VSEBINA – CONTENTS IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES A new version of the theory of ductility and creep under cyclic loading Nova verzija teorije o duktilnosti in lezenju pri cikli~ni obremenitvi L. B. Getsov, M. G. Kabelevskiy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Thermoelectrical properties of a monocrystalline Al64Cu23Fe13 quasicrystal Termoelektri~ne lastnosti monokristalnega kvazikristala Al64Cu23Fe13 I. Smiljani}, A. Bilu{i}, @. Bihar, J. Lukatela, B. Leonti}, J. Dolin{ek3, A. Smontara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 265 Faze v kvazikristalni zlitini Al64,4Cu22,5Fe13,1 Phases in a quasicrystalline alloy Al64,4Cu23,5Fe13,1 T. Bon~ina, B. Markoli, I. An`el, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 271 Hydrogen absorption by Ti–Zr–Ni-based alloys Absorpcija vodika v zlitinah Ti–Zr–Ni I. [kulj, A. Kocjan, P. J. McGuiness, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 279 Microstructural evaluation of rapidly solidified Al–7Cr melt spun ribbons Ovrednotenje mikrostrukture hitrostrjenih trakov Al-7Cr P. Jur~i, M. Dománková, M. Hudáková, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283 The influence of different waste additions to clay-product mixtures Vpliv razli~nih odpadkov na izhodno surovino za proizvodnjo ope~nih izdelkov V. Ducman, T. Kopar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289 Electrochemical and mechanical properties of cobalt-chromium dental alloy joints Elektrokemijske in mehanske lastnosti razli~nih spojev stelitne dentalne zlitine R. Zupan~i~, A. Legat, N. Funduk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Development of microstructure of steel for thermal power generation Razvoj mikrostrukture jekel za termi~no generacijo energije Kvackaj T., Kuskulic T., Fujda M., Pokorny I., Weiss M., Bevilaqua T. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 301 MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH, PORTORO@, 8.–10. OKTOBER, 2007 YOUNG SCIENTISTS AWARDS, 15th CONFERENCE ON MATERIALS AND TECHNOLOGY, PORTORO@, 8–10 OCTOBER, 2007 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 305 DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES . . . . . . . . . . . 319 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 41(6)255–330(2007) MATER. TEHNOL. LETNIK VOLUME 41 [TEV. NO. 6 STR. P. 255–330 LJUBLJANA SLOVENIJA NOV.-DEC. 2007 L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP UNDER CYCLIC LOADING NOVA VERZIJA TEORIJE O DUKTILNOSTI IN LEZENJU PRI CIKLI^NI OBREMENITVI Leonid B. Getsov1, M. G. Kabelevskiy2 1St. Petersburg State Technical University, Russia 2ZNIITMASh, Moscow, Russia guetsovonline.ru Prejem rokopisa – received: 2007-04-06; sprejem za objavo – accepted for publication: 2007-10-18 The model of deformation for isotropic metallic materials aimed at obtaining an increased accuracy for forecasting their behavior during complex cyclical loading, in particular cyclic loading where a significant creep role is played in the processes of creep, was developed. As with the majority of models, this new model has applicability limitations and the reliability of use for calculations is acceptable only for cases of small differences in loading from the proportionality. Key words: metallic materials, cyclic loading, creep deformation, modelling, reliability Razvit je bil model o deformaciji izotropnih kovinskih materialov s ciljem, da bi se dosegla ve~ja natan~nost pri napovedi vedenja pri kompleksni cikli~ni obremenitvi, {e posebej v primeru pomembne vloge procesov lezenja. Kot pri ve~ini modelov ima tudi novi model omejitve pri uporabi in zanesljivost uporabe izra~unov je sprejemljiva samo za primer, ko se obremenitev malo odmika od proporcionalnosti. Klju~ne besede: kovinski material, cikli~na obremenitev, deformacija z lezenjem, modeliranje, zanesljivost, natan~nost 1 INTRODUCTION Several models have been developed so far 1, and some of them are in commercial use as software packages, for example, ANSYS, MARC, NASTRAN, ABAQUS, LUSAS, LS-DYNA, COSMOS, ALGOR. However, these models fail to adequately describe the case of complex cyclic loading, when creep processes also play an important role. In analyzing the lines of access to the development of a theoretical explanation for the straining under cyclic non-isothermal loading, which is necessary for practical calculations of the strain-stressed state (SSS), the authors stood at a crossroads. It was possible to use structural and physical models, which made it possible to describe a wide range of peculiarities of a material’s behaviour under complex loading using a rather small number of experimental material parameters 2. It is worth noting that the deformation and the instantaneous plastic deformations are not separated and their interconnection is included as a property of the developed model. Analo- gous models have not provided sufficient reliability for a quantitative calculation since the monotonic change in the material properties differs from the experimental values. It was also clear that a modification of the traditional models for plastic flow and the different creep theories, as applied to specific loadings to achieve good accuracy with the calculation, would require a large number of basic experiments to obtain an acceptable fit for a description of real material behaviour. We chose the first solution because of its obvious advantages. The model of deformation for isotropic metallic materials was designed to make a very accurate prediction of their behavior. This article looks at the case of complex cyclical loading. 2 THEORY 1. The variations in non-isothermal theories of plastic flow and of the theory of work hardening during creep will be included in the analysis 3, allowing us to consider the mutual effect of the two forms of deformation within the framework of the traditional approach. In rating the correctness of these proposals, we will start from the necessity of fulfilling the following requirements: a) a description on a non-isothermal cyclic deformation; b) a consideration of the cyclic instability of the mate- rial properties; c) a description of the conditions of deformation for complex loading, in particular of alternating sign; d) a consideration of the mutual influence of time- dependent and time-independent deformation. The approach based on the separation of irreversible deformation into time-dependent and time-independent has a physical basis. 2. We will assume that the total material deformation consists of the elastic deformation rij, the creep defor- mations pij, and the plasticity ij, 4, thus: eij = rij + pij + ij (1) In the formulation of the rules of deformation we will consider the effect of accumulated plastic deformation Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 257 UDK 539.42:539.4:669 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)257(2007) on the creep and of the temperature-time prehistory on the elasto-plastic properties. Although the assumption that only the second invariant of the tensor of stresses enters into the relationship for the increase of the defor- mation and stresses, it is only a special case of the relationship recommended for the description of the complex stressing of bodies 5. At present this is a traditional approach. This conclusion is related to the fact that with comparatively small plastic deformations, it provides, as a rule, quite good agreement with experi- mental data for many cases when plastic deformation occurs during stress that is not very different from the uniaxial. Having accepted the hypothesis of work hardening that, in particular, means the neglecting of processes of the reverse elastic secondary effect and assuming that the recurrence of loading and preliminary plastic defor- mation affects only the scalar properties, we may write the equation for creep rate for cyclic loading with an alternating sign: p F p p pij n n n ( ) ( ) ( ) ( ) ( ) ( ) ( ) ( , , ..., , , , ,= − 0 1 0 1 1 ε ε ..., ε  , , , , )( )ε σ σ n ijn T S (2) Here, Sij are the components of the deviator of the stresses, σ is the intensity of the stresses, and ε ( )n and p n( ) are the intensities of the plastic deformation and the creep deformation determined from the equations p p p tn ij ij t t n n ( ) /  = ⎛ ⎝⎜ ⎞ ⎠⎟ − ∫ 2 3 1 1 2 d (3) ε ε ε( ) /  n ij ij t t n n t= ⎛ ⎝⎜ ⎞ ⎠⎟ − ∫ 2 3 1 1 2 d (4) where n is the number of the half cycle p pij ij∆ ≥ 0 (5) Experiments relating to uniaxial stressing 6 showed that by counting the creep deformations during cyclic loading from the start of a half cycle, the curves of the irreversible deformations accumulated during the non-steady creep are similar to the creep curves for the initial condition. In general, the coefficient of similarity depends on the number of the half cycle, the time, the amount of creep deformation accumulated per half cycle, and the temperature. The rate of steady creep is virtually independent of the number of half cycles. Comparatively small previous creep and instanta- neous plastic deformations 0.2–5.0 % may have a sub- stantial influence on the creep rate. Plastic deformations of the opposite sign accelerate the creep of high-tempe- rature materials of different classes (the Bauschinger effect in creep), while plastic deformations of the same sign can accelerate or retard creep, depending upon their size and the type of material. In cases when the material is submitted to elasto- plastic steady-stage creep the deformation of the opposite sign, further creep, as a rule, starts with the non-steady stage. An analysis of the experimental data on cyclic creep makes it possible to select two variations of the concre- tization of Eq. (2). The first approach is based on the use of the graph in Figure 1, based on the assumption that the effect of plastic deformations on the creep rate may be considered with an appropriate change of the value of creep defor- mation in the relationship p = f(p,). In this case, in the plastic deformation pl for the total deformation p, corresponding to the point k, we have a creep rate that corresponds to the point b, the point d, and not to the point b. For p = 0 the value of p corresponds to the point e. The material cyclic creep instability will be considered with the use of the function  of the number of cycles 5. In this case  ( ~, )p f p= σ (6a) Here [ ]~ ~ ( ~ ( ) ( , p p n p p  = + − −1   ψ ε ψ ε ϕ−) where pn is the intensity of the creep deformation accu- mulated as a result of non-steady creep; ψ1 , ψ2 , and ε are functions taking into consideration the effect of plastic deformations and the number of cycles and satisfying the following conditions: for ε+ = 0 ~ψ 1= ; for ε− = 0 ~ψ  1= ; for n = 1 ~ϕ = 1. ε+ and ε− are the sum L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... 258 Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 Figure 1: Graph for taking into consideration the effect of plastic deformation on the creep rate: OL is the creep curve p = f(); OM is the tangent to the creep curve at  = 0; a, b are the relations between the creep rate and the accumulated creep deformation; ae is the tangent to the curve ab at the point a; pn is the creep deformation in the non-steady stage. The arrows show the method of determining the effect of pl on p. Slika 1: Grafikon za upo{tevanje vpliva plasti~ne deformacije na hitrost lezenja. OL-krivulja lezenja p = f(); OM-tangenta na krivuljo lezenja v to~ki  = 0; a, b-odnos hitrosti lezenja in nakopi~ene defor- macije z lezenjem; ae-tangenta na krivuljo ab v to~ki a; pn-defor- macija z lezenjem v nestabilnem stanju. Pu{~ice prikazujejo metodo dolo~itve vpliva pl na p. of the intensities of the plastic deformations in those half cycles where ∆ε ij n ij np( ) ( ) > 0 and ∆ε ij n ij np( ) ( ) < 0; ∆ ( )ε ij n ij t t e t n n = − ∫ d 1 (7) The problem of the use of Eq. (6a) is related to the correctness of the extrapolation of the relationship p = f(p,) in the area of negative values of p. The second possible approach to the treatment of existing data provides, as a solution of Eq. (2), the following expression [ ]  ( , , (( ) min ( ) minp p n t p p pn n= + −−ϕ 1 0 (6b) where  minp is the steady-stage creep rate;  minp = f(σ,T); p0 is the initial creep rate of the material, calculated according to the theory of work hardening; p0 = f(σ,T,p); (n,t,p(n–1) is a function considering the effect of the number of cycles on the non-steady creep; ψ ε τ+1 ( , ) and ψ ε τ−2 ( , ) are functions considering the effect of plastic deformations; t and  are the times counted from the start of the cycle and the moment of the start of the plastic deformation. The form of the functions  minp and p0 may be deter- mined in creep tests under conditions of the uniaxial stressed condition at constant values of the stresses and the temperature. In Equations (2) and (6) the effect of the temperature at which the instantaneous plastic deformation was accumulated on creep rate is neglected. Such an effect is possible; however, the existing experimental results indicate that it is insignificant. 