mean a higher risk of new low-conducting SrZrO3 phase formation. The reactions between LSM and YSZ are more pronounced at higher temperatures, lower oxygen partial pressures and for prolonged reaction times. The microstructure and homogeneity characteristics of LSM depend to a large extent on the processing conditions and are hard to control in the conventional ceramic synthesis process, which is based on the diffusion of the components in the solid state at high temperatures. For this reason alternative preparation routes for LSM synthesis based on wet chemical methods have been proposed. A characterization of perovskite powders for the cathode and oxygen Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 85 UDC/UDK 661.871:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)85(2007) membranes prepared by different chemical routes was made by Sefir et al. 7, while the effect of the synthesis route on the catalytic activity of LSM was studied by Bell et al. 8. The results of these studies implied that the carbonate co-precipitation synthesis route delivers a finer powder with more homogeneous composition and surface structure and is thus more suited for mixed conductor applications in SOFC systems. They both based the LSM preparation on the carbonate co-preci- pitation method, following the route described by Tanaka et al., where Na2CO3 was used as the preci- pitating agent 9. Another co-precipitation-based LSM preparation was developed by Ghosh et al. 10 using ammonium carbonate as the precipitating agent. In the present work we applied the carbonate co-precipitation route for batch LSM preparation. Sintering and microstructure-development studies were carried out for the calcined powder. Prepared LSM bodies were also characterised in terms of their electrical conductivity. Since a requirement of the cathodes for a SOFC is a porous body with almost 40 % open porosity, we also modified the cathode material by the addition of a former. 2 EXPERIMENTAL In order to prepare stock solutions of 0.5-M lanthanum nitrate and strontium nitrate, lanthanum oxide (La2O3) and strontium carbonate (SrCO3) were separately carefully dissolved in concentrated nitric acid, while manganous nitrate (Mn(NO3)2.4H2O) was dissolved in distilled water. Predetermined amounts of each solution were then mixed for La0.85Sr0.15MnO3 preparation. The mixed solution was added drop-wise to a precipitating bath containing an aqueous solution of ammonium carbonate in which the amount of ammonium carbonate was in 50 % excess for complete precipitation of the mixed La-Mn-Sr-precursor. The pH of the precipitating bath was kept constant at 8.0 by small additions of aqueous ammonia. The temperature of the reaction mixture in the precipitating bath was adjusted to 65 °C and kept constant. The precipitation reaction took place in a CO2 atmosphere to prevent any undesired oxidation of the manganese to MnO2. The reaction time for the complete co-precipitation process was approximately 3 h. The precipitate was easily filterable, and after filtration was washed several times with 0.125-M aqueous ammonium-carbonate solution. The washed precipitate was dried, first at room temperature under a CO2 atmosphere for several hours, and afterwards, for 6 hours at 110 °C in air. Prior to calcination in a muffle furnace at 1000 °C for 1 hour, the dried powder was ground in an agate mortar. Calcined LSM powders were wet milled in isopropanol, pressed into tablets (Φ = 6 mm, h = 4 mm, p = 70 MPa) and sintered at (1150, 1200 and 1250) °C for 1 h. The carbonate co-precipitation route described above is also suitable for the preparation of larger quantities of mixed La-Sr-Mn-carbonate. The method is also very reliable and the final precipitate is easy to reproduce. To corroborate the above statement we prepared mixed La-SrMn carbonate in three different batches (S1, S2, and S3) and compared some of their properties. Typical quantities of initial precursors consumed in one cycle of the bath precipitation were 100 g of (NH4)2CO3 dissolved in 2.5 L of H2O, 854 mL of 0.5-M Mn-precursor solution, 363 mL of 0.5-M La precursor solution, and 128 mL of 0.5-M Sr precursor solution. Such a reaction mixture yielded approximately 100 g of final calcined LSM powder. Samples were characterized with the X-ray powder-diffraction technique using D4 ENDEAVOR apparatus. The shrinkage during sintering was measured by a LEITZ WETZLAR heating microscope. LSM samples with the mass fraction of 10 % of pore former (methyl cellulose) were mixed in a ball mill for (3, 30 and 180) min, pressed into pellets and then sintered at 1200 °C for 1 h. After sintering the samples were polished and thermally etched and analyzed with an SEM (Zeiss FE SUPRA 35 VP). The quantitative analyses of the microstructures was performed using a Zeiss KS300 3.0 image-analysis system. Samples for the electrical measurements were prepared by screen-printing LSM pastes on alumina substrates. Prior to printing the LSM powder was dispersed in an organic liquid phase in ratios from 40 % solid LSM powder and 60 % liquid phase, and then submitted to a homogenization procedure. 3 RESULTS AND DISCUSSION The calcined powders (for XRD study purposes the calcination of the precursor carbonate-hydroxide powder was carried out in the temperature range 1000–1300 °C) M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL 86 Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 Figure 1: XRD patterns of the powder samples calcined at 1000 °C Slika 1: Pra{kovni posnetki vzorcev, kalciniranih pri 1000 °C were submitted to XRD examination. The XRD patterns of the powder samples prepared from the three different batch experiments calcined at 1000 °C are shown in Figure 1. According to Figure 1 perovskite LSM phase formation (LSM XRD peaks are denoted with the letter p) is practically complete; after calcination at 1000 °C the LSM perovskite phase is quite well crystallised. XRD analysis revealed only traces of the residual secondary phases Mn3O4 and La2O3 or La(OH)3 resulting from the reaction of lanthanum oxide with moisture, forming lanthanum hydroxide (all denoted in Figure 1). The relatively small amount of secondary phases in the sample calcined up to 1000 °C makes the co-precipi- tation method favourable when compared to synthesis processes that are based on the diffusion of the components in the solid state (in this case the amount of secondary phases is normally greater). If the calcination temperature is increased (1200 °C or 1300 °C, Figure 2) the secondary phases are completely dissolved in the perovskite structure. The shrinkage behaviour of the LSM powders prepared with other batch experiments and calcined at 1000 °C are shown in Figure 3. All the samples exhibit a normal thermal expansion in the temperature range below 1000 °C. This region of thermal expansion is followed by a single region of rapid sintering. The rate of rapid shrinkage varies slightly for each sample, which may be a consequence of some differences during the wet milling process. Above 1300 °C the powder shrinkage curves show a flattened region, this is due to completion of the rapid sintering stage and the initiation of the final sintering stage with closing of the entrapped porosity. The final shrinkage (25.8 ± 0.4) % of different LSM batches does not diverge substantially. The shrinkage behaviour of LSM powders prepared with other batch experiments and calcined at 1000 °C are shown in Figure 3. All samples exhibit normal thermal expansion in the temperature range below 1000 °C. This region of thermal expansion is followed by a single region of rapid sintering. The rate of rapid shrinkage varies slightly for each sample, which may be a consequence of some differences during the wet milling process. Above 1300 °C the powder shrinkage curves show a flattened region, this is due to completion of the rapid sintering stage and initiation of the final sintering stage with closing of the entrapped porosity. The final shrinkage (25.8 ± 0.4) % of different LSM batches does not diverge substantially. Considering the results of the sintering curves, the sintering temperatures and sintered densities were calculated for all batch-prepared LSM materials. The sintering started at approximately 1050 °C. The sintering temperature and the sintered density varied from 1250 °C to 1280 °C with an average value calculated to be 1265 °C, and 5.61 g/cm3 to 6.10 g/cm3, with an average calculated value 5.84 g/cm3, respectively. The shrinkage process finished at approximately at 1390 °C. Scanning electron micrographs of an average LSM sample (batch S3) sintered at different sintering temperatures (1 hour of sintering) are presented in Figure 4. In general, with increasing sintering tempera- ture the LSM microstructure becomes denser. However, the three microstructures presented (Figure 4a–d) do not differ significantly, implying that LSM can be successfully sintered, even at 1150 °C. The above micrographs show a microstructure with denser regions as well as regions with some porosity in the micrometre range. If the dense regions are examined in greater detail (Figure 5), excellent particle-to-particle packing can be seen. These regions contain practically no porosity. Such a microstructure with completely dense regions separated by micrometer-range pores could be a consequence of insufficient powder homoge- nization prior to the sintering tests. Namely, any larger agglomerates left in the powder material after the wet-milling process would cause differences in the M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 87 Figure 3: Typical shrinkage curves of LSM powders prepared in different batch experiments Slika 3: Zna~ilne krivulje sintranja vzorcev LSM za tri poskuse obarjanja Figure 2: XRD patterns of the powder S1 calcined at different temperatures Slika 2: Pra{kovni posnetki vzorca S1, kalciniranega pri razli~nih temperaturah density through the sample. However, the fact that the LSM powder can be sintered at relatively low tempe- ratures (less than 1200 °C, if the material is properly treated prior to sintering) is very important from the applicability point of view. If the material is used as a cathode in SOFC systems then sintering temperatures above 1200 °C may cause some highly undesirable reactions with neighbouring materials. The specific conductivities of screen-printed LSM layers on Al2O3 substrates, sintered at 1200 °C and 1300 °C, are shown in Figure 6. The electrical conductivity was measured by the four-probe technique. The electrical conductivity of the sample sintered at 1300 °C and measured at 800 °C was found to be 60 S/cm, and for the sample sintered at 1200 °C, 35 S/cm. As expected, the results show that a higher sintering temperature increased the material’s electrical conduc- tivity. According to the literature, in samples sintered at 1400 °C the electrical conductivity