3. We will describe the instantaneous plastic defor- mations for non-isothermal cyclic deformation of alter- nating sign with the incremental theory of thermo- plasticity with a piecewise-spherical surface 3, modified by applying a relationship for the accumulated creep deformation. Let us assume that the original material is isotropic in the space of the deviators of stresses and it has the paths of the cyclic load for each point of the body given by the cone with a small spatial angle  (Figure 2). The area where the vector deviator of stresses during the whole load cycle must be found is cross-hatched. For the k-th half cycle we have f S S r p p k ij ij k k= − −( , , ..., , , , , .. ( ) ( ) ( ) ( ) ( ) ( ) ε ε ε ε0 1 1 0 1 ., , , )( )p k Tk = 0 (8) where the plastic work hardening does not depend on the temperature prehistory. As for creep, the number of the half cycle is increased per unit with a breakdown in the condition ε εij ij∆ ≥ 0 (9) where ∆ ∆ε = 2 3 εij ijS R (10) ∆ε is the intensity of the increase in deformations. A determination of the radius of the surface of flow R is significantly easier, applying for the material and temperature T, the generalized diagrams of cyclic defor- mation representing the relationships between the increases in stresses and deformations counted from the start of a given half cycle and independent of the amplitude of the cycle deformation. If we assume that between the increase of stresses and the accumulation of plastic deformation in a given half cycle at a given temperature and constant values of p a single relationship exists, we can determine the value of the radius of the surface of flow as an algebraic sum of its value in the zero half cycle and its increase in the subsequent half cycles: R R Rk i k i= + = ∑0 1 ∆ (11) The treatment of the results of the experiments made for uniaxial stressing showed that a small creep deformation increases the yield strength of a material in cases when the directions of deformation coincide (Figure 3), and decrease it when the directions of deformation in creep and in instantaneous elastoplastic deformation are opposite. We can assume that the creep deformation p influ- ences the value of the radius of the surface of flow by an additional plastic deformation of ~ ( / )ε ε ε= p   , where ε c and ε t are the deformation capacity of the material for creep and short-term tension. In this case, the expression for the surface of flow is written in the form: L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 259 Figure 2: Graph of the surface of flow during cyclic loading: f) original surface of flow; fk, fk+1) parts of the surface after the k-th and the (k+1)-th half cycles; R'k+1 and R''k+1) radius of the surface of flow in the (k+1)-th half cycle with the related increase of stresses in the k-th half cycle. Slika 2: Grafikon za povr{ino lezenja pri cikli~ni obremenitvi: f) izvirna povr{ina lezenja; fk, fk+1) deli povr{ine po k in (k+1) polo- vi~nem ciklu; R'k+1 in R''k+1) polmer povr{ine lezenja pri (k+1) polovi~nem ciklu z od njega odvisnim pove~anjem napetosti pri k polovi~nem ciklu f S S r p p k ij ij c t c t k = − + + ⎛ ⎝ ⎜ ε ε ε ε ε ε ε ( ) ( ) ( ) ( ) , , ... ..., 0 1 1 + ⎞ ⎠ ⎟ = p k T k c tε ε , , 0 (12) A comparison of the calculated curve (the dashed line in Figure 3) with the experimental (the solid lines) confirms the acceptability of this assumption. (The calculated curve for p = 0,5 % was obtained from the curve for p = 0 by replacement to the left by the value  = (p/c)t.) This form of surface flow makes it possible to explain, in particular, the phenomenon of plastic defor- mations for a creep test cycle with alternating sign and stresses lower than the elastic limit of the original material 6. Equations (1), (6), and (12), the expanded equations of the equilibrium and the consistency of the defor- mation of uniform medium and also the necessary boundary conditions make it possible to calculate the stressed-strained condition of a body in an arbitrary program of cyclic loading and heating by a step method. At the same time, Eq. (6) and (12) satisfy the require- ments given earlier. It was shown in 7 that the calculation for stresses and deformations during the loading stage may be developed as a solution to the problem of the deformation theory of plasticity with varying the parameters of elasticity 4. The process of determining successive approximations for the stresses and deformations is carried out with the separation of the deformation into elastic and an increase in the instantaneous plastic deformation and creep according to Eqs. (6) and (12) using the method of successive approximations. As a parameter of change of materials properties related to its cyclic instability by a change of cyclic loading in place of the number of semi-cycles, the path of cyclic creep, 1, and the path of cyclic plastic deformation 2, should be accepted. Let us assume that the cyclic creep and its increment are determined by the equations: λ ε ε ε ε ε ε ε ε 1 0 52 3 2 3 = − = ⋅ ⋅ = ⋅ ⋅ ∫ d d d dp p p pij pij p pij / / ; ( ) ; ( . pij pij pij pi p p p / p p p / p ) ; ( ) ; ( . . 0 5 1 0 52 3 2 3 λ = − = ⋅ ⋅ = ⋅ ∫ d d d d j pijp⋅ ) .0 5 (13) ∆λ λ λi i k i k= − ≥−( ) ( )1 0 (14) Here, the following requirements must be met: ∂ ∂ λ τ i k( ) > 0 at t < tk ; ∂ ∂ λ τ i k( ) = 0 at t > tk (15) By satisfying conditions (14) and (15), the value of n increases by unity. The increments of the inelastic deformation and the values of the intensity of the inelastic deformation are determined from the equations: d d dε εij pij ijp ℵ = + ; ε ε εℵ ℵ ℵ= ⋅ ⋅( ) .2 3 0 5/ ij ij 4. An estimation procedure for the material characte- ristics and design procedure of creep curves for some typical examples of uniaxial loading with cyclic creep at varying temperatures can be applied. The proposed method can be used to calculate the strains at all three stages of creep for heat-resistant steels and alloys with an arbitrary law of change in the stress and temperature at the working temperatures at which the material is structurally stable. The method cannot be used for calculations at the third stage of creep in materials that are fractured after necking or in cases of compression or alternating loading of materials with a highly anisotropic initial creep resistance (in tension and compression). The duration of the creep process during one cycle may range from a minute to hundreds of hours. In developing the method, we analyzed data on creep and stress relaxation in 20 grades of heat-resistant steels and alloys in a uniaxial stress state. 3 COVERNING EQUATIONS The method is based on a creep theory of the following type 8 p = f(,p,T,pl,1) (17) where λ = −∫ ( )d dp p is the path of cyclic creep. Tests involving a single loading are approximated with the following analytical formula: p = F(,T,pl,t) (18) L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... 260 Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 Figure 3: Effect of creep on the resistance of 25Kh2M1F steel to instantaneous deformation at 550 °C. Full lines are the experimental results; dashed line is the calculated. Slika 3: Vpliv lezenja na odpornost jekla 25Kh2M1F proti hipni deformaciji pri 550 °C, polne ~rte – eksperimentalne meritve; ~rtkane ~rte – izra~unano The relation for creep rate in the same tests can be determined by differentiating Eq. (18): p F t = d d = Φ(,T,pl,t) (19) The behaviour of the material by complex loading programs is assumed to be described by creep theory (17); thus, it is obvious that the required relationship p = f(,p,T,pl,) should be obtained by excluding t from Eqs. (18) and (19). However, this cannot be done analytically in a general form with the present form of Eq. (19); it is therefore proposed that in the solution of any specific problem, the value of function 17 should be found nume- rically by excluding t from Eqs. (18) and (19) at each step of the integration over time. For alternating loading and the absence of instantaneous plastic strains, the specific form of Eq. (18) to describe the creep curves in the first and second stages is: [ ]p A C t B tk l m= − − +σ σ σ1 exp( ) (20) where A,B,C,k,l,m are coefficients that are constant for a given test temperature. To describe the third creep stage, we replace the stresses in Eq. (20) with the ratio /(l – p/f), where f is the strain at failure. The change of the creep curve is, for the case of plastic deformation, accounted for by replacing (20) with the expression: [ ]{ }p A C t t S B t tk l a m a= − − − + −σ σ ε σ1 exp ( ( ) ( ) ( )pl (21) where ta is the time to the last plastic deformation. Here, we have in mind, not creep that is accompanied by a continuous change of instantaneous plastic strain, but the effect of discrete instantaneous plastic strain at the moment of the application of the plastic deformation pl. When the plastic strain pl is accumulated under stress in creep tests, the effect of pl is automatically accounted for with the coefficients A,B,C,k,l,m. The values of the function S, describing the effect of plastic strain on creep, depend on the sign of pl. The form of the function S(pl) describing the effect of plastic strain on creep rate, dependent on the sign of pl, is determined from a series of tests with different values of pl and is either specified exactly or given by the appro- ximating function: S h q( )ε εpl pl= +1 (22) It has been established that the values of the material parameters h and q are slightly dependent on the stress level. The dependence of p on the sign of the stresses is taken into account as follows:  ( )p f= σ σsign (23) The effect of cyclic loading on the creep rate is con- sidered with inclusion of the function f1() in Eq. (19) as a multiplier: [ ]{ }p A C t t S f B t t k l a m a = − − − ⋅ ⋅ + − σ σ ε λ σ 1 1 exp ( ( ) ( ) ( ) ( )pl (24) where ta is the time to the last plastic deformation or to the last change in the stress sign. The function f1() is determined from tests in cyclic loading with constant cycle parameters: under these conditions, f1() = f2(k), where k is the number of the cycle. For k = 1, f2(k) = f1() = 1. Thus, with the chosen test temperatures, Ti (i = 1 ... N), we have the following expressions for p and p: expp A C p C p i i k l i i i = ⎛ ⎝ ⎜ ⎜ ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ − ⎛ ⎝ ⎜ ⎜ + σ 1− ε σ 1− εfi fi ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ − ⎡ ⎣ ⎢ ⎢ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ ⎥ ⎥ l a i t t S( ) ( , ( , ))ε ε σpl plsign f B p m i 1 1( )λ σ 1− ε + ⎛ ⎝ ⎜ ⎜ ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ ⎧ ⎨ ⎪ ⎪ ⎩ ⎪ ⎪ ⎪ ⎫ ⎬ ⎪ ⎪ ⎭ ⎪ ⎪ fi sign σ (25) p A p C p i k i i = ⎛ ⎝ ⎜ ⎜ ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ − − ⎛ ⎝ ⎜ ⎜ ⎜ ⎜ ⎜ ⎞ σ 1− ε σ 1− εfi fi 1 exp ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ − ⎡ ⎣ ⎢ ⎢ ⎢ ⎢ ⎢ ⎤ ⎦ ⎥ ⎥ ⎥ ⎥ ⎥ ⎧ ⎨ ⎪ ⎪ ⎩ ⎪ ⎪ ⎪ ⎫ ⎬ ⎪ ⎪ ⎭ ⎪ ⎪ l a i t t S( ) ( ,ε pl sign pl fi ( , )) ( )ε σ λ σ 1− ε f B p m i 1 1+ ⎛ ⎝ ⎜ ⎜ ⎜ ⎜ ⎜ ⎞ ⎠ ⎟ ⎟ ⎟ ⎟ ⎟ ⎧ ⎨ ⎪ ⎪⎪ ⎩ ⎪ ⎪ ⎪ ⎫ ⎬ ⎪ ⎪⎪ ⎭ ⎪ ⎪ ⎪ sign σ (26) To find p with a known value of p at each step of the integration in accordance with (17), it is necessary to exclude the parameter t – ta from these equations. The value of p with an arbitrary temperature T is determined by interpolating the values of p, found using the above-described method with several of the nearest values of Ti. As an additional material parameter we consider the limiting temperature T0: at T T≤ 0 , p = 0. If the calculations are performed with a change of sign for the stress  at any number of times (i.e., when the continuous function  passes through zero), the value of p changes as follows. Instead of (17) we have p = f(,p – a,T,pl,1) (17a) where a is the creep accumulated up to the moment of the change in the sign of the stresses. We should additionally assume that for a complex stress state the cyclic loading and the preliminary plastic deformation affects only the scalar properties of the materials. Then, instead of (17) we can write 9  ( , , ..., , , , , ,( ) ( ) ( ) ( ) ( ) ( ) ( )p F p p p pij n n n= −0 1 1 0 1ε ε ..  ., ε ε σ λ σ ( ) ( ), , , , )n n ij T S −1 (27) Here, Sij are the components of the stress deviator; σ is the stress intensity; p n( ) and ε ( )n are the intensities of the creep strain and (non-creep) plastic strain: p p pn ij ij t t n n ( ) /  = ⎛ ⎝⎜ ⎞ ⎠⎟ − ∫ 2 3 1 1 2 dτ; ε ε ε τ( ) /  n ij ij t t n n = ⎛ ⎝⎜ ⎞ ⎠⎟ − ∫ 2 3 1 1 2 d (28) where n is the number of half-cycles. L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 261 4 METHOD OF DETERMINING THE MATERIAL’S CHARACTERISTICS The system of equations (18–20) contains several material constants A B C k l m t h q A B C k l m t h q p p 1 1 1 1 1 1 1 1 2 2 2 2 2 2 2 2 ε ε 1 2 ( ) ( ) . . . . . . . . . ( )A B C k l m t h qN N N N N N pN N Nε T0 and the functions S[pl, sign(pl,)], f1(). The constants Ai,Bi,Ci,ki,li,mi are determined from an analysis of the creep data for uniaxial stress in the first and second stages at N temperatures Ti and several (at least six) stresses. We select three numbers, a, b and c, determining the first and second creep stages for each experimental creep curve (Figure 4). The data on a, b and c for different values of  and T = const are used to find the values of A, B, C, k, l, m, applying the method of least squares and the basis of the following power relations: a A b B c Ck k l= = =σ σ σ; ; (29) The parameters hi and qi of the function S(pl) are determined with the analysis of the results of creep tests at n temperatures Ti, one stress for each value of Ti, and several values of preliminary plastic strain pl within the range from –3 % to –5 % to +3 % to + 5 % (at least three values pl < 0, three values pl > 0, and pl = 0). For materials with the function S(pl) depending on the stress level, it is better if the results of the determi- nation of S(pl) are obtained using exact specifications for the values of S(pl,). The function f1() is found with the analysis of results of the creep tests with alternating loading in the program shown in Figure 5 and the determination of the depen- dence on the cycle (the creep strain accumulated within a half-cycle) Figure 6. 5 EXAMPLES OF THE CALCULATION The values of the coefficients in Eqs. (25) and (26) for the alloy KhN70VMYuT (EI765) are shown in Table 1. The creep strain that occurs with an arbitrary law of change of stress, temperature, and instantaneous plastic strain is determined by the numerical integration of Eqs. (17) and (17a) using the fourth-order Runge-Kutt method and a computer algorithm. The algorithm can be obtained without a computer by using a simple law of L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... 262 Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 Figure 6: Method for determining the function f1 () from cyclic creep curves. Slika 6: Metoda za dolo~itev funkcije f1 () iz krivulje cikli~nega lezenja Figure 4: Creep curve Slika 4: Krivulja lezenja Figure 5: Diagram of loading of the specimens Slika 5: Diagram obremenitve preizku{ancev the change in stress (t), temperature T(t), and plastic strain pl(t). As examples of the calculation, the following va- riants were analysed 6,10: creep during the first and third stages ( = const, T = const); for a temperature change ( = const); for an alternating stress (T = const); for conditions of cyclic plastic deformation ( = const, T = const); for cyclic creep with alternating plastic strain; with plastic strain and a changing stress (T = const); for alternating stress during a changing temperature. Also, the stress relaxation with 0 < y; with 0 > y and addi- tional preliminary plastic deformation; with additional loadings and cyclic plastic strain were examined. The analysis of the agreement for theoretical and experimental data for twenty different grades of steels and alloys showed that the proposed creep model and the method of determining its parameters were valid (see Figure 7, for example). The studies 11,12 proposed variants of the above method for calculating the creep to determine the stress-strain state of blades (uniaxial stress state) and disks (complex stress state) with multiple starts of gas-turbine engines. 6 CONCLUSION The variations in aniso-thermal theories of plastic flow and of the theory of work hardening in creep with structural parameters have been considered, making it possible to include the mutual effect of two forms of deformations within the framework of the traditional approach. 7 REFERENCES 1 Handbook of Materials Behavior Models. Editor Jean Lemaitre. Academic Press, 2001, 900 p 2 D. A. Gokhfeld, O. S. Sadakov, The plasticity and creep of elements constructions by repeated loadings (in Russian), Moscow, 1984, 256 p 3 Yu. N. Rabotnov, Creep problems in structural members, North- Holland, 1969, 750 p 4 I. A. Birger, B. F. Shorr et al., The high temperature strength of machine parts (in Russian), Mashinostroenie, Moscow, 1975, 454 p 5 G. S. Pisarenko, A. A. Lebedev, Deformation and strength of materials in the complex stressed condition (in Russian), Naukova Dumka, Kiev, 1976, 415 p 6 L. B. Getsov, Materials and strength of gas turbine parts (in Russian), M: Nauka, 1996, 591 p 7 M. G. Kabelevskii, Mechanika. Tverdogo. Tela, (1972) 1, 169–174 8 L. B. Getsov, Strength of materials, (1978) 2, 22 9 L. B. Getsov, M. G. Kabelevskii, Strength of materials, (1978) 6, 44 10 A. M. Borzdyka, L. B. Getsov, Stress relaxation in metals and alloys (in Russian), Metallurgia, Moscow, 1978, 255 p 11 L. B. Getsov, V. K. Dondoshanskii, Sudovye Energ. Ustanovki, (1975) 7, 56 12 M. G. Kabelevskii, L. B. Getsov, Mashinovedenie, (1977) 4, 82 L. B. GETSOV, M. G. KABELEVSKIY: A NEW VERSION OF THE THEORY OF DUCTILITY AND CREEP ... Materiali in tehnologije / Materials and technology 41 (2007) 6, 257–263 263 Table 1: Parameters of the creep resistance of the alloy EI765 Tabela 1: Parametri odpornosti proti lezenju za zlitino E1765 Temp., T/°C n/% p/% h1 l1 f2(15) y/(kg/mm2) Stress level –lg A –lg B –lg C k0 l0 m 650 8 17 – – – –  < y 22,3 37,9 12,96 10,3 7,63 18,75 700 8-14 14 26 0,5 – 64,5  < y 13,62 16,3 9 5,56 5,92 7,41  > y 27,7 39,2 9 13,33 5,92 20 750 13-15 16 350 1,3 4 55,6  < y 5,8 16,62 2,8 1,14 2,68 8,33  > y 18,4 28,9 2,8 8,33 2,68 15,38 Figure 7: Experimental (full lines) and theoretical (dashed lines) curves describing the effect of cyclic plastic deformation on the creep resistance (a) and the relaxation (b) for the alloy EI765 at 700 °C Slika 7: Eksperimentalne (cele ~rte) in teoreti~ne (~rtkane ~rte) krivulje, ki opisujejo vpliv cikli~ne plasti~ne deformacije na odpornost proti lezenju (a) in relaksacijo (b) pri 700 °C za zlitino EI765 I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL TERMOELEKTRI^NE LASTNOSTI MONOKRISTALNEGA KVAZIKRISTALA Al64Cu23Fe13 Igor Smiljani}1, Ante Bilu{i}1,2, @eljko Bihar1, Jagoda Lukatela1, Boran Leonti}1, Janez Dolin{ek3, Ana Smontara1 1Institute of Physics, Bijeni~ka 46, HR-10000 Zagreb, Croatia 2Faculty of Natural Sciences, University of Split, N. Tesle 12, HR-21000 Split, Croatia 3J. Stefan Institute, Jamova 39, SI-1000 Ljubljana, Slovenia ismiljanicifs.hr Prejem rokopisa – received: 2007-07-12; sprejem za objavo – accepted for publication: 2007-08-01 We performed investigations of the electrical resistivity, thermopower and thermal conductivity of a monocrystalline i-Al64Cu23Fe13 as well as a polycrystalline i-Al63Cu25Fe12 icosahedral quasicrystal, for comparison. The electrical resistivity of both samples, the monocrystalline i-Al64Cu23Fe13 and the polycrystalline i-Al63Cu25Fe12, exhibits a negative temperature coefficient with 4K = 3950 µ cm and 4K= 4900 µ cm, and the ratio 4K/300K = 1.8, 4K/300K = 1.7, respectively. The thermopowers are large and have a negative sign. In addition, the thermopower of the monocrystalline i-Al64Cu23Fe13 exhibits a sign reversal at T = 278 K. The thermal conductivity is anomalously low, of the order of 1 W/mK at room temperature, with a slightly different temperature variation at low temperatures. On the basis of these results, we concluded that there are no systematic differences between the high-quality monocrystalline and polycrystalline icosahedral i-Al-Cu-Fe quasicrystals. Moreover, the reported transport properties of i-Al-Cu-Fe appear to be intrinsic to this family of icosahedral quasicrystals. Keywords: quasicrystals; i-AlCuFe, physical properties, resistivity, thermal conductivity Raziskali smo elektri~no prevodnost, termonapetost in toplotno prevodnost monokristalnega i-Al64Cu23Fe13 in za primerjavo tudi polikristalnega ikozaedri~nega kvazikristala i-Al63Cu25Fe12. Elektri~na upornost obeh vzorcev ima negativen temperaturni koeficient z 4K = 3950 µ cm in 4K = 4900 µ cm, ter razmerje 4K/300K = 1.8, 4K/300K = 1.7. Termonapetosti so velike in z negativnim predznakom, termonapetost monokristalnega i-Al64Cu23Fe13 pa ima spremembo predznaka pri T = 278 K. Toplotna prevodnost je anormalno majhna, je reda velikosti 1 W/mK pri sobni temperaturi in z nekoliko druga~no temperaturno odvisnostjo pri nizki temperaturi. Na podlagi rezultatov meritev sklepamo, da ni sistemati~ne razlike med visokokakovostnima monokristalnima in mnogokristalnima ikozaedri~nima kvazikristaloma i-Al-Cu-Fe. Poleg tega so transportne lastnosti i-Al-Cu-Fe zna~ilne za to dru`ino ikozaedri~nih kvazikristalov. Klju~ne besede: kvazikristali, i-AlCuFe, fizikalne lastnosti, upornost, toplotna prevodnost 1 INTRODUCTION The family of icosahedral i-Al-Cu-Fe quasicrystals is currently one of the most studied, due to its excellent thermal stability. Most studies reported so far were performed on polycrystalline samples, and include inve- stigations of the electrical resistivity and magnetoresi- stance,1-9 thermoelectric power,9-12 thermal conducti- vity,2,9,13 magnetism,3,8,14 and Hall coefficient.1,5,7,8 Though polycrystalline samples may acquire quite a high structural perfection through a proper thermal annealing procedure, rapid quenching to room temperature after annealing inevitably results in a strained material that also contains high thermal vacancy concentration for the room temperature conditions (i.e., the quenched-in vacancy concentration is in equilibrium for the much higher temperature of annealing). In addition, grain boundaries may hinder the propagation of electrons and phonons, thus affecting long-range electrical and heat-transport phenomena. In order to test for the true intrinsic properties of i-Al-Cu-Fe quasicrystals, it is desirable to compare the physical properties of the polycrystalline material with those measured on high-quality monocrystalline samples, where structural imperfections are largely absent. Therefore, we have performed a study by investigating the electrical resistivity, the thermoelectric power and the thermal conductivity of a monocrystalline i-Al64Cu23Fe13 and a polycrystalline i-Al63Cu25Fe12 quasicrystal. 2 EXPERIMENTAL PROCEDURE We investigated two samples with slightly different compositions, a monocrystalline i-Al64Cu23Fe13 (in the following text abbreviated as i-Al64Cu23Fe13) and a polycrystalline i-Al63Cu25Fe12 icosahedral quasicrystal (in the following text abbreviated as i-Al63Cu25Fe12). The i-Al63Cu25Fe12 were made from large polycrystalline ingots prepared by conventional casting and subsequent annealing, and it was verified with X-ray diffraction that the samples are single-phase icosahedral. A large mono- crystalline i-Al64Cu23Fe13 quasicrystal was prepared by the Czochralski technique and annealing removed the strains. It has an almost phason-free quasicrystalline structure and shows superior quasicrystallinity on both the macro- and microscopic scales. The samples were Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 265 UDK 669'71'3'1:620.17 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)265(2007) shaped in the form of a prism, with dimensions 3.9 mm × 1.5 mm × 1.4 mm (i-Al64Cu23Fe13) and 7.2 mm × 1.6 mm × 1.2 mm (i-Al63Cu25Fe12). The electrical resistivity was measured by a standard four-probe technique with applied currents of 0.1 mA to 1 mA, while the thermo- electric power was measured with respect to high-purity gold lead wires, using a deferential technique. The thermal conductivity was measured using an absolute steady-state heat-flow method. The thermal flux was generated by a 1 k RuO2 chip-resistor glued to one end of the sample, while the other end was attached to a copper heat sink. The temperature gradient across the sample was monitored by a chromel-constantan differen- tial thermocouple. 