is increased to 65 S/cm, exclusively due to the additional effect of densification (10). The relatively high measured specific electrical conductivity at 1300 °C (and even at 1200 °C) M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL 88 Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 d) Figure 4 (a–d): SEM micrograph of LSM material sintered at different temperatures Slika 4 (a–d): Posnetki vzorcev LSM po sintranju pri razli~nih temperaturah ( ( Figure 6: Measurements of specific electrical conductivity for samples sintered at 1200 °C and 1300 °C Slika 6: Specifi~na elektri~na prevodnost vzorcev, sintranih pri 1200 °C in 1300 °C Figure 5: Microstructure of polished und thermally etched sample S1 after sintering at 1150 °C (the dense region) Slika 5: Mikrostruktura vzorca S1 po sintranju (1150 °C), poliranju in termi~nem jedkanju (gosto podro~je) also means that good particle-to-particle contact had already been established at these temperatures. Some special applications of LSM-based materials require not only excellent contact between the particles but also a relatively high open porosity. To test the ability to form such porous LSM bodies with good particle-to-particle contact, methyl cellulose as a pore-former was admixed prior to tablet shaping and sintering. After methyl cellulose addition an increase in the porosity after different mixing times in the ball mill was observed for the sintered LSM samples containing mass fraction 10 % of pore former. The scanning electron micrographs in Figure 7 show that for a mixing time of 2 minutes the sample is composed of dense and less-dense regions. Less-dense regions contain larger several-micrometre-sized pores, while the more-dense regions contain smaller micron or sub-micrometre pores. This relatively high inhomogeneity in the material microstructure is a consequence of an insufficient homogenization time. Using image-analysis software (Karl Zeiss KS 300 3.0) the porosity in the dense regions was estimated as 5.8 %, while the overall porosity (calculated from the material’s geometrical density) reached 11.5 %. After 30 minutes of homogenization time the microstructure of the material still revealed differences in density throughout the sample, yielding 7 % porosity in dense regions and 23.9 % overall porosity. A superb microstructure was achieved after 180 min of homogenization, with no substantial differences in the material’s microstructure throughout the sample (33.7 % of overall porosity). Considering the fact that the electrical conductivity in the latter case still remains very close to that of the dense sample, such a microstructure could already be applied at the cathode site in a potential SOFC system. 4 CONCLUSIONS LSM was synthesized by a non-conventional wet-chemical method. This co-precipitation route included mixed La-Sr-Mn carbonate preparation and its subsequent thermal treatment. Calcined LSM powders after treatment at 1000 °C were very crystalline, with only traces of secondary phases, which completely dissolved in the LSM perovskite structure after heating to 1200 °C. The sintering characteristics of the prepared LSM powders revealed that dense structures could be prepared at sintering temperatures as low as 1150 °C, but the problem of powder homogenization prior to sintering has still to be solved. The specific electrical conductivity of sintered LSM elements reached 35 S/cm (measured at 800 °C) after sintering at 1200 °C and 60 S/cm (measured at 800 °C) after sintering at 1300 °C. Such a relatively high electrical conductivity is believed to be a consequence of the good powder sinterability and hence M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 89 Figure 7: SEM micrographs of LSM material with the addition of the mass fraction 10 % pore former for different times of homogenization in the ball mill Slika 7: SEM-posnetek materiala LSM z 10-odstotnim dodatkom por, ki so nastale pri razli~nih ~asih homogenizacije pri mletju the excellent particle-to-particle contact. Homogeneous sintered bodies with a higher porosity, but still good contact between the particles, were achieved after admixing methyl cellulose as a pore former and a prolonged homogenization of 180 min in a ball mill, with subsequent sintering at 1200 °C. 5 REFERNECES 1 N. Q. Minh, Ceramic Fuel Cells, J. Am. Ceram. Soc., 76 (1993) 563 2 R. Mahesh, R. Mahendiran, A. K. Raychaudhuri, C. N. R. Rao, Effect of particle size on the giant magnetoresistance of La0.7Ca0.3MnO3, Appl. Phys. Lett., 68 (1996), 2291 3 Y. Huang, Z. Xu, C. Yan, Z. Wang, T. Zhu, C. Liao, Soft chemical synthesis and transport properties of La0.7Sr0.3MnO3 granular perovskites, Solid State Commun., 114 (2000), 43 4 N. Zhang, W. Ding, W. Zhong, D. Xing, Y. Du, Tunnel-type giant magnetoresistance in the granular perovskite La0.85Sr0.15MnO3, Phys. Rev., B56 (1997), 8138 5 L. Kilizendermann, D. Das, D. Bahadur, R. Weiss, H. Nikel, K. Hilpert, Chemical interaction between La-Sr-Mn-Fe-O based perovskites and yttria stabilized zirconia, J. Am. Ceram Soc., 80 (1997) 4, 909–914 6 C. Clausen, C. Bagger, J. B. Bilde-Sorensen, A. Horsewell, Microstructural and microchemical characterization of the interface between La0.85Sr0.15MnO3 and Y2O3-stablized ZrO2, Solid State Ionics, 70/71 (1994), 59–64 7 R. J. Bell, G. J. Millar, J. Drennan, Influence of synthesis route on the catalytic properties of La1-xSrxMnO3, Solid State Ionics, 131 (2000), 211 8 J. Sefir, S. Vaucher, P. Holtappels, U. Vogt, H.-J. Schindler, J. Van herle, E. Suvorova, P. Buffat, D. Perret, N. Xanthopoulos and O. Bucheli, Characterization of perovskite powders for cathode and oxygen membranes made by different synthesis routes, Journal of the European Ceramic Society, 25 (2005), 1991–1995 9 J. Tanaka, K. Takahashi, Y. Yajima, M. Tsukioka, Lattice constants of monoclinic (La0.8Ca0.2)MnO3, Chemistry Letters, (1982), 1847 10 A. Ghosh, A. K. Sahu, A. K. Gulnar, A. K. Suri, Synthesis and characterization of lanthanum strontium manganite, Scripta Materialia, 52 (2005), 1305–1309 M. MARIN[EK ET AL.: A CO-PRECIPITATION PROCEDURE FOR THE SYNTHESIS OF LSM MATERIAL 90 Materiali in tehnologije / Materials and technology 41 (2007) 2, 85–90 J. ZELI] ET AL.: THE EFFECT OF SILICA FUME ADDITIONS ON THE DURABILITY ... THE EFFECT OF SILICA FUME ADDITIONS ON THE DURABILITY OF PORTLAND CEMENT MORTARS EXPOSED TO MAGNESIUM SULFATE ATTACK VPLIV DODATKA SILICA FUME NA TRAJNOST CEMENTNIH MALT PORTLAND, IZPOSTAVLJENIH DELOVANJU MAGNEZIJEVEGA SULFATA Jelica Zeli}, Ivana Radovanovi}, Dra`an Jozi} Faculty of Chemical Technology, University of Split, Teslina 10/5, HR-21 000 Split, Croatia jelica.zelicktf-split.hr Prejem rokopisa – received: 2006-09-18; sprejem za objavo – accepted for publication: 2006-10-25 The deterioration of concrete structures due to the presence of sulfate in soils, groundwater and marine environments is a well-known phenomenon. The use of blended cements incorporating materials such as natural pozzolans, fly ash or silica fume have an important role in the long-term durability of concrete exposed to sulfate attack. In this work, the partial replacement (mass fractions 2–15 %) of commercially blended Portland cement (Type CEM IIA-S 42.5N) with silica fume (SF) was investigated in terms of the resistance of SF mortars to a magnesium sulfate (MgSO4) solution (ω(SO4)2– = 2.5 %; exposure time up to 6 months at 20 °C). The changes in the length of the mortars due to expansion, as well as the changes in the elasticity modulus and mechanical strength loss as a function of silica fume replacement were determined. Structural changes in the mortars due to magnesium sulfate attack were also evaluated by differential thermal analysis (DTA) and by X-ray diffraction (XRD) analysis. The results show that the replacement of SF in most cases has a positive effect on the magnesium sulfate resistance. Mortars containing more than 8 % of mass fraction of SF replacement are characterized by a good sulfate resistance and show lower expansion than the control mortar on the basis of the sulfate-resisting Portland cement, due to the absence of the gypsum and ettringite, detected by the XRD analysis. It is also shown that the elastic modulus is proportional to the compressive strength, but higher compressive strength does not necessary correlate with a better resistance to sulfate attack. Key words: blended cement, silica fume, sulfate attack, durability, elastic modulus Propadanje strukture betona v okolju sulfata (tla, podzemne vode in morska voda) je znan pojav. Uporaba me{anih cementov z dodatki, kot so naravni pucolani, lete~i pepel ali "silica fume" (SF), ima pomembno vlogo pri dolgotrajni trdnosti betona, izpostavljenega sulfatni koroziji. V tem delu je preu~evan vpliv SF (masni dele` 2–15 %) kot nadomestni dodatek industrijskemu portlandskemu cementu (CEM IIA-S 42.5N) na odpornost cementne malte proti delovanju raztopin magnezijevega sulfata (w(SO4)2– = 2.5 %; ~as izpostavljanja do 6 mesecev pri 20 °C). Merili smo naslednje parametre v odvisnosti od dele`a SF: sprememba dol`ine malt zaradi ekspanzije, spremembe modulov elasti~nosti in izguba mehani~ne ~vrstosti. Strukturne spremembe v maltah, izpostavljenih vplivu sulfatne korozije, smo spremljali z diferencialno termi~no analizo (DTA) in rentgensko difrakcijsko analizo (XRD). Rezultati ka`ejo, da nadomestni dodatek SF v ve~ini primerov pove~uje odpornost malt proti vplivu magnezijevega sulfata. Malte z masnim dele`em dodatka SF, ve~jim od 8 %, so okarakterizirane kot dobro sulfatno odporne in ka`ejo manj{o ekspanzijo kot kontrolna malta na osnovi sulfatno odpornega portlandskega cementa, ker nimajo gipsa in etringita, kar je potrjeno z analizo XRD. Opa`eno je tudi, da je modul elasti~nosti proporcionalen kompresijski ~vrstosti, toda ve~je kompresijske ~vrstosti niso nujno povezane z ve~jo odpornostjo malt proti korozijskemu delovanju sulfata. Klju~ne besede: me{ani cement, «silica fume», sulfatna korozija, trajnost, modul elasti~nosti 1 INTRODUCTION Sulfate attack on cement concrete is a complex process and many factors such as cement type, sulfate cation type, sulfate concentration and exposure period can affect the sulfate resistance1,2. The sulfate ions react with the C3A and Ca(OH)2 in the concrete to produce expansive and/or softening types of deterioration associated with both the ettringite and gypsum formation. When the attacking solution contains Mg2+ ions, such as in magnesium sulfate (MgSO4), the formation of magnesium hydroxide (brucite) and the conversion of the main cement hydrated product, calcium silicate hydrate (CSH gel), into magnesium silicate hydrate (MSH) are also observed3-5. Several ideas have been suggested to increase the resistance of concrete against sulfate attack either by decreasing the porosity (high cement content, low water-to-cement ratio) or by using more resistant types of binders (sulfate-resisting Portland cement, addition of pozzolanas and blast furnace slag). Studies3,6-11 have shown that the addition of pozzolans such as silica fume enhances the strength and durability of mortars exposed to sulfate attack. This improvement in sulfate resistance for the Portland cement–silica fume blended concrete/ mortars or pastes is attributed to the pore refining and pore refinement effect occurring due to a pozzolanic reaction between the calcium hydroxide liberated during the cement hydration process and the silica fume when a new CSH gel was formed. This CSH gel is more susceptible to the magnesium than to sodium sulfate attack. In the advanced stages of attack, therefore, the Ca Materiali in tehnologije / Materials and technology 41 (2007) 2, 91–94 91 UDC/UDK 691.5:624 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)91(2007) ions in the CSH can be completely replaced by Mg ions, leading to the decalcification of the CSH, and the loss of the cementitious structure12. The aim of the this work is to give a comparative evolution of the sulfate resistance and durability of SF-blended cement mortars immersed in 2.5 % magnesium sulfate solution, as well as to evaluate the optimum dosage level of silica fume replacement to minimize the deleterious effects of the Mg-sulfate attack. 2 EXPERIMENTAL Commercial blended Portland cement, CEM IIA-S 42.5N (the Dalmacijacement, Croatia), that was produced by grinding cement clinker and mixing with up to 5 % of gypsum and about 25 % of blast furnace slag was used in this investigation. Its Blaine specific surface area was 3,530 cm2/g. The Bogue composition of the cement clinker was (in mass fractions): C3S = 64.7 %; C2S = 9.6 %; C3A = 7.5 %; C4AF = 10.1 %. The silica fume was collected from the production of ferrosilicon (ex the Dalmacija Ferroalloys Works, Dugi Rat, Croatia), containing mass fraction of SiO2 about 90 %, having a surface area of 18 m2/g, as measured by the BET nitrogen adsorption method, with extremely fine spherical and amorphous particles9,10. The sulfate- resisting cement (the Dalmacijacement, Croatia), Type V according to the ASTM C-150, was used as the control cement. Salts MgSO4 x 7H2O, p. a. were used in the preparation of the magnesium sulfate solution, ω(SO4)2– = 2.5 % (25,000 mg/L). The mortars were prepared from a mixture of blended cement, standard quartz sand and mass fractions of silica fume (0, 2, 5, 8, 11 and 15) %, according to the Croatian Standard HRN EN 196-1 : 2005. The mortar samples were designated as P-0, P-2, P-5, P-8, P-11 and P-15, respectively. The water-to- binder (cement+SF) ratio varied from 0.50 to 0.67; the samples were prepared to have the same flow-table consistency. No super-plasticizer was added. Two kinds of mortar samples were prepared: (i) for mechanical strength tests, prismatic specimens of dimensions (40 × 40 × 160) mm, according to the Croatian Standard HRN EN 196-1:2005; (ii) for the potential sulfate-resistance tests, and for the modulus of elasticity tests, prismatic specimens of dimensions (25.4 × 25.4 × 285.75) mm with two stainless-steel inserts cast into the ends to facilitate accurate monitoring of the changes in length, according to the ASTM C.452-68. After a 24-hour setting period in a humid environment (20 °C, 90 % RH), the specimens were taken out of the moulds and immersed into tap water for 27 d. The first tests were carried out 28 d after the sample preparation, and then the samples were exposed to a magnesium sulfate solution. The changes in the samples due to magnesium sulfate attack were measured after 28 d and every 28 d thereafter for a 180-day period. The structural changes in the samples during sulfate attack were identified by X-ray diffraction analysis (XRD; a Philips PW 1010, Netherlands) with CuKα radiation in the ranges of Bragg's angles of 3–60° (2 theta). The quantity of cal- cium hydroxide and magnesium hydroxide (brucite) in the mortar samples during sulfate attack was determined by differential thermal analysis (DTA-DTG-TG; a Derivatograph MOM, Hungary). 3 RESULTS AND DISCUSSION Figure 1 shows the compressive strength of the mortars exposed in the Mg-sulfate solution at room temperature. There was a continuous increase in compressive strengths with exposure time up to 90 d for all the SF-mortars, except the P-5 sample. After this period, the strengths decreased for all the samples exposed to sulfate attack. The increase observed in the compressive strengths for up to 90 d can be attributed to two types of reaction: (i) the continuous hydration of the unhydrated cement components to form more hydration products in addition to the reaction of SF (or slag from the blended cement used) with the liberated lime to form more CSH leading to increasing compressive strength, and (ii) the reaction of sulfate ions with the hydrated cement components to form gypsum and ettringite. Both reactions lead to a denser structure as a result of the precipitation of the reaction products within voids and micropores. After longer periods the magnesium sulfate attack becomes more dominant, leading to the formation of microcracks, which decreases the strength11. The results obtained also show that the sulfate resistance of these SF mortars seems to be slightly higher than the P-0 mortar, but lower than the control mortar based on the sulfate-resisting cement (designated SPC). Figure 2 presents the results for the expansion of mortars exposed to the sulfate attack. It can be seen that the smallest expansion, even smaller than the control J. ZELI] ET AL.: THE EFFECT OF SILICA FUME ADDITIONS ON THE DURABILITY ... 92 Materiali in tehnologije / Materials and technology 41 (2007) 2, 91–94 Figure 1: Mortars’ compressive strengths vs. exposure time to MgSO4 solution Slika 1: Kompresijska ~vrstost malt vs. ~as izpostavljanja raztopini MgSO4 mortar on the basis of sulfate-resisting Portland cement (the SPC sample), comes from the samples containing mass fractions 11 % and 15 % of silica fume. The highest expansion of 107.82 % and 59.91 % and the intensive trend of the breakdown were observed in the sample without (P-0) and with 2 mass percent (P-2) of silica fume, respectively: the P-0 sample disintegrated after only 90 d, while the P-2 one lasted for 150 d of immersion in the MgSO4 solution before breakdown due to corrosion13. The results obtained show that the greatest expansion was observed in the P-0 mortar, which did not exhibit the lowest compressive strength. The lack of correlation between the results of the compressive-strength measurements and the damage due to sulfate corrosion (microcracks) was also pointed out by other authors14. Figure 3 presents the XRD patterns of the P-0 and P-8 samples exposed to Mg-sulfate attack for 120 d. While the formation both of gypsum (G) and ettringite (E) as the result of sulfate corrosion is detected for the P-0 sample, in the P-8 sample these corrosion products were not found. Figure 4 presents the results of the changes in the elasticity modulus (E-modulus) of the mortars immersed in the Mg-sulfate solution, which were determined by measuring the dynamic resonance frequency. An electrosonometer15 was used for this purpose. Details for the determination of the dynamic resonance frequency and for calculating the modulus of elasticity were described elsewhere13,15. It can be seen that all samples tend to show a continuous decrease in the E-modulus with the time of exposure. Mortars containing silica fume showed a similar or slightly higher E-modulus with respect to the mortar containing no silica fume (the P-0), but these samples showed a significant reduction in modulus with respect to the reference SPC mortars. Figure 5 presents the E-modulus data mortars versus the square root of compressive strength for all the J. ZELI] ET AL.: THE EFFECT OF SILICA FUME ADDITIONS ON THE DURABILITY ... Materiali in tehnologije / Materials and technology 41 (2007) 2, 91–94 93 Figure 3: XRD diffractograms of the P-0 and P-8 mortar samples exposed to MgSO4 solution for 120 d, where Q-quartz, CH-lime, G-gypsum, E-ettringite Slika 3: XRD-difraktogrami za vzorce malt P-0 in P-8 po 120 d izpostavljanja raztopini MgSO4, kjer so: Q-kremen, CH-kalcijev hidroksid, G-gips, E-etringit Figure 4: Variation of the dynamic modulus of elasticity with exposure time of mortars to MgSO4 solution Slika 4: Sprememba dinami~nega modula elasti~nosti glede na ~asovno obdobje izpostavljanja malt v raztopini MgSO4 Figure 5: The dynamic modulus of elasticity vs. compressive strength for mortars exposed to MgSO4 solution Slika 5: Dinami~ni modul elasti~nosti vs. kompresijska ~vrstost za malte, izpostavljene delovanju raztopine MgSO4 Figure 2: Expansions of mortars vs. exposure time to MgSO4 solution Slika 2: Ekspanzija malt vs. ~as izpostavljanja raztopini MgSO4 mortars exposed to the MgSO4 solution. Based on the experimental data following the correlation between the dynamic modulus of elasticity (Ed) and the compressive strength (fc) was observed to be Ed = 19.40 +2.06 fc , which is in good agreement with the relationship Ed = 22 +2.8 fc , which is given in the British Standard B.S.CP 100:1972. Figure 6 presents the results of the quantity of unleached lime and brucite formed in the specimens after 120 d of the magnesium sulfate immersion. The curve, designated W-CH, illustrated the relationship between the quantity of unleached lime and the SF replacement for samples cured in tap water9,10. More leaching in the MgSO4 solution than in the water was demonstrated by the mortars with and without up to mass fraction 8 % of silica fume, suggesting that lime (calcium hydroxide) was converted to gypsum and brucite4. The presence of brucite, determined by DTA, in all the samples after 120 d of sulfate immersion was found, although the brucite formation was not observed by the XRD analysis. The relatively small concentration of brucite formed was related to the gypsum in the samples and explains why the observed XRD peaks due to brucite were always weak16. The presence of brucite obtained in the SF-mortar samples suggests a rapid attack by MgSO4. However, as the pozzolanic reaction proceeded, the reduction in the permeability and the refinement of the pore structure overcame the negative effect of the sulfate attack. 