3 RESULTS AND ANALYSIS 3.1 Electrical resistivity and thermopower The electrical resistivity ((T)) and thermopower (S(T)) of i-Al64Cu23Fe13 and i-Al63Cu25Fe12 were measured in the temperature range from 4 K to 300 K. The results are shown in Figure 1 and Figure 2. The resistivities of i-Al64Cu23Fe13 and i-Al63Cu25Fe12 exhibit a negative temperature coefficient, the room temperature values are 300K = 2200 µ cm and 300K = 2900 µ cm respectively and the total increases of resistivity is by factors of R = 4K/300K = 1.8 and R = 4K/300K = 1.7, respectively. In addition, (T) of the i-Al64Cu23Fe13 exhibits a weakly pronounced maximum with the peak value 300K= 4040 µcm, at 20 K. The thermopowers S(T) are, in general, large and, in addition, exhibit an interesting feature of a sign reversal (Figure 2) in the case of i-Al64Cu23Fe13. Below 120 K, S(T) is negative with a negative slope, whereas around 120 K it exhibits a minimum and the slope is reversed. Consequently, the S(T) of i-Al64Cu23Fe13 changes sign to positive at T = 278 K. For the analysis of (T) and S(T) we used the spectral resistivity model of Landauro and Solbrig,11,12,16 where both quantities are analyzed simultaneously by presu- ming a specific structure- and composition-related form of the energy-dependent spectral resistivity function () (or its inverse, the spectral conductivity ()=1/()). Using the Kubo-Greenwood formalism, the tempera- ture-dependent electrical conductivity is calculated according to σ εσ ε ε ε ( ) ( ) ( , ) T f T= −⎛ ⎝⎜ ⎞ ⎠⎟∫ d ∂ ∂ (1) whereas the thermopower is obtained from S T k e T T k T f T ( ) ( ) ( ) ( ) ( , )= − − ⎛ ⎝ ⎜ ⎞ ⎠ ⎟ −⎛⎝⎜ ⎞ ⎠∫ B B d σ εσ ε ε µ ε ε ∂ ∂ ⎟ (2) Here, f(,T) = {exp[( – µ)/kBT] + 1}–1 is the Fermi- Dirac function and µ(T) is the chemical potential, which is written in the low-temperature representation as17 µ ε ε ε ε ( ) ( ) ln ( ) T k T n= − ⎛ ⎝⎜ ⎞ ⎠⎟F B d d F 2 2 6 π = F – T 2 (3) The electronic density of states n() is related to the spectral conductivity via the Einstein relation ()= (e2/V)n()D() with D() being the electronic spectral diffusivity. The only material-dependent quantity in Eqs. (1–3) is (), so that a proper model of the spectral conductivity should reproduce both (T) and S(T) at the same time. The ab-initio-derived spectral resistivity could be modeled by the superposition of two Lorentzians I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL 266 Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 0 100 200 300 -30 -20 -10 0 i-Al Cu Fe (polycrystalline) fit i-Al Cu Fe (single-crystalline) fit S /( µ V /K ) T/K Figure 2: Thermopower of the monocrystalline i-Al64Cu23Fe13 and the polycrystalline i-Al63Cu25Fe12, respectively. The fits (solid lines) of (T) and S(T) were made simultaneously with the Kubo-Green- wood formalism using the spectral resistivity function (). Slika 2: Termonapetost monokristalnega i-Al64Cu23Fe13 in polikristal- nega i-Al63Cu25Fe12. Pribli`ek (cele ~rte) za (T) in S(T) je pripravljen isto~asno s Kubo-Greenwood formalizmom in z uporabo funkcije spektralne upornosti () 0 100 ρ Ω /( µ c m ) 2 T/K 00 300 2000 3000 4000 5000 i- Al Cu Fe (polycrystalline) fit i-Al Cu Fe (single-crystalline) fit Figure 1: Electrical resistivity of the polycrystalline i-Al63Cu25Fe12 and monocrystalline i-Al64Cu23Fe13, respectively. The fits (solid lines) of (T) and S(T) were made simultaneously with the Kubo- Greenwood formalism using the spectral resistivity function (). Slika 1: Elektri~na upornost polikristalnega i-Al63Cu25Fe12 in monokristalnega i-Al64Cu23Fe13. Pribli`ek (cele ~rte) (T) in S(T) je napravljen isto~asno s Kubo-Greenwood formalizmom z uporabo funkcije spektralne upornosti () ρ ε γ (ε−δ ) γ α γ (ε−δ ) γ1 2 1 2 2 2 2 2 ( ) = + ⎡ ⎣ ⎢ ⎤ ⎦ ⎥ + + ⎡ ⎣ ⎢ ⎤ ⎦ ⎥A 1 11 2 π π ⎧ ⎨ ⎩⎪ ⎫ ⎬ ⎭ (4) where 1/i is the height of a Lorentzian, 2i its FWHM, i its position with respect to the Fermi energy F (taken to be at the origin of the energy scale; F = 0) and  is the relative weight of the Lorentzians. The position of the narrow resistivity peak with respect to the Fermi energy F is responsible for the anomalous electronic transport properties. As this peak is due to a specific distribution of Fe atoms in the structure, quasiperiodicity alone cannot account for the anomalous transport properties of i-Al-Cu-Fe quasicrystals; a right chemical decoration is also needed. The Fermi energy can be shifted on the scale of a few 100 meV by deviations in the stoichiometry and/or by defects in both structure and chemical decoration,18,19 so that the relative position of the narrow peak can change on this energy scale in samples of slightly different compo- sition and annealing treatment. Consequently, solely on the basis of small shifts of F, the thermopower of i-Al-Cu-Fe samples of similar composition can switch between large positive and large negative values and it may also change sign with temperature, as demonstrated for the i-Al62Cu25.5Fe12.5 polycrystalline sample.9,10,12 The fits of the experimental (T) and S(T) data were performed simultaneously with Eqs. (1-4) by adjusting the set of parameters (A, , 1, 2, 1 and 2) pertinent to the shape of the spectral resistivity (). The starting value of the parameter  entering the temperature- dependent chemical potential of Eq. (3) was determined by recognizing that in the case when the spectral variation of the electronic diffusivity can be neglected, one can replace n() by () in Eq. (3). The initial value was obtained by using the Mott formula S T k e TMott B d d F ( ) ln ( )= ⎛ ⎝⎜ ⎞ ⎠⎟ π 2 2 3 σ ε ε ε so that  = –0.5e(SMott(T)/T). The fits are shown as solid lines in Figure 1 and Figure 2; the fit parameters are collected in Table 1. The fits of both (T) and S(T) are excellent in the whole investigated temperature range. The magnitude of the electrical resistivity of the monocrystalline i-Al64Cu23Fe13 is in-line with the values reported for the polycrystalline i-Al-Cu-Fe. In poly- crystalline i-Al-Cu-Fe the total variation of (T) over the relevant Fe range is merely a factor of two,5 so that the material is best classified as a semi-metal over the whole icosahedral concentration range with no indication of a large resonant increase of the resistivity. The 4K resistivity value of our i-Al64Cu23Fe13 matches well with that of the polycrystalline samples from the study5 with the same Fe concentration. However, since the Fermi energy of our i-Al64Cu23Fe13 is located nearly at the maximum of the spectral resistivity20, further shifts of F over the resistivity peak due to a small variation of the Fe composition would not result in an additional increase of the resistivity, but can only make it smaller. This hints that the factor-of-two larger peak resistivity of the polycrystalline i-Al-Cu-Fe material at the Fe 12.5 % concentration,5 as compared to the monocrystalline i-Al64Cu23Fe13, could originate in extrinsic factors like grain boundaries and other lattice imperfections that act as additional scattering centers for the conduction electrons. 3.2 Thermal conductivity The measured thermal conductivities (T) of i-Al64Cu23Fe13 and i-Al63Cu25Fe12 are displayed in Figure 3. The conductivity value at room temperature of i-Al64Cu23Fe13 amounts 300K = 1.7W/mK, whereas in the case of the i-Al63Cu25Fe12 is 300K = 2.9 W/mK. These values are surprisingly low for an alloy of regular metals and are comparable to the thermal conductivities of known thermal insulators, amorphous fused silica21 and the technologically widespread thermally insulating material, yttrium-doped zirconia ceramics22. The (T) data were analyzed with a semi-quantitative model, appropriate for icosahedral quasicrystals and their approximants22-29. The thermal conductivity parameter (T) is divided into three terms I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 267 Table 1: Parameters of the spectral resistivity () of Eq. (4), obtained from the simultaneous fits of (T) and S(T) for i-Al64Cu23Fe13 (monocrystalline) and i-Al63Cu25Fe12 (polycrystalline) Tabela 1: Parameter posebne upornosti () ena~ba (4), dolo~en z isto~asnim pribli`kom (T) in S(T) za i-Al64Cu23Fe13 (monokristalen) in i-Al63Cu25Fe12 (polikristalen) sample Al-Cu-Fe A/(µ cm eV) d1/(meV) 1/(meV)  2/(meV) 2/(meV) mono-crystalline 392 –43 241 1.13 –9 38 polycrystalline 847 –5.2 587 1.07 –16 55 Table 2: Fit parameters of the thermal conductivity (T) i-Al64Cu23Fe13 (monocrystalline) and i-Al63Cu25Fe12 (polycrystalline), respectively Tabela 2: Parametri pribli`ka toplotne prevodnosti (T) za i-Al64Cu23Fe13 (monokristalen) in i-Al63Cu25Fe12 (polikristalen) sample Al-Cu-Fe Leff κH 0 (W/mK) Ea/(meV) A/(s–1K–2) B/(s–1K–4) ß mono-crystalline 2.1 0.7 6.3 1.2x107 2.8x104 3.2 polycrystalline 2.5 2.2 16.2 3.5 x 10 6 7.2 x 103 2.0 κ κ κ κ( ) ( ) ( ) ( )T T T T= + +el D H (5) The electronic contribution e is obtained using the empirical Wiedemann-Franz law ( el=LoT) with a tem- perature-dependent effective Lorenz number L T T T T ( ) ( ) ( ) = × κ σ el (6) The lattice contribution ( – el) is analyzed by considering (i) the propagation of long-wavelength acoustic phonons (for which the quasicrystal structure is an elastic continuum) within the Debye model and (ii) hopping of localized vibrations within the icosahedral cluster substructure, which participate in the heat transfer via thermally activated hopping. In the simplest model, the hopping of localized vibrations is described by a single activation energy Ea, yielding a contribution to the thermal conductivity κ κH H 0 a B = −⎛ ⎝ ⎜ ⎞ ⎠ ⎟exp E k T (7) where κH 0 is a constant. The Debye thermal conductivity is written as30 κ τ( ) θ D D D d= −∫C T x x e e x T x x 3 0 4 21 / ( ) (8) where C k /D B= 4 2 32π ν , ν is the average sound velo- city, D the Debye temperature,  the phonon relaxation time and x /k T= ω B , where ω is the phonon energy. The different phonon-scattering processes are incorporated into the relaxation time (x) and we assume that Matthiessen’s rule is valid, τ τ−1− = ∑1 j , where τ−1j is the scattering rate related to the j-th scattering channel. In analogy with the Al-Pd-Mn approximant and quasicrystal phases,24,25,30 we consider two dominant scattering processes in the investigated temperature range: (1) the scattering of phonons on structural defects of stacking-fault type with the scattering rate τ sf − =1 2 2Ax T and (2) umklapp processes with the phenomenological form of the scattering rate pertinent to quasicrystals,23,24,29 τ um ß ß− −=1 4Bx T , so that τ τ τ− − −= +1 1 1sf um . The Debye temperature of i-Al-Cu-Fe was estimated from the specific heat32 as D 560 K and the Debye constant CD was determined from ultrasonic data30. The fits (solid lines in Figure 4) are excellent and the fit parameters are collected in Table 2. The electronic (el), Debye (D) and hopping (H) con- tributions are shown separately on the graph. The temperature-dependent effective Lorenz number L(T) of i-Al64Cu23Fe13 and i-Al63Cu25Fe12 deviates considerably from the Wiedemann-Franz value L0, amounting L/L0= 2.1 and L/L0= 2.5 at 300 K, respectively, and the elec- trons carry around 40 % of the total heat in both cases. The Debye contribution exhibits a maximum at about 30 K and declines above this temperature, whereas the hopping contribution becomes significant at elevated temperatures. The activation energy for the hopping of i-Al64Cu23Fe13 and i-Al63Cu25Fe12 was determined as Ea 16 meV and Ea 16 meV, respectively. These energies correlate with the inelastic neutron and X-ray scattering experiments on i-Al-Pd-Mn quasicrystals, where dispersionless vibrational states were identified for energies higher than 12 meV. Such dispersionless states indicate localized vibrations and are considered to be a consequence of a dense distribution of energy gaps in the phonon excitation spectrum of quasicrystals. The parameters B and define phonon scattering by umklapp processes in a phenomenological way. The fit-determined = 3.2 value for i-Al64Cu23Fe13 yields I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL 268 Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 0 100 200 300 0 1 2 3 i-Al Cu Fe (polycrystalline) i-Al Cu Fe (single-crystalline) K /( W /m K ) T/K Figure 3: Thermal conductivity (T) of the monocrystalline i-Al64Cu23Fe13 and the polycrystalline i-Al63Cu25Fe12 , respectively Slika 3: Toplotna prevodnost (T) za monokristalni i-Al64Cu23Fe13 in polikristalni i-Al63Cu25Fe12 10 100 0 1 2 3 κ(i-Al Cu Fe , single-crystalline) κ κ κ fit κ(i-Al Cu Fe , polycrystalline) κ κ κ fit κ/ (W /m K ) T/K Figure 4: Thermal conductivity (T) with the fits to the total (T), of the monocrystalline i-Al64Cu23Fe13 and the polycrystalline i-Al63Cu25Fe12. The three contributions to the total (T), electronic el, Debye D and hopping H, are shown separately Slika 4: Toplotna prevodnost (T) s pribli`kom za skupen (T), za monokristalni i-Al64Cu23Fe13 in za polikristalni i-Al63Cu25Fe12. Posebej so prikazani trije prispevki k skupni (T), elektronski el, Debye D in presko~na H the frequency- and temperature dependence of the umklapp term τ ω−1um ∝ 3 2 0 8. .T and indicates similarity with the modified quasi-umklapp scattering rate τ ω−1um ∝ 3 T , used for the analysis of the thermal conductivity of i-Zn-Mg-Y quasicrystals, while for the polycrystalline sample, the fit-determined = 2.0 τ ω−1um ∝ 2 2T indicates the quasi-umklapp scattering rate obtained for the i-Al-Pd-Mn quasicrystals25,29. Here it should be mentioned that the Debye and hopping contributions slightly compensate for each other in the fit procedure, so that the parameter values characte- rizing D and H should be considered at the qualitative level. Regarding the comparison of the thermal conductivity of monocrystalline and polycrystalline i-Al-Cu-Fe, we are not aware of other quantitative analyses of (T) of polycrystalline samples in the sense of Eqs. (5–8), so that the comparison has to be made at the level of experimental thermal conductivities. The room-temperature value for i-Al64Cu23Fe13 amounts to 300K = 1.7 W/mK, whereas for the i-Al63Cu25Fe12 it is 300K = 2.9 W/mK. Though the scatter of the known values of thermal conductivity of i-Al-Cu-Fe34 is relatively large, there seems to be no systematic diffe- rence between the polycrystalline and monocrystalline samples. 