4 CONCLUSIONS Silica fume replacement enhances the durability of mortar exposed to magnesium sulfate attack due to the lowering of the lime content, and therefore the increase of the initial compressive strength, on account of the pozzolanic reaction. Thereafter, by the decreasing of lime content in mortars during the Mg-sulfate immersion the formation of gypsum and ettringite, which are responsible for the decreasing of mortar durability, decreases. Mortars containing the mass fraction more than 8 % of silica fume replacement are characterized by a good sulfate resistance and show lower expansion than a control sulfate-resisting mortar due to the absence of gypsum and ettringite, detected by XRD analysis. Although the formation of brucite, determined by DTA analysis, suggests a rapid attack by MgSO4 solution and an affect on the compressive strength of SF-mortars, the pozzolanic reaction (reduction in permeability and refinement of the pore structure) overcame this negative effect. The elastic modulus is proportional to the compressive strength, but a higher compressive strength does not necessarily correlate with a better durability to sulfate attack. For these experimental conditions, the optimum dosage level of 15 % silica fume replacement to minimize the deleterious effects of magnesium sulfate attack can be proposed. Acknowledgments The authors would like to acknowledge the Commissioners of the European Union for funding under the REINTRO Project, No: ICA2-CT-2002-10003. 5 REFERENCES 1 A. Neville, Cem. Concr. Res. 34 (2004), 1275–1296 2 F. Bellmann, B. Möser, J. Stark, Cem. Concr. Res. 36 (2006), 358–363 3 M. D. Cohen, A. Bentur, ACI Mater. J. 83 (1988), 147–157 4 R. S. Gallop, H. F. W. Taylor, Cem. Concr. Res. 26 (1996), 1013–1028 5 P. K. Mehta, Cem. Concr. Res. 13 (1983), 401–406 6 K. Torii, M. Kawamara, Cem. Concr. Res. 24 (1994), 361–370 7 P. S. Mangat, J. M. Khatib, ACI Mater. J. 92 (1995), 542–552 8 F. Türker, F. Aköz, S. Koral, N. Yüzer, Cem. Concr. Res. 27 (1997), 205–214 9 E. Tkal~ec, J. Zeli}, Zement-Kalk-Gips 11 (1987), 574–579 10 J. Zeli}, R. Krstulovi}, E. Tkal~ec, P. Krolo, Cem. Concr. Res. 29 (1999), 819–826 11 E. E. Hekal, E. Kishar, H. Mostafa, Cem. Concr. Res. 32 (2002), 1421–1427 12 M. Santhanan, M. D. Cohen, J. Olek, Cem. Concr. Res. 33 (2003) 341–346. 13 I. Radovanovi}, Diploma Thesis, Faculty of Chemical Technology, University of Split, 2004 14 E. F. Irassar, A. Di Maio, O. R. Batic, Cem. Concr. Res. 26 (1996), 113–123 15 D. Ton~i}, P. Gugi}, I. Mandi}, Croatian J. Cem. Ind. 1 (1982/83), 3–6 16 S. A. Hortshorn, J. H. Sharp, R. N. Swamy, Cem. Concr. Res. 29 (1999), 1331–1340 J. ZELI] ET AL.: THE EFFECT OF SILICA FUME ADDITIONS ON THE DURABILITY ... 94 Materiali in tehnologije / Materials and technology 41 (2007) 2, 91–94 Figure 6: Lime and brucite content in the SF-mortar samples exposed to MgSO4 solution for 120 d. Control SF-mortars are stored in tap water (designated CH-W) Slika 6: Vsebnost kalcijevega hidroksida (CH) in brucita (MH) v vzorcih SF-malt po 120 d izpostavljanja raztopini MgSO4. Kontrolni vzorci SF-malt so hranjeni v vodovodni vodi (ozna~ba CH-W) V. MARTINAC ET AL.: ACTIVATED SINTERING OF MAGNESIUM OXIDE OBTAINED FROM SEAWATER ACTIVATED SINTERING OF MAGNESIUM OXIDE OBTAINED FROM SEAWATER AKTIVIRANO SINTRANJE MAGNEZIJEVEGA OKSIDA, DOBLJENEGA IZ MORSKE VODE Vanja Martinac, Miroslav Labor, Meri Miro{evi}-Anzulovi}, Nedjeljka Petric Faculty of Chemical Technology, Department of Thermodynamics, Teslina 10/V, 21000 Split, Croatia martinacktf-split.hr Prejem rokopisa – received: 2006-09-18; sprejem za objavo – accepted for publication: 2006-11-19 The process of isothermally sintering magnesium oxide obtained from seawater by substoichiometric (where precipitation of magnesium hydroxide took place with 80 % of the stoichiometric quantity of dolomite lime) and by overstoichiometric precipitation (which took place with 120 % of the stoichiometric quantity of dolomite lime) was examined, with the addition of mass fractions (1, 2 and 5) % TiO2, at temperatures in the range 1550–1750 °C, with (1, 3 and 5) h of soaking at the maximum temperature. The process was followed by a determination of the product density, as well as of the B2O3 content in the sintered samples. The results indicate that, besides favorably affecting the density of the sintered samples, the addition of TiO2, with increasing temperature and soaking time, also significantly reduces the B2O3 content in the sintered magnesium oxide obtained from seawater. A statistical analysis of the obtained data was performed with the “Statistica” package in order to obtain a model of the density () of sintered samples of magnesium oxide relative to the temperature (t), time of isothermal heating (), and the percentage of added TiO2 (w). Key words: substoichiometric precipitation, overstoichiometric precipitation, magnesium oxide from seawater, TiO2 addition, activated sintering Prou~evan je proces izotermnega sintranja magnezijevega oksida, dobljenega iz morske vode z nestehiometrijskim (dodatek dolomitnega apna je bil 80 % od stehiometrijsko potrebne koli~ine) in prestehiometrijskim usedanjem (z dodatkom 120 % od stehiometrijske koli~ine dolomitnega apna) z dodatkom masnih dele`ev TiO2 (1, 2 in 5) %, v temperaturnem intervalu 1550–1750 °C, v trajanju izotermnega sintranja (1, 3 in 5) h. Proces je bil spremljan z dolo~evanjem gostote in dolo~evanjem vsebnosti B2O3 v sintranih vzorcih. Rezultati preu~evanja ka`ejo, da dodatek TiO2 pozitivno deluje na gostoto sintranih vzorcev, s pove~anjem temperature in obdobjem izotermnega sintranja pa ob~utno vpliva tudi na zmanj{anje B2O3 v sintranem magnezijevem oksidu, dobljenem iz morske vode. Da bi dobili model odvisnosti gostote () sintranih vzorcev magnezijevega oksida od temperature (t) in ~asa izotermnega sinteriranja (), ter dele`a dodanega TiO2 (w) je bila narejena statisti~na obdelava dobljenih podatkov z uporabo paketa “Statistica”. Klju~ne besede: nestehiometrijsko usedanje, prestehiometrijsko usedanje, magnezijev oksid iz morske vode, TiO2 dodatek, aktivirano sintranje 1 INTRODUCTION Magnesia1 (MgO) is a very important material for the refractory industry. Due to its high refractory properties (MgO melts at (2823 ± 40) °C), MgO ceramics are chemically inert, resistant to the effect of metal melts, acid gases, alkali slag, neutral salts, and react with carbon only above 1800 °C. Raw magnesites for refractory use are obtained from natural ore or are synthetically processed from seawater. Magnesium oxide obtained from seawater2-9 is a high-quality refractory material, and its advantages lie not only in the huge reserves of seawater10 (1 m3 contains 0.945 kg of magnesium), but in the higher purity of the sintered magnesium oxide (> 98 % MgO). The magnesium oxide used here was obtained from seawater either by substoichiometric (MgO 80 % pptn – where pptn stands for precipitation) or overstoichiometric (MgO 120 % pptn) precipitation of magnesium hydroxide in seawater using dolomite lime. The purpose of this work was to examine the effect of TiO2 additions on the properties (density and the B2O3 content) of sintered samples of magnesium oxide obtained from seawater at elevated temperatures. The examinations described were carried out in order to obtain a model of the density of sintered samples of magnesium oxide relative to the temperature and the time of isothermal heating and the percent of added TiO2. 2 EXPERIMENTAL The composition of the seawater used for the pre- cipitation of magnesium hydroxide was: MgO = 2.124 g dm–3; CaO = 0.5573 g dm–3 The composition of the dolomite lime used as the precipitation agent was as follows (mass fractions, w/%): CaO = 57.17 % MgO = 42.27 % SiO2 = 0.099 % Al2O3 = 0.051 % Fe2O3 = 0.079 % The experimental procedure used to obtain magne- sium hydroxide from seawater was similar to that employed in our previous investigations8,9. The sedimentation rate was increased by the addition of the Materiali in tehnologije / Materials and technology 41 (2007) 2, 95–98 95 UDC/UDK 546.46:546.27:551.464 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)95(2007) optimum amount of the anionic 818A flocculent (polyacrylamide), produced by the Dutch firm Hercules. The experimental procedure used to determine the optimum quantity of the anionic 818A flocculent has been described in a previous study11. The magnesium hydroxide obtained from seawater was dried at 105 °C and then calcined at 950 °C. The composition of the magnesium oxide obtained by precipitation with 80 % of the stoichiometric quantity of dolomite lime was (mass fractions, w/%): 98.76 % MgO, 0.88 % CaO and 0.1934 % B2O3. The composition of the magnesium oxide obtained by precipitation with 120 % of the stoichiometric quantity of dolomite lime was (mass fractions, w/%): 97.92 % MgO , 1.45 % CaO and 0.1055 % B2O3. Mixtures of magnesium oxide were then prepared with the addition of a mass fraction of 1 %, 2 % and 5 % TiO2. The doping oxide used was analytical reagent grade titanium (TiO2 p. a.) in the rutile form, produced by Merck. The samples were homogenized by manual stirring in absolute ethanol (C6H5O p. a.). The mixtures were cold pressed into compacts in a hydraulic press at a pressure of 625 MPa. The compacts were then sintered in a gas furnace, made by a French firm Mecker (type 553), with a zirconium (IV)-oxide lining, at 1550 °C, 1650 °C and 1750 °C with 1h, 3h and 5h of soaking at the maximum temperature. It took approximately 2 h to reach the maximum temperature in the furnace. After sintering, the samples were left to cool in the furnace. The sample density after sintering () was determined from the volume of water displaced from a calibrated cylinder. The boron content in the magnesium oxide samples examined was determined potentiometrically. The variation coefficient for the method used was ± 1 %12. The results shown represent the average of a number of measurements. 