4 CONCLUSION We performed investigations of electrical resistivity, thermoelectric power and thermal conductivity on a monocrystalline i-Al64Cu23Fe13 and a polycrystalline i-Al63Cu25Fe12 icosahedral quasicrystal, for comparison. The electrical resistivity and thermopower analysis shows that the Fermi energy is located at the minimum of the pseudogap in the spectral conductivity (). All this gives evidence that we are dealing with icosahedral quasicrystal samples of exceptional quality, so that its physical properties may be considered as intrinsic to the i-Al-Cu-Fe phase. A comparison of the investigated monocrystalline i-Al64Cu23Fe13 to the polycrystalline i-Al63Cu25Fe12, however, shows that there are no pronounced differences between the two forms of the material. While there are essentially no differences in the magnetic properties of the monocrystalline and polycrystalline materials, the electrical resistivity of the polycrystalline material is larger by a factor of two. This difference can be easily accounted for by the grain boundaries and other lattice imperfections that act as additional scattering centers for the conduction electrons. Comparing the thermopowers of the monocrystalline and polycrystalline i-Al-Cu-Fe materials, in general, the S(T) magnitude and temperature dependence depend strongly on the position of F relative to the spectral resistivity peak, so that slight differences in the samples’ stoichio- metry and structural perfection may lead to very different thermopowers in both magnitude and sign. A quantitative comparison of the thermal conductivities of monocrystalline and polycrystalline i-Al-Cu-Fe is less straightforward due to the random scatter of the reported values, but 300K of the monocrystalline i-Al64Cu23Fe13 fits within the range of values for the polycrystalline i-Al63Cu25Fe12. To conclude, we found no systematic differences in the electrical resistivity, thermoelectric power and thermal conductivity between the high-quality monocrystalline and polycrystalline i-Al-Cu-Fe quasi- crystals, and the reported physical properties of the other i-Al-Cu-Fe quasicrystals appear to be intrinsic to this family of icosahedral quasicrystals. Acknowledgements We would like to thank Y. Yokoyama and Y. Calvayrac for giving us the samples of monocrystalline i-Al64Cu23Fe13 and polycrystalline i-Al63Cu25Fe12 icosahedral quasicrystal, respectively. This work was done within the activities of the 6th Framework EU Network of Excellence "Complex Metallic Alloys" (Contract No. NMP3-CT-2005-500140), and has been supported in part by the Ministry of Science, Education and Sports of Republic of Croatia through the Research Projects Nos. 035-0352826-2848 and 177-0352828- 0478. 5 REFERENCES 1 T. Klein, A. Gozlan, C. Berger, F. Cyrot-Lackmann, Y. Calvayrac, A. Quivy, Europhys. Lett. 13 (1990), 129–134 2 A. Smontara, J.C. Lasjaunias, C. Paulsen, A. Bilu{i}, Y. Calvayrac, Mat. Sci. Eng. 294-296 (2000), 706–710 3 T. Klein, C. Berger, D. Mayou, F. Cyrot-Lackmann, Phys. Rev. Lett. 66 (1991), 2907–2910 4 T. Klein, H. Rakoto, C. Berger, G. Fourcaudot, F. Cyrot-Lackmann, Phys. Rev. B 45 (1992), 2046–2049 5 P. Lindqvist, C. Berger, T. Klein, P. Lanco, F. Cyrot-Lackmann, Y. Calvayrac, Phys. Rev. B 48 (1993), 630–633 6 D. Mayou, C. Berger, F. Cyrot-Lackmann, T. Klein, P. Lanco, Phys. Rev. Lett. 70 (1993), 3915–3918 7 M. Ahlgren, P. Lindqvist, M. Rodmar, Ö. Rapp, Phys. Rev. B 55 (1997), 14847–14854 8 R. Escudero, J. C. Lasjaunias, Y. Calvayrac, M. Boudard, J. Phys.: Condens. Matter 11 (1999), 383–404 9 A. Bilu{i}, A. Smontara, J.C. Lasjaunias, J. Ivkov, Y. Calvayrac, Mat. Sci. Eng. 294-296 (2000), 711–714 10 A. Bilu{i}, I. Be{li}, J. Ivkov, J.C. Lasjaunias, A. Smontara, Fizika A (Zagreb) 8 (2000), 183–194 11 C.V. Landauro, H. Solbrig, Mat. Sci. Eng. A 294-296 (2000), 600–603 12 H. Solbrig, C.V. Landauro, Quasicrystals, Structure and Physical Properties, Wiley-VCH, Weinheim 2003, 254 13 A. Perrot, J.M. Dubois, M. Cassart, J.P. Issi, Proc. of the 5th Inter. Conf. on Quasicrystals, World Scientific, Singapore, 1995, 588–601 14 K. Fukamichi, Physical Properties of Quasicrystals, Springer, New York 1999, p. 295 and references therein 15 Y. Yokoyama, Y. Matsuo, K. Yamamoto, K. Hiraga, Mater. Trans., JIM, 43 (2002), 762–765 16 C. V. Landauro, H. Solbrig, Physica B 301 (2001), 267–275 I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 269 17 N. W. Aschroft, N. D. Mermin, Solid State Physics, Saunders College Publishing, London 1976, 46 18 F. S. Pierce, P. A. Bancel, B. D. Biggs, Q. Guo, S. J. Poon, Phys. Rev. B 47 (1993), 5670–5676 19 H. Solbrig, C. V. Landauro, A. Löser, Mat. Sci. Eng. A 294–296 (2000), 596–599 20 J. Dolin{ek, S. Vrtnik, M. Klanj{ek, Z. Jagli~i}, A. Smontara, I. Smiljani}, A. Bilu{i}, Y. Yokoyama, A. Inoue, C.V. Landauro, Phys. Rev. B (in press) 21 D.-M. Zhu, Phys. Rev. B 50 (1994), 6053–6056 22 R. Mévrel, J.-C. Laizet, A. Azzopardi, B. Leclercq, M. Poulain, O. Lavigne, D. Demange, J. Eur. Cer. Soc. 24 (2004), 3081–3089 23 @. Bihar, A. Bilu{i}, J. Lukatela, A. Smontara, P. Jegli~, P. J. McGuiness, J. Dolin{ek, Z. Jagli~i}, J. Janovec, V. Demange, J. M. Dubois, J. Alloys Compd. 407 (2006), 65–73 24 J. Dolin{ek, P. Jegli~, P. J. McGuiness, Z. Jagli~i}, A. Bilu{i}, @. Bihar, A. Smontara, C.V. Landauro, M. Feuerbacher, B. Grushko, K. Urban, Phys. Rev. B 72 (2005), 064208–11 25 A. Bilu{i}, A. Smontara, J. Dolin{ek, P. J. McGuiness, H. R. Ott, J. Alloys Compd. 432 (2007), 1–6 26 A. Smontara, I. Smiljani}, A. Bilu{i}, Z. Jagli~i}, M. Klanj{ek, S. Roitsch, J. Dolin{ek, M. Feuerbacher, J. Alloys Compd. 430 (2007) 29–38 27 A. Smontara, I. Smiljani}, A. Bilu{i}, B. Grushko, S. Balanetskyy, Z. Jagli~i}, S. Vrtnik, J. Dolin{ek , J. Alloys Compd., in press 28 J. Dolinsek, T. Apih, P. Jeglic, I. Smiljanic, A. Bilusic, @ . Bihar, A. Smontara, Z. Jaglicic, M. Heggen, M. Feuerbacher, Intermetallics 15 (2007) 1367–1376 29 A. Bilu{i}, @. Budrovi}, A. Smontara, J. Dolin{ek, P. C. Canfield., I. R, Fisher, J. Alloys Compd. 342 (2002), 413–415 30 R. Berman, Thermal Conduction in solids, Clarendon Press, Oxford 1978, 23 31 P. A. Kalugin, M. A. Chernikov, A. Bianchi, H. R. Ott, Phys. Rev. B 53 (1996), 14145–14151 32 J. C. Lassjaunias, Y. Calvayrac, H. Yang, J. Physique 17 (1997), 959–976 33 Y. Amazit, M. de Boissieu, A. Zarembowitch, Europhys. Lett. 20 (1992), 703–706 34 A. Bilu{i}, D. Pavuna, A. Smontara, Vacuum 61 (2001), 345–348 I. SMILJANI] ET AL.: THERMOELECTRICAL PROPERTIES OF A MONOCRYSTALLINE Al64Cu23Fe13 QUASICRYSTAL 270 Materiali in tehnologije / Materials and technology 41 (2007) 6, 265–270 T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 PHASES IN A QUASICRYSTALLINE ALLOY Al64,4Cu23,5Fe13,1 Tonica Bon~ina1, Bo{tjan Markoli2, Ivan An`el1, Franc Zupani~1 1Univerza v Mariboru, Fakulteta za strojni{tvo, Smetanova 17, SI-2000 Maribor, Slovenija 2Univerza v Ljubljani, Naravoslovnotehni{ka fakulteta, A{ker~eva 12, SI-1000 Ljubljana, Slovenija tonica.boncinauni-mb.si Prejem rokopisa – received: 2007-09-24; sprejem za objavo – accepted for publication: 2007-10-18 V ternarnem sistemu Al-Cu-Fe se pojavlja t. i. i-faza (ikozaedri~ni kvazikristal), ki je termodinamsko ravnote`na faza in s tem sestavni del ravnote`nega faznega diagrama. Na enofazno podro~je i-faze meji veliko {tevilo intermetalnih faz, ki so lahko glede na sestavo zlitine, razmere pri strjevanju in toplotni obdelavi v stabilnem ali metastabilnem ravnote`ju z i-fazo. Sinteza enofazne kvazikristalne zlitine je mogo~a samo v ozkem koncentracijskem obmo~ju in pri primernem na~inu toplotne obdelave, zato je poznanje in ugotavljanje faz klju~nega pomena. V raziskavi smo izdelali zlitino Al64,4Cu22,5Fe13,1 in vzorce toplotno obdelali na razli~ne na~ine. Posamezne faze in druge mikrostrukturne zna~ilnosti smo ugotovili z ustrezno metalografsko pripravo vzorcev in metodami svetlobne mikroskopije (SM), vrsti~ne elektronske mikroskopije (SEM), mikrokemi~ne analize (EDS), rentgenske fazne analize (XRD), presevne elektronske mikroskopije (TEM) ter z merjenjem mikrotrdote. Faze smo lo~evali tudi glede na njihovo morfologijo in mikrotrdoto (nanoidentifikacija). Prisotnost i-faze smo potrdili z metodama XRD in TEM. Klju~ne besede: Al-Cu-Fe, kvazikristal, metalografija In the ternary system Al-Cu-Fe an i-phase (icosahedral quasicrystal) is present. It is thermodynamically stable and a part of the equilibrium phase diagram. However, according to the chemical composition and conditions during the solidification and heat treatment, a considerable number of intermetallic phases can be in stabile or metastabile equilibrium with the i-phase. Consequently, synthesis of onephase quasicrystalline alloy is possible only in a narrow concentration range and after appropriate heat treatment. In the investigation, alloy Al64,4Cu22,5Fe13,1 was synthesized and heat treated to increase the fraction of the i-phase. The presence of phases and other microstructural characteristics were determined using appropriate metallographic preparation methods, light microscopy (LM), scanning electron microscopy (SEM), microchemical analyses (EDS), X-ray diffraction (XRD), transmission electron microscopy (TEM) and microhardness measurements. The presence of quasicrystalline phase i was clearly confirmed using XRD and TEM. Key words: Al-Cu-Fe, quasicrystal, metallography 1 UVOD Leta 1984 so Shechtman in sodelavci 1 objavili ~lanek o novi snovi, ki ima poseben elektronski uklonski vzorec – red dolgega dosega, vendar brez periodi~nosti. To snov so kasneje poimenovali kvazikristalna snov. H. R. Trebin v knjigi Quasicrystals 2 trdi, da je kvazikri- stalno stanje tretje stanje trdnih snovi poleg kristalnega in amorfnega. Atomi so urejeno razporejeni, toda z rotacijskimi simetrijami, ki imajo pet-, osem-, deset- ali dvanajst{tevne osi, ki jih nimajo snovi v kristalnem stanju. [tevilne raziskave kvazikristalnih faz temeljijo na poglobljenem {tudiju tvorbe teh faz v zlitinah Al-Cu-Fe, saj so elementi, ki jih sestavljajo lahko dostopni, poceni in niso strupeni. Poleg tega spada zlitina Al-Cu-Fe med najbolj primerne za preu~evanje nastanka kvazikristalnih faz, s tem pa tudi mo`nosti za uporabo kvazikristalnih zlitin. Velika verjetnost nastanka kompleksnih ternarnih spojin v zlitinah Al-Cu-Fe namre~ izhaja `e iz kon- stitucije robnih binarnih zlitinskih sistemov z ve~jim {tevilom binarnih intermetalnih spojin. Kristalografske zna~ilnosti spojin v sistemu Al-Cu- Fe v obmo~ju i-faze so navedene v tabeli 1. Ugotovljeno je, da v ternarnem sistemu nastopa poleg kristalnih intermetalnih spojin tudi kvazikristalna i-faza, za katero je zna~ilna ploskovno centrirana ikozaedri~na kvazi- kristalna zgradba (FCI) 3. Ta faza nastaja v skladu s ternarno peritekti~no reakcijo L +  +   i pri tempe- raturi 882 °C, kjer ima talina sestavo v to~ki P1, kar je razvidno iz vertikalnega prereza ternarnega sistema zli- tine Al-Cu-Fe pri x(Cu) = 25 % Cu (slika 1). Konsti- tucija vertikalnega prereza ternarnega sistema Al-Cu-Fe pri konstantni koncentraciji bakra (x(Cu) = 25 %) nazorno prikazuje razmere pri ternarni peritekti~ni reakciji in potrjuje ugotovitve raziskovalcev 3, da ima kvazikristalna i-faza ozko koncentracijsko obmo~je obstojnosti (blizu Al62Cu25,5Fe12,5), ki se spreminja s temperaturo in je prikazano na sliki 1 s ~rtkano ~rto. Iz izotermnega prereza aluminijevega kota ternarnega sistema Al-Cu-Fe pri 700 °C (slika 2) je razviden obstoj ve~ heterogenih ravnote`ij in intermetalnih spojin, ki v njih sodelujejo. Iz izotermnega prereza pri 700 °C (slika 2) je raz- vidno, da so lahko v ravnote`jem stanju z i-fazo {tiri faze: , ,  in talina (L), v metastabilnih stanjih pa se pojavijo {e dodatne faze. Cilj na{ega dela je bil raziskati faze, ki se pojavljajo v zlitini Al62Cu25,5Fe12,5 po litju in toplotni obdelavi. Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 271 UDK 669.71'3'1:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)271(2007) 2 EKSPERIMENTALNO DELO Zlitino Al-Cu-Fe smo izdelali v vakuumski pe~i LEYBOLD-HEREAUS IS pri tlaku 10–2 bar. Kot vlo`ek smo uporabili aluminij (w = 99,99 % in baker (w = 99,99 %) ter predzlitino AlFe45. Lili smo v jekleno kokilo v za{~itni atmosferi argona. Izdelana zlitina je bila krhka, zato smo jo lahko zdrobili na majhne ko{~ke, ki smo jih toplotno obdelali. Zlitino smo kemi~no analizirali z metodo ICP-AES (opti~na emisijska spektrometrija z indukcijsko sklopljeno plazmo), sestava zlitine je nave- dena v tabeli 2. Kemijska sestava je vedno navedena v mno`inskih dele`ih. T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 272 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 Tabela 1: Najpomembnej{e binarne in ternarne faze v sistemu Al-Cu-Fe v sose{~ini i-faze 8 Table 1: The most important binary and ternary phases in system Al-Cu-Fe in the vicinity of i-phase 8 Stehiometri~na formula spojine Stechiometric formula of the compound Kemi~na sestava Chemical composition x/% Mre`ni parametri Lattice parameters Kristalni sistem Crystal system Pearsonov simbol Pearson’s symbol  – AlCu 49,8 − 52,4 Cu a = 0,4015 nm b = 1,202 nm c = 0,8652 Ortorombi~ni Orthorombic oP16  – Al3Cu4 590−530 °C 55,2 − 59,8 Cu a = 0,81 nm c = 1,000 nm Heksagonalni Hexagonal hP42  – Al3Cu4 <570 °C 55,2 − 56,3 Cu a = 0,707 nm b = 0,408 nm c = 1,002 nm = 90,63° Monoklinski Monoclinic m*21 – Al2Cu 31,9 − 33 Cu a = 0,6063 nmc = 0,487 nm Tetragonalni Tetragonal tI12  – Al13Fe4 od Al78Fe22 do Al73Cu5Fe22 a = 1,5489 nm b = 0,8083 nm c = 1,247 nm = 107,72° Monoklinski Monoclinic mC102  - Al(FeCu) 22,0 − 54,5 Al a = 0,2909 nm Kubi~na telesno centrirana Cubic body centered cP2  – Al7Cu2Fe 70 Al a = 0,6336 nmc = 1,487 nm Urejeni tetragonalni Ordered tetragonal tP40 i – Al62Cu25,5Fe12,5 a = 0,63346 Ikozaedri~ni ploskovno centrirani Icosahedral phase centred –35m Slika 2: Izotermni prerez ternarnega faznega diagrama Al-Cu-Fe pri 700 °C v aluminijevem kotu 7 Figure 2: Isothermal section of the ternary Al-Cu-Fe system in the Al-rich corner at 700 °C 7 Slika 1: Vertikalni prerez ternarnega sistema Al-Cu-Fe pri 25 % Cu 6 Figure 1: Vertical cross-section of the ternary system Al-Cu-Fe at 25 % Cu 6 Tabela 2: Kemijska sestava preiskane zlitine Al-Cu-Fe Table 2: Chemical composition of the investigated Al-Cu-Fe alloy w/% / x(Al) % w/% / x(Cu) % w/% / x(Fe) % ZLITINA 2 Al-Cu-Fe 42,9 / 64,4 37 / 22,5 20 / 13,1 Vzorce zlitine Al64,4Cu22,5Fe13,1 smo toplotno obdelali v cevni pe~i v za{~itni argonski atmosferi. @arjenje je potekalo 24 h pri 750 °C (nato hitro ohlajanje) in 100 h pri 780 °C (vzorec je bil zataljen v kremenovo cevko v argonski atmosferi in v njej tudi toplotno obdelan, ohlajanje je bilo po~asno). Zlitino Al-Cu-Fe smo preiskali v izhodnem litem stanju in v toplotno obdelanem. Za karakterizacijo zlitine smo uporabili sodobne raziskovalne metode, kot so: svetlobna mikroskopija (SM), elektronska vrsti~na mikroskopija (SEM), elektronska presevna mikroskopija (TEM), mikrokemi~na spektroskopska analiza (EDS) in rentgenska fazna analiza (XRD). Pri nekaterih vzorcih smo tudi merili mikrotrdoto po Vickersu HV 0,05. 3 REZULTATI IN DISKUSIJA Zlitina Al-Cu-Fe v za~etnem litem stanju ima kemijsko sestavo v mno`inskih dele`ih 64,4 % Al, 22,5 % Cu in 13,1 % Fe. S svetlobno in vrsti~no elektronsko mikroskopijo smo ugotovili {tiri faze: i, ,  in  (slika 3 T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 273 Slika 3: Mikrostruktura in EDS-spektri faz v zlitini Al64,4Cu22,5Fe13,1 v za~etnem litem stanju: a) SEM-posnetek mikrostrukture zlitine, b) SEM-posnetek monokvazikristala, ki je nastal z neovirano rastjo, c) EDS-spekter faze , d) EDS-spekter faze i, e) EDS-spekter faze , f) EDS-spekter faze  Figure 3: Microstructure and EDS-spectra of phases present in the alloy Al64,4Cu22,5Fe13,1 in the as-cast condition: a) SEM-micrograph, b) SEM-micrograph of a quasicrystal with pentagonal octahedral morphology. c) EDS-spectrum of phase , d) EDS-spectrum of i-phase, e) EDS-spectrum of phase , f) EDS-spectrum of phase  a). Kemijsko sestavo posameznih faz smo ugotovili z analizo EDS (slike 3 c do f). Na prelomni povr{ini so vidni monokvazikristali (slika 3 b), ki so nastali s prosto rastjo i-faze s kemijsko sestavo 64,4 % Al, 12,4 % Fe in 23,2 % Cu (slika 4 d). Navzo~e so {e faze  (Al(Cu,Fe) (slika 4 c),  (Al13Fe4) (slika 4 e) in -(Al3Cu4) (slika 3 f). Primarno se je izlo~ala faza , vendar je bila njena koli~ina majhna, ker je sestava zlitine zelo blizu evtekti~nemu `lebu, zato kmalu pote~e binarna evtekti~na reakcija L   +  (vertikalni prerez, slika 1). Preostala talina se porabi pri ternarni peritekti~ni reakciji, ko nastane i-faza. Naj- svetlej{a faza  je bogata z bakrom in se nahaja v meddendritnem prostoru. 3.1 Mikrostruktura zlitine Al64,4Cu22,5Fe13,1 po DTA Vzorec zlitine smo kontrolirano segrevali do 1100 °C in ohlajali do sobne temperature s hitrostjo 10 K/min. S svetlobno in vrsti~no elektronsko mikroskopijo smo ugotovili, da so v mikrostrukturi vzorca po DTA {tiri faze: i, ,  in ' (sliki 4 a, b). T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 274 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 Slika 4: Mikrostruktura in EDS-spektri faz v zlitini Al64,4Cu22,5Fe13,1 po DTA: a, b) SEM posnetka mikrostrukture, c) EDS-spekter faze , d) EDS-spekter faze i, e) EDS-spekter faze , f) EDS-spekter faze ' Figure 4: Microstructure and EDS-spectra of phases present in the alloy Al64,4Cu22,5Fe13,1 after DTA: a, b) SEM-micrographs, c) EDS-spectrum of phase , d) EDS-spectrum of i-phase, e) EDS-spectrum of phase , f) EDS-spectrum of phase ' Pri strjevanju je primarno nastajala faza , kasneje je potekla {e binarna evtekti~na reakcija, kjer sta se iz taline hkrati izlo~ali fazi  in . Glede na verikalni prerez ternarnega faznega diagrama (slika 1) pote~e pri ravnote`nem strjevanju peritekti~na reakcija, kjer se tvori i-faza po reakciji:  +  + L  i. Pri ni`jih temperaturah i-faza evtektoidno razpade v fazi  in . Na sliki 4 a je  prikazana v dendritni obliki, ki je obdana z i-fazo. Na sliki 4 b je  prikazana v obliki peresastih delcev, ki je obdana s fazo . Prisotna je {e faza β' (naj- svetlej{a faza na sliki 4 a), ki ima glede na EDS-analizo (slika 4 f) od vseh analiziranih faz v tem vzorcu najve~ji mno`inski dele` bakra (42,7%). Strjevanje zlitine v za~etnem stanju in pri DTA je bilo neravnote`no, le da je bilo pri DTA po~asnej{e. Razlika se poka`e pri strjevanju preostale taline pri ni`jih temperaturah. V za~etnem litem stanju nastane faza -(Al3Cu4), ki ima okoli 3 % Fe in okoli 53 % Cu. Po drugi strani nastane v DTA-vzorcu, ki se je kontrolirano ohlajal s hitrostjo 10 K/min, nazadnje faza ' s sestavo okoli 10 % Fe in 43 % Cu. Nastanek faze ' je povezan s pomikom sestave zlitine k manj{im dele`em `eleza in z ote`enim nastankom -faze 4. 3.2 Toplotno obdelano stanje 24 h pri 750 °C Vzorec zlitine Al64,4Cu22,5Fe13,1 je bil 24 h `arjen pri 750 °C in hitro ohlajen. Na sliki 5a je prikazan vzorec v poliranem stanju. Zaradi toplotno aktiviranih procesov v T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 275 Slika 5: Mikrostruktura in EDS-spektri faz v zlitini Al64,4Cu22,5Fe13,1 v toplotno obdelanem stanju (24 h pri 750 °C): a) pregledni SEM-posnetek, b) SEM-posnetek, c) EDS-spekter faze , d) EDS-spekter faze i in e) EDS-spekter faze  Figure 5: Microstructure and EDS-spectra of phases present in the alloy Al64,4Cu22,5Fe13,1 after heat treatment (24 h at 750 °C): a, b) SEM-micrographs, c) EDS-spectrum of phase , d) EDS-spectrum of i-phase, e) EDS-spectrum of phase  trdnem je celoten vzorec mo~no porozen, pojavi pa se tudi sprememba fazne sestave. Mikrostruktura je sestavljena iz najsvetlej{e faze  (Al(Cu,Fe))(slika 5 c), i-faze (slika 5 d) in faze  (Al13Fe4) (slika 5 e). Glede na vertikalni prerez ternar- nega faznega diagrama Al-Cu-Fe (slika 1) je bila tempe- ratura 750 °C za zlitino Al64,4Cu22,5Fe13,1 prenizka in ~as `arjenja prekratek, da bi dosegli enofazno podro~je i-faze ali dvofazno podro~je ( + i). Prisotnost faz smo potrdili z EDS-analizo in rentgenska fazno analizo (XRD) (slika 6). 3.3 Toplotno obdelano stanje 100 h pri 780 °C Vzorec zlitine smo `arili tudi 100 h na temperaturi 780 °C in po~asi ohladili, kar je omogo~ilo pribli`anje ravnote`nim razmeram ohlajanja. Majhna temperaturna razlika (30 °C) pri `arjenju in dalj{i ~as `arjenja glede na toplotno obdelavo pri 750 °C povzro~ita velike razlike v mikrostrukturi. Po kon~anem `arjenju dobimo enofazno kvazikristalno strukturo (sliki 7 a in b). Zna~ilno kvazikristalno strukturo smo potrdili s presevno elektronsko mikroskopijo (TEM) (slika 8) in meritvami mikrotrdote, prisotnost kvazikristalne i-faze pa tudi z rentgensko fazno analizo (XRD) (slika 6). Z uklonske slike 8 je razvidna zna~ilnost kvazi- kristalne strukture – ni periodi~nega vzorca. V {estih smereh se razdalje med uklonskimi lisami pove~ujejo s , ki je (1 + 5)/2. Uklonske lise so med seboj oddaljene v razmerju   1,6. Tako razmerje velja med razdaljami BC/AB, DC/CB in DE/CD (slika 8). Z merjenjem mikrotrdote smo ugotovili, da je HV 0,05 942 ± 15, kar je skladno s podatki iz literature 5. 4 SKLEPI Zlitina Al64,4Cu22,5Fe13,1 vsebuje v za~etnem litem stanju {tiri faze: , i,  in . Enako {tevilo faz je bilo tudi po nadzorovanem ohlajanju z 10 K/min, le da se je namesto faze  pojavila faza '. Pri po~asnem ohlajanju zlitine v za~etnem litem stanju so bile ustvarjene razmere za neovirano rast kvazikristalov iz taline. Nastali monokvazikristali so imeli obliko pentagonalnega dodekaedra. T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 276 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 Slika 8: Uklonska slika (TEM) na zdrobljenih delcih zlitine Al64,4Cu22,5Fe13,1 v toplotno obdelanem stanju (100 h pri 780 °C, po~asno ohlajanje) Figure 8: Selected area diffraction pattern of broken particles of the alloy Al64,4Cu22,5Fe13,1 in the heat-treated condition (100 h at 780 °C, slow cooling) Θ β β β λ Slika 6: Rentgenska fazna analiza (XRD) toplotnoobdelane zlitine Al64,4Cu22,5Fe13,1 Figure 6: X-ray diffraction of the heat-treated alloy Al64,4Cu22,5Fe13,1 Slika 7: Mikrostruktura zlitine Al64,4Cu22,5Fe13,1 v toplotno obdela- nem stanju (100 h pri 780 °C): a) pregledni SM posnetek, b) SEM posnetek Figure 7: a) Optical micrograph and b) SEM-micrograph of the alloy Al64,4Cu22,5Fe13,1 in heat-treated condition (100 h at 780 °C) Z ustrezno toplotno obdelavo se je dele` faze i pove~al. Po 24-urnem `arjenju zlitine Al64,4Cu22,5Fe13,1 na 750 °C se je mo~no pove~al dele` faze i, vendar pa sta bili v zlitini {e vedno fazi  in . Toda po 100-urnem `arjenju na 780 °C in po~asnem ohlajanju do sobne temperature, je ternarna peritekti~na reakcija L +  +   i potekla v celoti, tako je nastala enofazna kvazikristalna mikrostruktura. Kvazikristalna faza je imela trdoto HV okoli 1000. Na osnovi raziskav lahko sklenemo, da lahko v zlitini Al64,4Cu22,5Fe13,1 s primerno toplotno obdelavo dose`emo enofazno kvazikristalno mikrostrukturo. Prisotnost kvazikristalne faze i v raziskovanih zlitinah iz sistema Al-Cu-Fe smo zanesljivo potrdili s presevno elektronsko mikroskopijo in rentgensko fazno analizo. 5 LITERATURA 1 D. S. Shechtman, I. Blech, D. Gratias, J. W. Cahn: Phy. Rev. Lett., 53 (1984), 1951–1953 2 Quasicrystals, Structure and Physical Properties. Edited by Hans-Rainer Trebin: Wiley-VCH GmbH & Co. KgaA, Weinheim, 2003, 2–23 3 A. P. Tsai, A. Inoue, T. A. Masumoto, Jpn. J. Appl. Phys., 26 (1987), L1505–L1507 4 L. Zhang, R. Lück, Z. Metallkunde 94 (2003) 2, 774–781 5 E. Giacometti, N. Baluc, J. Bonneville and J. Rabier, Scripta Materialia, 41 (1999) 9, 989–994 6 L. Zhang, R. Lück Z. Metallkunde 94 (2003)2, 98–107 7 L. Zhang, R. Lück, Z. Metallkunde 94 (2003)2, 108–115 8 Ternary Alloys. Edited by G. Petzow and G. Effenberg: VCH Ver- lagsgesellscaft, Weinheim, 1988, 475–489, 361–362 T. BON^INA ET AL.: FAZE V KVAZIKRISTALNI ZLITINI Al64,4Cu22,5Fe13,1 Materiali in tehnologije / Materials and technology 41 (2007) 6, 271–277 277 I. [KULJ ET AL.: HYDROGEN ABSORPTION BY Ti–Zr–Ni-BASED ALLOYS HYDROGEN ABSORPTION BY Ti–Zr–Ni-BASED ALLOYS ABSORPCIJA VODIKA V ZLITINAH Ti–Zr–Ni Irena [kulj1, Andra` Kocjan2, Paul J. McGuiness2, Borivoj [u{tar{i~1 1Institute for Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia 2Jo`ef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia irena.skuljimt.si Prejem rokopisa – received: 2007-05-15; sprejem za objavo – accepted for publication: 2007-11-07 Some transition metals and their alloys have the ability to reversibly absorb considerable amounts of hydrogen. The amount of absorbed hydrogen and the absorption kinetics depend on interactions between the hydrogen atoms and the alloy. Titanium and zirconium show a high affinity for hydrogen, and Ti–Zr–Ni alloys with either amorphous or quasicrystalline structures have proved to be excellent absorbers of hydrogen. In this study we have focused on processing Ti–Zr–Ni ribbons with four different compositions and investigating their hydrogenation behaviour. A melt-spinning process was used as the alloy-preparation technique, and samples from each composition were examined before and after the hydrogenation process. The ribbons were analysed by X-ray diffraction (XRD) and examined with a scanning electron microscope (SEM) equipped with an energy-dispersive X-ray spectrometer (EDS). Only the sample with the lowest nickel concentration was found to absorb any significant quantity of hydrogen under the applied experimental conditions. Keywords: Quasicrystals, Ti–Zr–Ni alloys, Hydrogenation Nekatere kovine prehoda in njihove zlitine imajo sposobnost reverzibilne absorpcije vodika. Koli~ina absorbiranega vodika in kinetika absorpcije sta odvisni od interakcij med atomi vodika in atomi zlitine. Titan in cirkonij imata visoko afiniteto do vodika in zato zlitina Ti–Zr–Ni z amorfno ali kvazikristalno strukturo odli~no absorbira vodik. V tem delu smo se osredinili na postopek izdelave trakov Ti–Zr–Ni s {tirimi razli~nimi sestavami in preu~ili postopek hidrogenacije. Kot postopek izdelave trakov smo uporabili litje taline na vrte~i se valj. Hitrostrjene trakove razli~nih sestav smo preu~ili pred hidrogenacijo in po njej. Trakove smo analizirali z rentgensko spektroskopijo (XRD) in vrsti~nim elektronskim mikroskopom (SEM) z energijsko disperzijo rentgenskih `arkov (EDS). Le vzorec z najni`jo vsebnostjo niklja je pri danih pogojih absorbiral znatnej{e koli~ine vodika. Klju~ne besede: kvazikristali, zlitine Ti–Zr–Ni, hidrogenacija 1 INTRODUCTION Amorphous and quasicrystalline alloys have recently attracted a lot of attention for hydrogen-storage appli- cations. Much of the research has concentrated on Ti- and Zr-based alloys and related materials. These alloys mostly contain an i-phase, which has an icosahedral quasicrystalline structure. These quasicrystals are able to absorb/desorb considerable amounts of hydrogen1, 2, 3, 4. Containing more tetrahedral sites in their structure than other crystals, and with Ti and Zr having high affinities for hydrogen, makes these alloys potentially excellent hydrogen-storage materials 5. The i-type quasicrystal is formed either by rapid quenching or solid-state trans- formation at 500–600 °C, generally leading to a micro- structure of quasicrystal and crystal phases with a grain size of several microns 1. The melt-spinning process is a technique used for the rapid cooling of molten metals and alloys. It is used to develop materials that require extremely high cooling rates in order to form, for example, metallic glasses. The cooling rates achieved are of the order of 104–107 K/s. Therefore, the process can be successfully used for the preparation of amorphous and quasicrystalline samples through the rapid solidification of a molten alloy on a cold spinning wheel made of copper. The cooling rate of the process has also been defined for this type of alloy with additions of silicon 6. The hydrogenation of these alloys was successful at high temperatures and low pressures. However, only 10 % of the absorbed hydrogen can be evolved from the alloy during the heat treatment 7. The aim of this study was to optimise the melt-spinning process for the production of ribbons with amorphous and quasicrystalline structures. 2 EXPERIMENTAL DETAILS Samples with the compositions in Table 1 were prepared from Ti78.6Ni21.4, Zr65.9Ni34.1 and Ti60Zr40 binary-alloy ingots. The starting alloys were crushed into smaller lumps (3 cm), and suitable amounts of each composition were put in to a melting crucible in an induction furnace. The crucible we used was made of graphite, coated with a zirconia-water suspension that was dried before use. The coating was used to prevent the formation of carbides. Induction melting in a vacuum was used to prepare the pre-alloy with the desired composition. The homogenous pre-alloys were then re-melted and spun to obtain the ribbons. The melt-spinning process was carried out with a 25.2 m/s wheel speed, and the size of the crucible’s nozzle was Materiali in tehnologije / Materials and technology 41 (2007) 6, 279–282 279 UDK 669'295'296'24:669.788 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)279(2007) 2.2 mm (). To obtain melt-spun ribbons with the desired structure it was necessary to choose a suitable wheel speed and the correct size of nozzle. The process was carried out in a slight under-pressure of argon. Some of the prepared ribbons were examined and others were hydrogenated. The hydrogenation process was carried out in a furnace in a hydrogen atmosphere. First, the ribbons were put in to the furnace and the furnace was sealed. The furnace was then evacuated and refilled with an 8-bar over-pressure of hydrogen. The furnace was heated up to 350 °C, and the temperature was maintained at this temperature. When the absorption was complete the furnace was allowed to cool down and the remaining hydrogen was released from the system. All the ribbons, as-melt-spun and hydrogenated, were investigated with an X-ray diffractometer (XRD) and examined with a scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS). The samples for the SEM examinations and the EDS analyses were ground and polished prior to the examination. 3 RESULTS AND DISCUSSION The cross-sections of the samples obtained using an optical microscope can be seen in Figure 1. A com- parison was made between sample A1, with the highest amount-of-substance fraction of Ni present in the composition (26.3 %), and sample A2, with the smallest amount-of-substance fraction of Ni present in the I. [KULJ ET AL.: HYDROGEN ABSORPTION BY Ti–Zr–Ni-BASED ALLOYS 280 Materiali in tehnologije / Materials and technology 41 (2007) 6, 279–282 Figure 2: Cross-sectional views of the as-melt-spun ribbons with compositions A1 (a), A2 (b), A3 (c) and A4 (d) Slika 2: Pre~ni prerez hitrostrjenih trakov s sestavami A1 (a), A2 (b), A3 (c) in A4 (d) Figure 1: Cross-sectional view of the ribbon with the most (a) and the least (b) amount of Ni present in the microstructure Slika 1: Pre~ni prerez traku z najve~jo (a) in najmanj{o (b) vsebnostjo Ni v mikrostrukturi composition (17 %). It seems that the sample with less Ni in the overall composition forms a two-phase micro- structure and the sample with more Ni forms a single- phased microstructure. The assumption was confirmed by the SEM investigation. From the images shown in Figure 2 it is clear that sample A2 (Figure 2b) is the only one consisting of two phases. All the other samples, with higher Ni concentrations, have single-phase micro- structures. All the samples and phases were analysed with EDS, and the results are collected in Table 2. All the phase compositions are in accordance with the starting compositions listed in Table 1. The ribbons with the A2 composition are the only samples containing Ti–Zr particles in a matrix structure, which again agrees with the original composition (Table 1). The particles can be seen in Figures 1 b and 2 b. XRD scans belonging to the as-melt-spun samples are collected in Figure 3. It is clear that sample A3 was the only sample that had an amorphous structure. All other samples, A1, A2 and A4, show patterns that indicate the presence of crystalline phases. From this we can conclude that the solidification resulting from melt spinning at 25.2 m/s with a 2.2 mm nozzle was insufficiently rapid to form an amorphous structure except in the sample with the lowest amount of Ni. The samples with compositions A1, A2 and A4 cooled down I. [KULJ ET AL.: HYDROGEN ABSORPTION BY Ti–Zr–Ni-BASED ALLOYS Materiali in tehnologije / Materials and technology 41 (2007) 6, 279–282 281 Figure 4: XRD scans of as-melt-spun Ti–Zr–Ni ribbons after hydrogenation Slika 4: XRD spektri hitrostrjenih trakov Ti–Zr–Ni po hidrogenaciji Figure 3: XRD scans of as-melt-spun Ti–Zr–Ni ribbons Slika 3: XRD spektri hitrostrjenih trakov Ti–Zr–Ni Figure 5: Ribbons with compositions A1 (a), A2 (b), A3 (c) and A4 (d) after hydrogenation Slika 5: Trakovi s setavo A1 (a), A2 (b), A3 (c) in A4 (d) po hidrogenaciji too slowly and so had time to form a crystalline structure. Based on these results we can state that the conditions used during the melt-spinning process were not suitable for producing ribbons with an amorphous structure. The oversized nozzle as well as a too-slow wheel speed during the melt-spinning process resulted in a relatively slow cooling rate and slow solidification, giving the material time to crystallise. At this stage it is also worth mentioning that the melting temperature of the alloy is increasing with decreasing Ni content. From the Ti–Zr–Ni phase diagram 9 it is clear that when the decreasing Ni content exceeds the amount-of-substance fraction ≈22 % the solidus-liquidus area starts widening. A consequence of this is the presence of solid particles in the melt that could cause difficulties during melt- spinning. The images in Figure 5 and the XRD scans in Figure 4 were obtained from ribbons after hydrogenation. Comparing the scans for A1, A3 and A4 before and after the hydrogenation it is clear that the scans are the same, proving that no hydrides were formed during the hydrogenation process. The compositions in Tables 2 and 3 also confirm that the samples did not absorb any hydrogen. Sample A2 was the only affected sample. When exposed to hydrogen the hydrogen penetrated via grain boundaries through the ribbons and forced the grains to decrepitate 9. Consequence of the process was the formation of a powder with an unaffected compo- sition. It can also be concluded that the hydrogenation conditions were not sufficient for the hydrogenation to proceed. It is most likely that the hydrogen pressure applied to the sample was insufficient. Unfortunately, the existing equipment used for the hydrogenation is not capable of supporting higher pressures. 4 CONCLUSIONS From a range of samples with different amounts of Ni in their compositions only the sample with smallest amount-of-substance fraction Ni (17 %) content forms a two-phase structure. Some Ti–Zr particles are present in the microstructure. All the other compositions formed a single-phase microstructure. The cooling rate of the melt-spun samples was not fast enough to provide the conditions for fast solidifi- cation and, consequently, ribbons with an amorphous structure, except for the sample with least amount of Ti. The conditions under which the hydrogenation process was carried out were found to be insufficient to initiate the formation of hydrides for three of the four samples. Acknowledgements Many thanks go to the Jo`ef Stefan Institute for their support and contribution to the research of which this work forms a part. Thanks also to Slovenian research agency ARRS for their financial support. 