3 RESULTS AND DISCUSSION Tables 1 and 2 show the experimentally obtained values for the density under the operating conditions described for sintered magnesium oxide samples (80 % precipitation) and magnesium oxide samples (120 % precipitation). The results indicate that at higher temperatures, such as 1550 °C, 1650 °C and 1750 °C, the increase in the density of the samples of magnesium oxide obtained from seawater is not very significant when compared to the un-doped samples, even though they were sintered under the same thermal conditions. Therefore, at higher temperatures in the range 1550–1750 °C, the effect of the TiO2 addition is less evident because the concentration of the added ions is evenly distributed over the whole grain mass. One can assume that the mass transfer is the same as with pure MgO, determined by the diffusion of O2– ions through the MgO lattice as the slower diffusion species. The increased temperature leads to the increased mobility of the elements forming the crystal lattice, a contact surface is formed between the particles of compacted powder, the porosity is eliminated, and the whole system shrinks13-15. Tables 3 and 4 show the results of the examination of the effect of TiO2 on the boron content in the sintered magnesium oxide samples (80 % and 120 % precipitation) for the operating conditions described. The results indicate that the TiO2 addition significantly influences the reduction of the B2O3 content in the process of the isothermal sintering of magnesium oxide from seawater, even at higher temperatures. The different behavior of the magnesium oxide obtained by 80 % from that obtained by 120 % precipitation of magnesium hydroxide in seawater is due to the different contents of CaO in these samples. CaO retains the B2O3 in the samples during sintering. With MgO (120 % precipitation) the CaO content is much higher (w = 1.45 %) than with MgO (80 % precipitation) (w = 0.88 %), V. MARTINAC ET AL.: ACTIVATED SINTERING OF MAGNESIUM OXIDE OBTAINED FROM SEAWATER 96 Materiali in tehnologije / Materials and technology 41 (2007) 1, 95–98 Table 1: Density (ρ) of sintered magnesium oxide samples (80 % precipitation) with the addition of a mass fraction of 1 %, 2 % and 5 % TiO2, and no sintering aid, t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa Tabela 1: Gostota (ρ) za sintrane vzorce magnezijevega oksida (80-odstotno usedanje) z masnimi dele`i TiO2 in 1, 2 in 5 %, in brez dodatka TiO2, t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa sample t/°C τ/h ρ/(g cm–3) TiO2 addition, w/% no sint. aid 1 % 2 % 5 % 80 % pptn 1550 1 3 5 3.2401 3.3181 3.3514 3.2691 3.3498 3.3643 3.3243 3.3439 3.3953 3.3361 3.3744 3.4017 1650 1 3 5 3.3463 3.3513 3.4230 3.3798 3.3803 3.4128 3.4115 3.4245 3.4385 3.4257 3.4351 3.4510 1750 1 3 5 3.3522 3.3660 3.4244 3.4007 3.4593 3.4881 3.4199 3.4513 3.4933 3.4391 3.4719 3.5289 Table 2: Density (ρ) of sintered magnesium oxide samples (120 % precipitation) with the addition of a mass fraction of 1 %, 2 % and 5 % TiO2, and no sintering aid, t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa Tabela 2: Gostota (ρ) za sintrane vzorce magnezijevega oksida (120-odstotno usedanje) z masnimi dele`i TiO2 in 1, 2 in 5 %, in brez dodatka TiO2, t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa sample t/°C τ/h ρ/(g cm−3) TiO2 addition, w/% no sint. aid 1 % 2 % 5 % 120 % pptn 1550 1 3 5 3.2880 3.3179 3.3739 3.3088 3.3557 3.3887 3.3980 3.4156 3.4409 3.4036 3.4383 3.4619 1650 1 3 5 3.3350 3.3570 3.4244 3.3726 3.4029 3.4412 3.4253 3.4378 3.4662 3.4357 3.4603 3.5047 1750 1 3 5 3.3711 3.4153 3.4591 3.4033 3.4516 3.4786 3.4292 3.4860 3.5018 3.4473 3.5180 3.5434 and favors the Ca2B2O5 formation reaction. In a previous paper16 the X-ray diffraction method was used to prove the content of di-calcium borate (Ca2B2O5) in sintered samples of magnesium oxide from seawater. Therefore, in the process of sintering, B2O3 reacts with CaO to form Ca2B2O5. In papers17-20 the X-ray diffraction method and EDAX analysis have helped determine that during the sintering process the added TiO2 reacts with CaO from the solid MgO-CaO solution and forms calcium titanate CaTiO3. Thus, TiO2 binds a part of the CaO in the CaTiO3 and as a result reduces the content of CaO that reacts with B2O3. In this way a smaller quantity of Ca2B2O5 remains in the sintered sample, i.e., a higher quantity of B2O3 evaporates. With MgO (120 % precipitation) there is an excess of CaO, which favors the formation of Ca2B2O5, while with MgO (80 % precipitation) a larger part of the B2O3 evaporates from the sample into the air during sintering. A positive effect of the TiO2 addition on the reduction of the B2O3 content in the magnesium oxide obtained from seawater makes it possible to achieve a high-purity product, because the hot-strength properties of certain refractory products are significantly affected by their boron content. The results shown in Tables 1 and 2 have also been considered in a regression analysis. The "Statistica" package was used to analyze statistically the data obtained, in order to obtain a model of the density of the sintered samples of magnesium oxide relative to the temperature (t), the time of isothermal sintering (), and the mass fraction of TiO2 added (w). A model of multiple regression has been proposed, = o + 1 t + 2  + 3 ln (w+1) +  for MgO samples (80 % precipitation) and for MgO samples (120 % preci- pitation), where is the density (g cm–3), t is the tempe- rature (°C),  is the time of isothermal sintering (h), w is the mass fraction (%) of added TiO2, o, 1, 2, 3 are unknown coefficients, and  is the random error in the model. Table 4: Effect of TiO2 on B2O3 content in the sintered magnesium oxide samples (120 % precipitation) at t = 1550 °C, 1650 °C, 1750 °C, = 1 h, 3 h, 5 h, p = 625 MPa Tabela 4: Vpliv TiO2 na vsebnost B2O3 v sintranih vzorcih magnezijevega oksida (120-odstotno usedanje) pri t = 1550 °C, 1650 °C, 1750 °C, = 1 h, 3 h, 5 h, p = 625 MPa sample t/°C /h B2O3, w/% TiO2 addition, w/% no sint. aid 1 % 2 % 5 % 120 % pptn 1550 1 3 5 0.0705 0.0582 0.0312 0.0618 0.0354 0.0270 0.0360 0.0264 0.0245 0.0222 0.0119 0.0058 1650 1 3 5 0.0592 0.0315 0.0180 0.0405 0.0286 0.0171 0.0331 0.0254 0.0106 0.0148 0.0068 0.0055 1750 1 3 5 0.0331 0.0177 0.0148 0.0312 0.0119 0.0100 0,0216 0,0106 0.0084 0.0093 0.0046 0.0032 Table 5: Correlation matrix for MgO (80 % precipitation) samples Tabela 5: Matrica korelacije za vzorce MgO (80 – odstotno usedanje) variable correlations (MgO (80 % pptn) samples) t/°C /h TiO2 addition, w/% density, /(g cm–3) t/oC 1.00 0.00 -0.00 0.69 /h 0.00 1.00 -0.00 0.46 TiO2 addition, w/% -0.00 -0.00 1.00 0.43 density, /(g cm–3) 0.69 0.46 0.43 1.00 Table 6: Correlation matrix for MgO (120 % precipitation) samples Tabela 6: Matrica korelacije za vzorce MgO (120 – odstotno useda- nje) variable correlations (MgO (120 % pptn) samples) t/°C /h TiO2 addition, w/% density, /(g cm–3) t/°C 1.00 0.00 -0.00 0.53 /h 0.00 1.00 -0.00 0.51 TiO2addition, w/% -0.00 -0.00 1.00 0.60 density,  /(g cm3) 0.53 0.51 0.60 1.00 Tables 5 and 6 show the correlation matrices for the examined samples. The results indicate that there is no correlation between the independent variables (tempera- ture t, time  and percentage of addition of TiO2). The estimate for the regression function for MgO (80 % precipitation) is: ρ = 2.461 + 0.000512 T + 0.0173 τ + + 0.0434 ln (w + 1) where s = 0.01899; R2 = 0.9057; F(3, 104) = 333.2251; p = 0.00 The estimate for the regression function for MgO (120 % precipitation) is: ρ = 2.686 + 0.000382 T + + 0.0181 τ + 0.0567 ln (w + 1) where s = 0.01413; R2 = 0.9433; F(3, 104) = 576.8233; p = 0.00 V. MARTINAC ET AL.: ACTIVATED SINTERING OF MAGNESIUM OXIDE OBTAINED FROM SEAWATER Materiali in tehnologije / Materials and technology 41 (2007) 2, 95–98 97 Table 3: Effect of TiO2 on B2O3 content in the sintered magnesium oxide samples (80 % precipitation) at t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa Tabela 3: Vpliv TiO2 na vsebnost B2O3 v sintranih vzorcih magnezijevega oksida (80-odstotno usedanje) pri t = 1550 °C, 1650 °C, 1750 °C, τ = 1 h, 3 h, 5 h, p = 625 MPa sample t/°C τ/h B2O3, w/% TiO2 addition, w/% no sint. aid 1 % 2 % 5 % 80 % pptn 1550 1 3 5 0.1894 0.1337 0.0759 0.1249 0.1075 0.0637 0.0913 0.0193 0.0145 0.0245 0.0174 0.0132 1650 1 3 5 0.1445 0.1020 0.0589 0.0750 0.0640 0.0257 0.0272 0,0161 0.0138 0.0216 0.0096 0.0085 1750 1 3 5 0.0862 0.0415 0.0319 0.0594 0.0261 0.0126 0.0193 0.0116 0.0090 0.0158 0.0074 0.0055 The regression functions and the regression coeffi- cients are significant at the level of p = 0.00 in both models. The statistical analysis indicates that these models are useful for an estimation of the density of the samples examined. 4 CONCLUSIONS The statistical analysis using the Statistica package indicates that the proposed models are useful and acceptable for an estimation of the density of the sintered magnesium oxide samples relative to the temperature (t), the time of the isothermal heating (), and the percent of added TiO2 (w). The favorable effect of the TiO2 addition on the reduction of the B2O3 content in the sintered magnesium oxide obtained from seawater is due to two inter- dependent reactions of the formation of Ca2B2O5 and CaTiO3, which lead to a reduction in the B2O3 content during sintering. 5 LITERATURE 1 Shand M. A., The Chemistry and Technology of Magnesia, Willey Interscience, N.York, 2006 2 Frith M., Buttrey T., Strawbridge I., Brit. Ceram. Trans., 97 (1998), 29 3 Sims C., Industr. Minerals, 7 (1997), 21 4 Gilpin W. C., Heasman N., Chem. Ind., 16 (1977), 567 5 Heasman N, Gas Wärme International, 28 (1979), 392 6 Rabad`hieva D., Ivanova K., Balarev Hr., Trendafelov D., @urnal prikla~noji himii, 70 (1997), 375 7 Bonney O. V., U. S. Patent 43 148 85 (to Amstar Corporation, New York) 9 Feb. 1982, Chem. Abstr. 96 (1982), 125549 8 Petric B., Petric N., Ing. Eng. Chem. Des. Dev., 19 (1980), 329 9 Martinac V., Labor M., Petric N., Mater. Chem. Phys., 46 (1996), 23 10 Brown E. et al., Seawater: Its composition, properties and behaviour, Butterworth Heinemann in association with The Open University, Walton Hall, Milton Keynes, 2nd Ed., 1997, 86 11 Petric N., Martinac V., Labor M., Jurin O., Kovine zlit. tehnol., 33 (1999), 475 12 Culkin F., The major constituents of seawater, in: Chemical Oceanography, Ed. by Riley J. P., Skirrow G., Vol.1, Academic Press, London, 1975, 136–151 13 German R. M., Sintering theory and practice, Wiley, New York, 1996 14 Kang Suk-Joong L., Sintering, Elsevier, 2005 15 Lee Y. B., Park H. C., OH K. D., J. Mat. Sci., 33 (1998), 4321 16 Petric N., Petric B., Tkal~ec E., Martinac V., Bogdani} N., Miro{evi}-Anzulovi} M., Sci. Sinter., 19 (1987), 81 17 Chaudhuri M. N., Kumar A., Bhadra A. K., Banerjee G., Sarkar S. L., Ceram. Bull., 71 (1992), 345 18 Chaudhuri M. N., Kumar A., Bhadra A. K., Banerjee G., Interceram., 39 (1990), 26 19 ]osi} M., Pavlovski B., Tkal~ec E., Sci. Sinter., 21 (1989), 161 20 ^eh M., Kolar D., J. Mater. Sci., 29 (1994), 6295 V. MARTINAC ET AL.: ACTIVATED SINTERING OF MAGNESIUM OXIDE OBTAINED FROM SEAWATER 98 Materiali in tehnologije / Materials and technology 41 (2007) 2, 95–98 J. BA@AN, K. STRÁNSKÝ: FLOWING OF THE MELT THROUGH CERAMIC FILTERS FLOWING OF THE MELT THROUGH CERAMIC FILTERS PRETOK TALINE SKOZI KERAMI^NE FILTRE Jií Ba`an1, Karel Stránský2 1Technical University of Mining and Metallurgy in Ostrava, FMMI, 17. listopadu 15, 708 33 Ostrava-Poruba, Czech Republic 2VUT Brno, FSI, Technická 2, 619 69 Brno, Czech Republic jiri.bazanvsb.cz Prejem rokopisa – received: 2006-10-11; sprejem za objavo – accepted for publication: 2006-12-04 It is important to clean melts by filtering them with ceramic filters. Particular attention must be paid to the hydraulic conditions as the melt flows through the filter. The Bernouilli equation is used to describe the filtration of the melts. The physical regularities of the capillary melt can be seen in the flow. The Hagen-Poiseuille law and Darcy’s law have been applied. The characteristics of laminar flow have been examined, as have the attributes of turbulent flow. There are dimensionless criteria characterising the melt flow through a ceramic filter. The principal mechanisms of melt filtration are assessed. Key words: metal melt, filtration, ceramic filter, flow of metals through pores, laminar and turbulent flow Pomembnost ~i{~enja kovinske taline s filtriranjem skozi kerami~ni filter. Hidravli~ni pogoji pretoka taline skozi filter. Uporaba Bernouillijeve ena~be za opis filtriranja taline. Fizikalne zakonitosti kapilarnega toka taline. Hagen-Poiseuillejev in Darcyjev zakon. Zna~ilnosti laminarnega toka. Atributi turbulentnega toka. Brezdimenzijska merila, karakteristi~na za pretok taline skozi kerami~ni filter. Glavni mehanizmi filtriranja kovinskih talin. Klju~ne besede: kovinska talina, filtriranje, kerami~ni filter, pretok taline skozi pore, laminarni in turbulentni pretok 1 INTRODUCTION Recently, many laboratory-, pilot- and industrial- scale tests were carried out in the field of filtration using the casting of cast irons and steels. The results showed that the filtration of the melts of these metallic alloys can be used advantageously in melting foundries and steel shops, using the bottom pouring of steel (i.e., ladles with a taphole at the bottom), as well as in tundish metallurgy during the continuous casting of steel, and in cases when the ceramic filter is applied directly in the entry part of the inlet system during the casting of steel and cast iron. 2 FLOWING A MELT THROUGH THE CERAMIC FILTER The overall efficiency of filtration technology is significantly influenced by the hydraulic conditions during the flowing of the melt through a ceramic filter. The basis for the description and evaluation of the hydraulic conditions during the flowing of a metallic melt (steels and cast irons) through a ceramic filter is the Bernouilli equation, expressing the principle of the conservation of mechanical energy of the flowing liquid 1. It is possible to express the Bernouilli law for an ideal liquid considering the individual energies in the following equation ρ ν ρ ⋅ + ⋅ ⋅ + = 2 2 h g p const. /Pa (1) in which the term (·2/2)/Pa represents the kinetic energy of a unit volume of liquid, the term (h··g)/Pa represents the position or gravitational energy of a unit volume of liquid, given by the earth’s gravity, and the term p/Pa represents the potential pressure energy of a unit volume of liquid, which is usually dependent on externally acting forces. The individual quantities in the equation are as follows: v, the flow velocity of an ideal liquid (m s–1); , the density of the flowing liquid (kg m–3); h, the true position, i.e., the real height of the flowing unit volume of liquid (m); g, the acceleration due to gravity (m s–2); and p, the pressure of a unit volume of flowing liquid in (Pa), the basic dimensions of which are kg s–2 m–1. According to the Bernouilli equation the sum of the kinetic, positional and pressure energies of an ideally flowing liquid remains constant at each point in the flow. However, if this equation is to be used for a description of the flow of a real liquid, e.g., the melts of steels and cast irons, it is necessary to add to the three terms on the left-hand side of the equation one more term, i.e., ez, which expresses the loss of energy per unit volume of a real liquid, i.e., of the melt of steel or cast iron. The loss of energy during the flow of a real liquid through the filter will be directly proportional to the kinematic viscosity of the liquid,  (m2 s–1), to the velocity of the flowing liquid, v (m s–1), indirectly proportional to the structure of the filter characterised by the length, the dimension and the geometry of the holes and capillaries in the filter, d/m, to the melt density, /(kg m–3), and to the dimensionless proportion constant, . It is therefore necessary to express the Bernouilli Materiali in tehnologije / Materials and technology 41 (2007) 2, 99–102 99 UDC/UDK 669.018.95:669.1 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(2)99(2007) equation for a filtered melt of steel or cast iron in the following form: ρ ν ρ ⋅ + ⋅ ⋅ + + = 2 2 h g p ez const. /Pa (2) or in the explicit form ρ ν ρ ξ υ ν ρ δ ⋅ + ⋅ ⋅ + + ⋅ ⋅ ⋅ = 2 2 h g p const. /Pa (3) The flow of a liquid in a round hole (i.e., in a capillary), which is characteristic for direct holes (strainer filters), is governed by the Hagen-Poiseuille law, indicating that the flow volume of a viscous liquid during laminar flow through the tube of circular cross-section is directly proportional to the hydraulic gradient, p/l, and the fourth power of the tube radius, and indirectly proportional to the dynamic viscosity of the flowing liquid. The Hagen-Poiseuille law is written in the form: Qv/(m 3.s –1) = π ⋅ ⋅ ⋅ r p l 4 8 ∆ ∆ η (4) in which Qv is the volume of liquid flowing through the tube, r/m is the internal radius of the tube (i.e., the capillary), /(kg m–1 s–1) is the dynamic viscosity, p is the loss of static pressure along the length of the tube and l is the length of the tube. The hydraulic gradient, (p/l)/(kg m–2 s–2), is in this case a measure of the liquid’s resistance to flow and it is proportional to the first power of the mean velocity of the liquid flowing through the capillary. The Hagen-Poiseuille law also determines implicitly the mean velocity of the liquid flowing through the tube. The dependence of the mean velocity of the laminar flow in a capillary can be expressed by the equation obtained by dividing the flow volume, Qv, by the area of the cross-section of the capillary, πr2 : ν/(m.s –1) = r p l 2 8 ⋅ ⋅ ∆ ∆ η (5) It follows from this equation that the mean velocity of the laminar flow through a capillary increases with the square of the capillary’s radius; it is also directly proportional to the hydraulic gradient and indirectly proportional to the octuple of the value of the dynamic viscosity. These relations are useful for an evaluation of the hydraulic conditions during the flow of melts of steel and ductile iron through ceramic filters, because they determine the causal relations between the physical quantities that characterise the melt, such as the viscosity and the density, the geometric characteristics of the filter, such as the diameters of the holes, the filter’s thickness, the parameters of the flow of the melt through the filter, such as the flow volume of the melt, its flow-rate and also its hydraulic gradient in the filter. The Darcy law is a special case of the Hagen- Poiseuille law. It expresses the velocity of the laminar flow of a liquid in a porous environment. According to this law the velocity of the flow in a porous environment is given by the linear relation νlam/(m s –1) = k p l ∆ ∆ (6) where k/(m3 s kg–1) is the coefficient of leakage (fil- tration) and p/l is the already-mentioned hydraulic gradient. The Darcy law expresses the penetration of the melt into the pores in the walls of the ceramic filter’s capillaries during filtration, which determines the physical-chemical reactions between the flowing melt and the filter ceramics. 2.1 Laminar flow The laminar flow of a real liquid through the capillary is distributed in such a way that the maximum of the flow velocity is achieved in the axis of the circular section of the capillary and the minimum velocity (in practice, a zero flow velocity) near the capillary walls. The distribution of the flow velocity in the capillary in this case approaches the shape of a rotating paraboloid. This distribution of velocity occurs because the outermost layers of the liquid stick to the capillary walls, since there is great friction between the wall and the flowing liquid (in practice, an almost infinitely large friction) and due to the internal liquid friction, individual layers of liquid mutually obstruct their movement. In this way the liquid counteracts during the flow through the capillary, the passing of which requires the pressure, expressed by equations (4) to (6), related to both the mentioned laws. The laminar flow is therefore characterised for a certain mean velocity, making it possible to subsequently calculate the flow volume. During the laminar flowing of a real liquid through capillary vortex rings, vortical threads with the shape of concentric circles with the centres lying in the capillary axis, are formed 2-5. 2.2 Turbulent flow The turbulent flow of liquids (and also gases) is the most widespread and also the most complex form of macroscopic movement of mass in technical equipment and systems. The laminar flow can only be maintained up to a critical mean velocity, because at higher mean velocities of the flowing liquid the perturbative influence of the vortices begins to prevail, the flow begins to change, the liquid threads begin to interlace and turbulent flow begins. During the turbulent flow, the resistance to flow, expressed as the hydraulic gradient, is not directly proportional to the mean velocity of the liquid flowing through the capillary, as in the case of laminar flow, but it increases approximately with its second power. Accordingly, for turbulent flow the J. BA@AN, K. STRÁNSKÝ: FLOWING OF THE MELT THROUGH CERAMIC FILTERS 100 Materiali in tehnologije / Materials and technology 41 (2007) 2, 99–102 following equation is valid for the dependence of the flow velocity of the melt through the filter and the hydraulic gradient: νn = ∆ ∆ p l (7) where v is the mean velocity of the liquid flow and n is the exponent with a value approximately in the interval from 1.75 to 2.00, while for laminar flow its value is n = 1. Let us add, in connection with this, that the turbulent flow is not in accordance with the relations (4) and (5), which follow from the Hagen-Poiseuille law. The turbulence occurs for high values of the dimensionless Reynolds number, when the following inequality is valid for this number: Re = ν υ ⋅ d >> 1 (8) where v is the characteristic velocity of liquid flow (m s–1), d/m is the characteristic length dimension of the filter capillary and /(m2 s–1) is the kinematic viscosity. Unlike laminar flow, turbulent flow has a large number of degrees of freedom, which results from the random vorticular movements of a large number of particles (aggregates, clusters) of liquid, i.e., turbulent vortexes. At each point of the turbulent flow of the liquid there are turbulent vortexes in the filter, which in various places of the capillary can even be formed upstream. The velocity profile during the laminar flow of a liquid through the capillary has an approximately para- bolic distribution of velocity with a distinct maximum at the capillary axis. On the other hand, the velocity profile during turbulent flow of the same liquid through the thin capillary is not parabolic, as it was for laminar flow, but the velocity is almost constant in all the internal parts of the capillary. The exception is the thin layer of liquid at the capillary wall, in which the flow velocity sharply increases from an almost zero value, approximately in proportion to the distance from the wall. The mean velocity of the turbulent flow of the liquid through the thin tube is therefore much closer to the maximum velocity than in the case of laminar flow. The gradient of the flow velocity of the liquid in the direction from the wall of a thin tube to its axis is thus greater during turbulent flow than during laminar flow, see Figure 1. When the melt of steel or ductile iron flows through the strainer or foam filter, regardless of the elemental character of the flow (laminar or turbulent), there exists for both types of filter and for both types of metallic melt, i.e., steel and cast iron, a common character of contact for the flowing metallic melt with the filter walls. The velocity of the filtered metallic melt that is in contact with the filter walls is virtually zero; therefore, the melt can penetrate into the microscopic pores of both types of filter (strainer and foam filter), to a large extent in agreement with the Darcy law, as it is expressed with equation (6). Thus, it is possible to expect on the walls of the holes of the strainer filter, as well as on the walls of the irregular spaces of the foam filter, the same elementary character of physical-chemical reactions of the filtered metallic melt with the filter ceramics. 3 CONCLUSION The hydraulic conditions have an important influence on the overall efficiency of the technology of filtration during the flowing of the melt through a ceramic filter. A knowledge of the hydraulic conditions during the flowing of the melts of steels and cast irons through ceramic filters, based on the physical regularities that govern the flow of metallic melts through the capillaries of ceramic filters, can be the basis for: a) the modelling of basic mechanisms for the flowing of melts of steels and cast irons through ceramic filters, b) a detailed explanation of the physical-chemical mechanisms of filtration occurring during the application of ceramic filters in the practice of foundries casting steels and cast irons, and c) the optimisation of the use of individual types of ceramic filters in the technology of the melting of steels and cast irons. The investigation was performed in the frame of the grant projects GA^R reg. No. 106/04/0393 and 106/04/1006. 4 LITERATURE 1 Horák, Z., Krupka, F., [indelá, V. Technická fyzika. Teoretická kni`nice in`enýra Technical physics, Engineer’s Theoretical Library. Praha, SNTL, 1961, 1436 p. 2 Happ, J., Frohberg, M. G. Untersuchungen zur Filtration von Eisenschmelzen. Giessereiforschung, 23 (1971) 1, 1–9 J. BA@AN, K. STRÁNSKÝ: FLOWING OF THE MELT THROUGH CERAMIC FILTERS Materiali in tehnologije / Materials and technology 41 (2007) 2, 99–102 101 (a) (b) Figure 1: Profile of the velocity of laminar (a) and turbulent flow (b) of the metallic melt through a hole with a circular cross-section Slika 1: Hitrostni profil za laminaren (a) in turbulenten (b) pretok kovinske taline skozi poro z okroglim prerezom 3 Flinn, R. A., Van Vlack, L. H., Colligan, G. A. Mold-metal reactions in ferrous and nonferrous alloys. AFS Transactions 86-09, 29–46 4 Schmahl, J. R., Aubrey, L. S. Filtration with re-circulated silicon carbide foam: an effective means for inclusion removal in gray and nodular iron casting. AFS Transactions 93-213, 1001–1020 5 Acosta G., F. A., Castillejos E., A. H., Almanza R., J. M., Flores V. A. Analysis of liquid flow through ceramic porous media used for molten metal filtration. Metallurgical and Materials Transactions B, 26B (1995), 159–171 J. BA@AN, K. STRÁNSKÝ: FLOWING OF THE MELT THROUGH CERAMIC FILTERS 102 Materiali in tehnologije / Materials and technology 41 (2007) 2, 99–102 S. ^AMPELJ ET AL.: SINTEZA MAGNETNIH NANODELCEV, FUNKCIONALIZIRANIH S TANKO PLASTJO SILIKE SINTEZA MAGNETNIH NANODELCEV, FUNKCIONALIZIRANIH S TANKO PLASTJO SILIKE SYNTHESIS OF MAGNETIC NANOPARTICLES FUNCTIONALIZED WITH THIN LAYER OF SILICA Stanislav ^ampelj1, Darko Makovec1, Marjan Bele2, Miha Drofenik1,3, Janko Jamnik2 1Oddelek za raziskavo sodobnih materialov, Institut "Jo`ef Stefan", Jamova 39, SI-1000 Ljubljana, Slovenija, 2Kemijski in{titut, Ljubljana, Hajdrihova 19, SI-1000 Ljubljana, Slovenija, 3Fakulteta za kemijo in kemijsko tehnologijo, Univerza v Mariboru, SI-2000 Maribor, Slovenija, stanislav.campeljijs.si Prejem rokopisa – received: 2006-09-20; sprejem za objavo – accepted for publication: 2006-11-27 V prispevku opisujemo pripravo maghemitnih nanodelcev prevle~enih s siliko. Postopek je potekal v treh lo~enih stopnjah. V prvi smo delce sintetizirali s koprecipitacijo ionov Fe2+ in Fe3+. V drugi smo delce prevlekli s citronsko kislino, da bi prepre~ili aglomeracijo, in jih dispergirali v vodi. Uspe{nost dispergiranja smo spremljali z ugotavljanjem masnega dele`a delcev, dispergiranih v vodi. Vsebnost delcev je odvisna od pH-vrednosti pri kateri se citronska kislina adsorbira na povr{ini in koncentracije raztopljene citronske kisline. Tretja stopnja postopka je vsebovala prevla~enje nanodelcev s siliko, ki je potekalo v vodni suspenziji. Klju~nega pomena za uspe{no prevla~enje nanodelcev s siliko je stabilnost suspenzije nanodelcev oble~enih s citronsko kislino. Klju~ne besede: maghemit, nanodelci, citronska kislina, silika, vodna suspenzija In this paper we describe the synthesis of maghemite nanoparticles coated with silica. The process was divided into three separate steps. In first step we have synthesized particles with co-precipitation of Fe2+ and Fe3+ ions. In the second step the particles were coated with citric acid in order to prevent agglomeration and dispersed in water. The success of dispersion was determined by weight percent of dispersed particles in water. The quantity of dispersed particles depends on pH value at which the citric acid was adsorbed on the surface of particles and on the concentration of dissolved citric acid. The third step was coating the nanoparticles with silica, which took place in the water suspension. The key factor for successful coating of nanoparticles with silica is the stability of suspension of nanoparticles coated with citric acid. Key words: maghemite, nanoparticles, citric acid, silica, water suspension 1 UVOD V zadnjih letih se velika pozornost namenjena uporabi nanodelcev v medicini. [e posebej zanimiva je uporaba magnetnih nanodelcev, ker nanje lahko vplivamo z magnetnim poljem. Zaradi njihovih magnet- nih lastnosti lahko take nanodelce spremljamo ali z njimi upravljamo na daljavo. Uporabljamo jih lahko bodisi za diagnosti~ne namene (na primer za pove~evanje kontrasta pri slikanju z NMR-tehniko) ali terapevtske tehnike, kot sta na primer magnetna hipertermija in ciljani vnos zdravilnih u~inkovin1. Pogoj za njihovo uporabo v medicini je poleg njihove "nano"-velikosti in zadovoljivih magnetnih lastnosti tudi njihova nestrupe- nost in specifi~ne povr{inske lastnosti. Kot magnetni material za uporabo v obliki nanodelcev v medicini se uporablja predvsem maghemit (-Fe2O3), ki velja za nestrupen material2. Nanodelci se za medicinske namene navadno uporabljajo v obliki suspenzij v fiziolo{kem mediju na vodni osnovi. Nanodelci morajo ostati disper- girani v mediju tudi ob ve~jih spremembah v svojem okolju, kot so na primer spremembe v ionski mo~i in pH-vrednosti medija. Da bi magnetne delce lahko disper- girali v teko~ini, morajo izkazovati superparamagnetne lastnosti. Take lastnosti imajo pri sobni temperaturi delci magnetnega materiala, manj{i od neke kriti~ne velikosti, ki je navadno manj{a od 15 nm. Superparamagnetni nanodelci zunaj magnetnega polja ne ka`ejo spontane magnetne polarizacije. Med njimi ni magnetnih inte- rakcij, ki bi povzro~ale magnetno aglomeracijo. Stabilne koloidne suspenzije superparamagnetnih nanodelcev imenujemo tudi magnetne teko~ine. Za uporabo v medicini je treba na povr{ino nanodelcev vezati razli~ne biolo{ke u~inkovine. Vezavo u~inkovin na povr{ino nanodelcev dose`emo prek funkcionalizacijskega sloja molekul, vezanih na povr{ino delcev. Te molekule zagotavljajo funkcionalne skupine za kemijsko vezavo razli~nih biolo{kih u~inkovin, hkrati pa zagotavljajo stabilnost suspenzije nanodelcev. Izka`e se, da je tako za stabilnost nanodelcev v fiziolo{kem mediju kot tudi za u~inkovitost vezave biolo{kih u~in- kovin na nanodelce klju~nega pomena jakost vezi med povr{ino in funkcionalizacijskimi molekulami. Povr{ina nanodelcev oksida (maghemita) je relativno inertna in navadno ne omogo~a mo~ne kovalentne vezave molekul. Zato je delce potrebno prevle~i s plastjo silike, ki ima na povr{ini silanolne OH-skupine, ki omogo~ajo nadaljnjo vezavo u~inkovin s kovalentno vezjo. Siliko je na povr{ino delcev mogo~e vezati s sol-gel-postopkom3,4,5 ali z Na-silikatom6. Materiali in tehnologije / Materials and technology 41 (2007) 2, 103–107 103 UDK/UDC 537.622:532/533:620.1 ISSN 1580-2949 Izvirni znanstveni ~lanek/Original scientific article MTAEC9, 41(2)103(2007) Pri na{em delu smo sintetizirali maghemitne nano- delce in jih prekrili s tanko plastjo silike za kasnej{o funkcionalizacijo. 2 EKSPERIMENTALNO DELO Maghemitne delce smo sintetizirali s koprecipitacijo Fe2+ (0,027 mol/L) in Fe3+ (0,023 mol/L) s koncen- triranim amoniakom (25 %) pri sobni temperaturi. Postopek je potekal v dveh stopnjah. V prvi smo dvignili pH-vrednost raztopine `elezovih ionov na pH = 3 in vzdr`evali konstantno pH-vrednost 30 min. V tej stopnji se obori Fe(OH)3. V drugi stopnji smo dodali koncen- triran amonijak, da smo dosegli pH = 11,6. V tej stopnji se `elezov hidroksid oksidira z zra~nim kisikom in nastane spinelni produkt. Tako sintetizirane delce smo nato magnetno lo~ili in jih trikrat sprali z amoniakalno raztopino s pH > 10,5. Pri tej pH-vrednosti imajo delci na povr{ini visok negativni naboj, kar se ka`e kot visok -potencial, ki prepre~uje njihovo mo~no agregiranje. -potencial sintetiziranih delcev kot tudi delcev, prevle~enih s citronsko kislino in s siliko smo dolo~ili z elektrokineti~nimi meritvami mobilnosti delcev. Meritve so bile opravljene z zetametrom (Brookhaven Instru- ments Corp., ZetaPALS). Nanodelce smo karakterizirali z uporabo rentgenske pra{kovne difrakcije (XRD) (Bruker AXS, D4 ENDEAVOR) in presevnega elektronskega mikroskopa (TEM) (JEOL 2010F). Velikost delcev smo dolo~ili iz TEM-posnetkov in z uporabo Scherrerjeve metode, ki temelji na {iritvi rentgenskih uklonov v XRD. Pri tem smo uporabljali ra~unalni{ki program Diffracplus Topas™. Pred prekritjem nanodelcev s hidrolizo tetraetoksi- silikana (TEOS) smo jih dispergirali v vodi, da smo omogo~ili prekritje posameznih delcev, in ne njihovih aglomeratov. Za dispergiranje v vodi smo delce prevlekli s citronsko kislino7. V me{anico vode in nanodelcev smo dodali razli~ne koli~ine citronske kisline ( = 0,005–0,04 g/mL) in uravnali pH-vrednost z dodatkom amonijaka (25 %). Vrednosti pH med prekrivanjem s citronsko ki- slino so bile med 2 in 11. Me{anico smo segreli na 80 °C in jo intenzivno me{ali 90 min. Med tem se je citronska kislina verjetno kemijsko vezala na povr{ino nanodelcev. V naslednji stopnji procesa smo me{anico {ibko aglomeriranih delcev ohladili in delce med intenzivnim me{anjem dispergirali, navadno s hkratno spremembo pH-vrednosti. Neaglomerirani delci v suspenziji so dolgoro~no stabilni in se ne posedajo niti v zmernem gradientu magnetnega polja. Aglomerate smo iz suspen- zije izlo~ili s centrifugiranje na 5000 r/min, 5 min. Suspenzija po centrifugiranju je stabilna magnetna teko~ina. Kot merilo za uspe{nost postopka prekritja nanodelcev s citronsko kislino in njihovega dispergiranja smo uporabili vsebnost nanodelcev v stabilni magnetni teko~ini. Vsebnost delcev je bila dolo~ena z `aroizgubo. Magnetne nanodelce, dispergirane v vodi, smo pre- vlekli s siliko. Magnetni teko~ini smo dodali etanol in TEOS (-Fe2O3 : TEOS = 1 : 23; TEOS : EtOH = 1 : 1, masna razmerja), segreli na 75 °C in pustili 24 h, da je TEOS v celoti hidroliziral. Debelino prevleke silike na nanodelcih smo kontrolirali s koli~ino TEOS-a. Kvaliteto prevleke silike na nanodelcih smo opazo- vali s visokolo~ljivostnim presevnim elektronskim mikroskopom (HRTEM). Izkazalo se je, da je za nastanek prevleke na posameznem delcu bistvenega pomena stabilnost njihove suspenzije, zato smo posebno pozornost namenili pripravi stabilnih magnetnih teko~in. 3 REZULTATI IN DISKUSIJA Slika 1 prikazuje rentgenski pra{kovni difraktogram sintetiziranih delcev. Ta prikazuje {iroke uklone, ki jih lahko pripi{emo spinelni fazi, tj. maghemitu. [irjenje uklonov XRD je posledica majhnosti krista- litov. Iz {irjenja XRD-uklonov je bila dolo~ena pov- pre~na velikost kristalitov 13 nm. Slika 2 prikazuje TEM-sliko sintetiziranih delcev. Delci, prikazani na sliki 2, so dobro kristalizirani in imajo ozko porazdelitev velikosti. Povpre~na velikost, izra~unana s slike 2, je S. ^AMPELJ ET AL.: SINTEZA MAGNETNIH NANODELCEV, FUNKCIONALIZIRANIH S TANKO PLASTJO SILIKE 104 Materiali in tehnologije / Materials and technology 41 (2007) 2, 103–107 Slika 2: TEM slika sintetiziranih delcev Figure 2: TEM image of synthesized particles Slika 1: XRD sintetiziranih delcev Figure 1: XRD of synthesized particles 13,7 ± 2,9 nm, kar se dobro sklada z velikostjo, dolo~eno z XRD. Delci so po sintezi aglomerirani in kot taki neprimerni za prevla~enje s siliko. Za prevla~enje s siliko je treba prepre~iti aglomeracijo delcev, torej pripraviti stabilno suspenzijo v vodi. To se lahko dose`e elektrostatsko, steri~no ali elektrosteri~no. V na{em primeru smo aglomeracijo prepre~ili z vezavo citronske kisline na povr{ino. Kislina, vezana na povr{ini, steri~no prepre~uje aglomeracijo nanodelcev, hkrati pa zagotovi na njihovi povr{ini visok naboj, ki prispeva k elektro- statski stabilizaciji suspenzije (elektrosteri~ena stabiliza- cija). Uspe{nost vezave citronske kisline na povr{ino delcev je razvidna iz premika izoelektri~ne to~ke. Slika 3 prikazuje odvisnost -potenciala na nanodelcih po sintezi, delcev, prevle~enih s citronsko kislino in delcev, prevle~enih s siliko, v odvisnosti od vrednosti pH. Izoelektri~na to~ka neprekritih delcev je pri pH = 7,2 (polna linija), izoelektri~na to~ka delcev po prekritju s citronsko kislino pa se premakne v kislo obmo~je pH = 2 (~rtkana linija). Premik izoelektri~ne to~ke nastane zaradi spremembe na povr{ini delcev. Pri sintetiziranih delcih so na povr{ini kisikovi ioni. Pri delcih, prevle- ~enih s citronsko kislino, pa so na povr{ini OH-skupine citronske kisline. Za stabilno magnetno teko~ino moramo zagotoviti na povr{ini vseh nanodelcev optimalno koli~ino vezane citronske kisline. Merilo za stabilnost suspenzije je vsebnost magnetnih delcev v pripravljeni magnetni teko~ini po centrifugiranju. Koli~ina dispergiranih delcev je odvisna od ve~ faktorjev, ki so povezani z vsebnostjo in vezavo citronske kisline na njihovo povr{ino, in od pH-vrednosti, pri kateri nanodelce, prekrite s citronsko kislino, dispergiramo. Nanodelci, prekriti s citronsko kislino pri nizkih pH-vrednostih, so se le ste`ka dispergirali. Izkazalo se je, da se koli~ina nanodelcev, ki preide v stabilno suspenzijo, mo~no pove~a, ~e se pH-vrednost dvigne na pribli`no 10. Iz krivulje odvisnosti -potenciala od pH-vrednosti (slika 3) izhaja, da s pove~anjem pH-vrednosti -potencial raste in s tem se pove~uje stabilnost suspenzije. Delce, prevle~ene s citronsko kislino, smo zato v vseh nadalj- njih poizkusih dispergirali pri pH-vrednosti 10. Vsebnost in na~in vezave citronske kisline na povr{ini nanodelcev pa je odvisna od pH-vrednosti, pri kateri je potekala adsorpcija citronske kisline na povr{ini delcev in tudi od koli~ine citronske kisline v suspenziji. Slika 4 prikazuje masni dele` magnetnih delcev v magnetni teko~ini, dolo~eni z `aroizgubo v odvisnosti od pH-vrednosti med adsorpcijo citronske kisline na njihovi povr{ini. S slike 4 je razvidno, da je najve~ja vsebnost magnetnih delcev pri pH-vrednosti 5,2. Pri pH-vred- S. ^AMPELJ ET AL.: SINTEZA MAGNETNIH NANODELCEV, FUNKCIONALIZIRANIH S TANKO PLASTJO SILIKE Materiali in tehnologije / Materials and technology 41 (2007) 2, 103–107 105 Slika 5: Masni dele` maghemitnih delcev v magnetni teko~ini v odvisnosti od koncentracije citronske kisline v suspenziji med procesom njene adsorpcije na povr{ini nanodelcev Figure 5: Weight percent of maghemite particles in ferrofluid as a function of concentration of citric acid in suspension during the adsorption of citric acid on the surface of nanoparticles Slika 3: ζ-potencial maghemitnih delcev (polna linija), maghemitnih delcev, prevle~enih s citronsko kislino (~rtkana linija), in maghe- mitnih delcev, prevle~enih s siliko (pik~asta linija) v odvisnosti od pH-vrednosti Figure 3: ζ-potential of maghemite particles (full line), maghemite particles coated with citric acid (dashed line) and maghemite particles coated with silica (dotted line) as a function of pH value Slika 4: Masni dele` maghemitnih delcev v magnetni teko~ini v odvisnosti od pH-vrednosti, pri kateri se je citronska kislina adsorbirala na povr{ini nanodelcev Figure 4: Weight percent of maghemite particles in ferrofluid as a function of pH value at which the citric acid has adsorbed on the surface of nanoparticles nostih ve~jih od 5,2 je na povr{ini delcev premajhen pozitiven naboj, nad pH = 7 pa je celo negativen, kar prepre~uje adsorpcijo citronske kisline. Zato je koli~ina adsorbirane kisline na povr{ini zanemarljiva, posledica pa je majhna vsebnost magnetnih delcev v suspenziji. Pri pH-vrednostih, manj{ih od 5,2, je vsebnost delcev manj{a zaradi sen~enja naboja na povr{ini delcev in pove~ane topnosti oksidnih delcev v raztopini citronske kisline. Slika 5 prikazuje masni dele` magnetnih delcev v stabilni suspenziji v odvisnosti od koncentracije citron- ske kisline med adsorpcijo. Ker je adsorpcija ravnote`ni proces med neadsorbirano kislino, raztopljeno v vodi, in kislino, adsorbirano na povr{ini delcev, je koli~ina adsorbirane kisline odvisna od koncentracije kisline v raztopini. Ve~ja kot je koncentracija citronske kisline v raztopini, ve~ji je dele` adsorbirane kisline in posledi~no ve~ja koli~ina stabilno dispergiranih delcev. Koli~ina dispergiranih delcev se zmanj{a pri koncentracijah citronske kisline, vi{jih od 0,02 g/mL, zaradi sen~enja naboja na povr{ini delcev. S pove~anjem koncentracije ionov v suspenziji se zmanj{uje naboj na povr{ini delcev. Ko je dose`ena kriti~na koagulacijska koncen- tracija ionov, se delci za~nejo agregirati. Delce prevle~ene s citronsko kislino, in dispergirane v vodi smo pri nadaljnjem delu prevlekli s siliko. Prevleka silike na nanodelcih je potekla s hidrolizo TEOS-a. Nukleacija silike mora potekati heterogeno na povr{ini nanodelcev, in ne homogeno v sami raztopini. Zato morajo biti delci stabilno dispergirani. Za uspe{no prevla~enje je klju~nega pomena tudi katalizator, ki je v na{em primeru amonijak. Ta je v suspenziji iz pred- hodnega postopka prekritja s citronsko kislino in se ka`e v vrednosti pH. Optimalni pogoji za prekritje delcev so pri pH-vrednostih 10, torej pri pogojih, kjer je suspenzija najbolj stabilna in je hkrati prisotna nizka koncentracija amonijaka. ^e suspenzija ni bila popolnoma stabilna, je pri{lo do prekritja aglomeratov nanodelcev. Slika 6 prikazuje TEM-posnetek delcev, oble~enih s siliko pri nevtralni pH-vrednosti. Pri pH-vrednostih okoli nevtralne to~ke je hidroliza potekla nehomogeno, s siliko so se prevlekli manj{i aglomerati nanodelcev. Hkrati je vidna tudi nehomogena debelina silike na aglomeratih delcev, kar nakazuje, da nukleacija silike ni potekala homogeno. Prisotnost aglomeratov maghemitnih delcev nakazuje na nestabilno suspenzijo delcev med procesom prekritja. Slika 7 prikazuje HRTEM-posnetek delcev, ki so bili prekriti z dodatkom TEOS-a 42 µL na 5 mL suspenzije maghemitnih nanodelcev, ki je imela pH-vrednost 10,1. Na povr{ini delcev je vidna zelo tanka ( 1 nm) plast amorfnega materiala. EDS-analiza posameznih delcev je pokazala relativno konstanten signal silicija, zato lahko sklepamo, da je amorfna plast na povr{ini delcev silika. Slika 8 prikazuje TEM-posnetek delcev, prekritih s siliko, v primeru, ko smo dodali 1,34 mL TEOS-a na 5 mL suspenzije mag- hemitnih nanodelcev. Vidna je debelej{a plast silike na posameznih delcih. Vsi nanodelci v vzorcu imajo na povr{ini plast silike S. ^AMPELJ ET AL.: SINTEZA MAGNETNIH NANODELCEV, FUNKCIONALIZIRANIH S TANKO PLASTJO SILIKE 106 Materiali in tehnologije / Materials and technology 41 (2007) 2, 103–107 Slika 8: TEM-posnetek maghemitnih delcev, prevle~enih s siliko Figure 8: TEM image of maghemite particles coated with silica Slika 6: TEM-posnetek delcev, prevle~enih s siliko, pri nevtralni pH-vrednosti Figure 6: TEM image of particles coated with silica at neutral pH value Slika 7: HRTEM-posnetek maghemitnih delcev, prevle~enih s tanko plastjo amorfnega materiala Figure 7: HRTEM image of maghemite particles coated with thin layer of amorphous material debelo pribli`no 10 nm. Poleg prekritih maghemitnih delcev je v vzorcu tudi veliko majhnih nanodelcev same silike. Sestava vzorca ka`e, da je pri{lo do hidrolize TEOS-a ne samo na povr{ini delcev, vendar tudi s homogeno nukleacijo silike v raztopini. Rezultat ka`e, da je uporabljeni na~in prekritja nanodelcev primeren, treba pa bo {e optimizirati pogoje, da bi se znebili delcev silike, ki ni prisotna kot prevleka na maghemitnih nanodelcih. 4 SKLEP Uporabljena metoda prekritja superparamagnetnih nanodelcev s siliko se je izkazala za u~inkovito. Klju~nega pomena za uspe{no prekrivanje posameznih nanodelcev je stabilnost na vodi osnovanih magnetnih teko~in. Sama metoda {e ni optimizirana, saj so poleg maghemitnih delcev prekritih s siliko prisotni tudi delci silike, ki ni prisotna kot prevleka na maghemitnih nanodelcih. 5 LITERATURA 1 I. [afaík, M. [afaíkova, Monatshefte für Chemie, 133 (2002), 737–759 2 U. Häfeli, W. Schüt, J. Teller, M. Zborowski, Scientific and clinical applications of magnetic carriers, Plenum Press, New York, 1997 3 W. Stöber, A. Fink, J. Colloid Interface Sci., 26 (1968), 62–69 4 A. Lu, Y. Yin, B. T. Mayers, Y. Xia, Nanoletters, 2 (2002), 183–186 5 Q. Liu, Z. Xu, J. A. Finch, R. 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