5 REFERENCES 1 Davis J. P., Majzoub E. H., Simmons J. M., Kelton K. F.; Mater. Sci. Eng.; 104 (2000), A294 2 Guo X. Q., Louzguine D. V., Yamaura S., Ma L. Q., Sun W., Hase- gawa M., Inoue A.; Mater. Sci. Eng.; A338 (2002), 97 3 Viano A. M., Majzoub E. H., Stroud R. M., Kramer M. J., Misture S. T., Gibbons P. C., Kelton K. F.; Phil. Mag. A; 78 (1998) 1, 131 4 Gibbons P. C., Hennig R. G., Huett V. T., Kelton K. F.; J. non- crystalline Solids; 334&335 (2004), 461 5 Batalla E., Strom-Olsen J. O., Altounian Z., Boothroyd D., Harris R.; J. Mater. Res.; 1 (1986), 765 6 Rud A. D., Schmidt U., Slukhovskii O. I., Zeliska G. M.; J. Alloys Comp.; 373 (2004), 48 7 Viano A. M., Stroud R. M., Gibbons P. C., McDowell A. F., Conradi M. S., Kelton K. F.; Phys. Rew. B; 51(1995) 17, 12026 8 Kelton K. F., Gangopadhyay A. K., Lee G. W., Hyers R. W., Rathz T. J., Robinson M. B., Rogers J.; Studies of nucleation and growth, specific heat and viscosity of undercooled melts of quasicrystal and ploytetrahedral-phase forming alloys, Proc. of NASA Microgravity Materials Science Conference, Huntsville, Alabama, USA, 2000, 327–337 8 McGuiness P. J., Harris, I. R., Rozendaal, E., Ormerod, J., Ward, M. Journal of Materials Science 21(1986) 11, 4107 I. [KULJ ET AL.: HYDROGEN ABSORPTION BY Ti–Zr–Ni-BASED ALLOYS 282 Materiali in tehnologije / Materials and technology 41 (2007) 6, 279–282 Table 1: Compositions Tabela 1: Sestava A1 Ti48Zr26Ni26 A2 Ti50Zr33Ni17 A3 Ti44Zr33Ni23 A4 Ti50Zr29Ni21 Table 2: Compositions of phases forming the microstructure of melt- spun ribbons Tabela 2: Sestave faz, ki tvorijo mikrostrukturo hitrostrjenih trakov Phase x(Ti)/% x(Zr)/% x(Ni)/% A1 Matrix 48.3 ± 0.5 26.1 ± 0.7 25.6 ± 0.5 A2 Matrix 50.1 ± 0.5 33.7 ± 0.7 16.2 ± 0.5 Particle 56.2 ± 0.5 43.8 ± 0.7 / A3 Matrix 43.9 ± 0.5 33.3 ± 0.7 22.8 ± 0.5 A4 Matrix 51.4 ± 0.5 27.8 ± 0.7 20.9 ± 0.5 Table 3: Compositions of the phases forming the microstructure of the hydrogenated ribbons Tabela 3: Sestave faz, ki tvorijo mikrostrukturo hidrogeniranih trakov x(Ti)/% x(Zr)/% x(Ni)/% A1 48.7 ± 0.5 26.1 ± 0.7 25.2 ± 0.5 A2 49.1 ± 0.5 33.3 ± 0.7 17.6 ± 0.5 56.3 ± 0.5 43.7 ± 0.7 / A3 43.2 ± 0.5 33.8 ± 0.7 23.0 ± 0.5 A4 50.3 ± 0.5 28.0 ± 0.7 21.7 ± 0.5 P. JUR^I ET AL.: MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS OVREDNOTENJE MIKROSTRUKTURE HITROSTRJENIH TRAKOV Al-7Cr Peter Jur~i1, Mária Dománková2, Mária Hudáková2, Borivoj [u{tar{i~3 1ECOSOND, Ltd., K Vodárn 531, 257 22 ^er~any, Czech Republic 2STU Trnava, J. Bottu 52, 917 24 Trnava, Slovak Republic 3IMT, Lepi pot 11, 1000 Ljubljana, Slovenia p.jurciseznam.cz Prejem rokopisa – received: 2007-09-20; sprejem za objavo – accepted for publication: 2007-10-18 The use of conventional bulk materials is limited by segregation phenomena, which are generated during the solidification and cannot be eliminated in the solid state. The introduction of rapid-solidification technology (RST) into material processing overcame some of the problems of unacceptable material quality and broadened the range of materials that it is possible to fabricate. The use of conventional ingot metallurgy for the fabrication of aluminium alloys containing a large amount of elements with a low diffusion coefficient is impossible because coarse, hard and brittle intermetallics are formed and the alloys have poor mechanical properties. The use of RST makes it possible to produce these alloys with an improved microstructure; however, before industrial production the structure and properties of the rapidly solidified semi-products as well as the consolidated bulk product must be evaluated systematically. In this paper, melt-spun ribbons, made under various conditions from a binary Al–7%Cr, alloy are investigated. The structure consisted of a supersaturated Al solid solution with a high dislocation density, precipitates of chromium-rich phases and rosette-like spherolites formed from the Al solid solution and the Al7Cr intermetallic phase. The type of phases is related to the processing conditions in only a very limited way. Keywords: rapid solidification, melt spinning, ribbons, microstructure Prakti~na uporaba konvencionalnih kovinskih materialov je omejena zaradi izcejanja zlitinskih elementov. To nastaja zaradi razli~nih vzrokov in pojavov med strjevanjem in ga je prakti~no nemogo~e odpraviti s kasnej{o toplotno obdelavo v trdnem stanju. Hitro strjevanje in njegova vpeljava kot tehnologije za procesiranje razli~nih kovinskih materialov premaguje to oviro in raz{irja izbiro materialov. Uporaba konvencionalne metalurgije ingotov za izdelavo Al zlitin z velikim dele`em legirnih elementov z majhnim difuzijskim koeficientom je prakti~no nemogo~a, ker v mikrostrukturi nastajajo grobi delci trdih in krhkih intermetalnih spojin in imajo zlitine nizke mehanske lastnosti. Uporaba tehnologije hitrega strjevanja omogo~a izdelavo teh zlitin z izbolj{ano mikrostrukturo, pred vpeljavo te tehnologije v redno proizvodnjo pa je treba sistemati~no analizirati strukturo in lastnosti hitrostrjenih polproizvodov ter iz njih izdelanega zgo{~enega izdelka. V tem ~lanku obravnavamo raziskave hitro strjenih trakov zlitine Al z masnim dele`em Cr 7 %, izdelanih pri razli~nih pogojih. Preiskave so pokazale, da mikrostrukturo sestavljajo: s Cr prenasi~ena trdna raztopina Al z veliko gostoto dislokacij ter izlo~ki na Cr bogatih faz, kakor tudi peritektik v obliki rozet nastal iz trne raztopine Al in intermetalne faze Al7Cr. Ugotovili smo, da je oblika in velikost posameznih faz relativno malo odvisna od izbranih pogojev izdelave hitro strjenih trakov. Klju~ne besede: hitro strjevanje, nalivanje na hitrovrte~i se valj, trakovi, mikrostruktura 1 FUNDAMENTAL The industrial use of metallic materials is limited by their microstructure and mechanical properties, characteristics that are greatly influenced by the initial casting operation. During slow cooling in large industrial ingots a considerable amount of segregation takes place due to the different solubilities in the solid and the liquid, and this cannot be improved via a solid-state thermal treatment. Only a rapid solidification can successfully overcome the problems connected with segregation and produce fine-grained, segregation-free materials with an unusual chemical composition and unique mechanical properties. Rapidly solidified (RS) materials differ a great deal from materials with the same chemical composition prepared by conventional casting procedures in terms of the refinement of the main structural constituents 1–3. As a result of non-equilibrium freezing, they may also contain supersaturated phases, metastable intermediate phases or, in limited cases, amorphous constituents 4,5. The results affect the microstructures and the properties of materials, in many cases favourably, and this positive effect of RS on materials’ characteristics has been clearly determined 6–8. One typical example where the rapid solidification is required to obtain material with acceptable properties is the group of aluminium alloys with elements that have a negligible solid solubility and a low diffusion coeffi- cient. In conventionally produced Al alloys, elements like Fe, Ni, V, Cr, etc. are considered as impurities, since they form coarse and brittle aluminides. On the other hand, Al alloys containing elements with a low diffusion coefficient made by the RS technique exhibit an Materiali in tehnologije / Materials and technology 41 (2007) 6, 283–287 283 UDK 669.715:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)283(2007) excellent combination of toughness and elongation and are stable up to relatively high temperatures 7,8. The Al–Cr system is a very typical example. The solid solubility of chromium in aluminium is very low. In slowly solidified material, for example, chromium forms large needle-like as well as branch-like particles of AlxCry intermetallics, which have a deleterious effect on the mechanical properties. On the other hand, powder- metallurgy materials based on this binary system, with the addition of some other elements, can easily achieve an ultimate tensile strength up to 600 MPa during an elongation of several percents 9,10. 2 EXPERIMENTAL The Al7Cr alloy was prepared from technically pure aluminium and chromium. The mixture made from raw materials was molten in a vacuum furnace and overheated up to 1150 °C in order to eliminate the occurrence of large and hardly soluble intermetallics in the melt. Melt-spun ribbons were prepared in an experimental device, i.e., the Melt-Spinner M-10, in IMT, Ljubljana. The metals were melted under an argon overpressure. The rotation speed of the copper wheel ranged between 900 r/s and 1350 r/s (16.8 ms–1 and 25.2 ms–1). Other important parameters of the melt-spinning process are given in Table 1. The microstructure of the alloy was investigated using light microscopy (slowly solidified material) and transmission electron microscopy (melt-spun ribbons). The microstructure of the slowly solidified material was revealed by using the Dix-Keller reagent. Thin foils were prepared directly from the rapidly solidified ribbons using a TENUPOL 2® device. A mixture of 30 % nitric acid and 70 % methanol was used as an etching agent. The thinning was carried out at a temperature of –30 °C and a bias of 19 V. 3 RESULTS AND DISCUSSION The microstructure of the slowly solidified material in an as-cast ingot is shown in Figures 1 and 2. Figure 1 shows an optical micrograph of the slowly solidified alloy with star-shaped formations of intermetallics, and Figure 2 shows the slowly solidified alloy with semi- globular particles of intermetallics. The alloy has a dendritic solidification morphology composed of a relatively large amount of intermetallic phases, which differ from each other mainly in shape and size. Some of them have a globular, convex shape, which indicates primary crystallization from the melt. In some cases, star-shaped formations are found in the micro- P. JUR^I ET AL.: MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS 284 Materiali in tehnologije / Materials and technology 41 (2007) 6, 283–287 Table 1: Important parameters of the melt spinning process Tabela 1: Pomebni parametri procesa hitrega strjevanja Sample designation Superheating of the melt Induction heating power Atmosphere (vacuum + Ar 5.9) Wheel speed Nozzle diameter °C m/s mm Al-1 1030 3 kW to 500 °C 7 kW to 1030 °C (heating rate approx. 70 °C/min.) 60 kPa overpressure of argon 16.8 2.2 Al-2 19.6 0.8 Al-3 25.2 0.8 Figure 2: Optical micrograph of the slowly solidified alloy with semi-globular particles of intermetallics Slika 2: Opti~ni posnetek po~asi strjene zlitine s polkroglastimi inter- metalnimi spojinami Figure 1: Optical micrograph of the slowly solidified alloy with star-shaped formations of intermetallics Slika 1: Opti~ni posnetek po~asi strjene zlitine z zvezdastimi interme- talnimi spojinami structure, Figure 1. Their occurrence can be related to the primary crystallization and a peritectic reaction between the intermetallic and the Al solid solution. From the binary Al–Cr equilibrium diagram only the Al7Cr intermetallic would be expected for our chosen composition, although with an increased Cr content other compounds would also be possible (Al11Cr2, Al4Cr, Al3Cr, Al9Cr4, etc.) 11,18. All these intermetallics are non- stoichiometric compounds (Bertholides) with a relatively narrow range of possible compositions. X-ray diffraction fixed the Al solid solution with sharp diffraction lines. This indicates that during a slow solidification no supersaturation of the solid solution occurred, Figures 3 and 4. The second phase was iden- tified as the Al13Cr2 compound, Figure 3. In the binary diagram reported in 11 there is only the isoconcentration Al7Cr phase. In reference 12, it is suggested that the phase is stable in a given concentration range. Taking into account the fact that the investigated alloy has a lower chromium content than the compound, the stoichiometry Al13Cr2 may correspond to the lower limit of the range for the Al7Cr compound. In addition, an intermetallic phase with the same stoichiometry was also found by Selke 13 in the bulk alloy Al–15 % Cr. In this alloy the chromium content is also below the concentration range of the phase Al7Cr 12; it is, however, twice as high as in the alloy investigated in this work. Pearson´s Handbook also mentions the Al45Cr7 phase. The stoichiometric ratio of 45:7 is between that of 13:2 and 7:1 and, with respect to the actual chemical composition of the alloy, the probability of its occurrence is lower than that for the Al13Cr2. Therefore, identifying this intermetallic as Al13Cr2 is considered to be correct. Figure 5 shows one of the features of the specimen Al-1. This type is represented by a primary crystallized rosette-like particle having a size of about 250 nm. Electron diffraction patterns fixed this phase as Al11Cr2 aluminide, in good agreement with the Al–Cr binary equilibrium diagram 12, where the Al11Cr2 phase is in equilibrium with the residual melt above 785 °C. In the slowly solidified alloy, the Al11Cr2 phase decomposes normally to the Al solid solution with a negligible Cr content and the phase Al7Cr. However, if the solidi- fication rate is rapid enough the phase can be conserved in the alloy down to room temperature. Figure 6 shows the aluminium solid-solution matrix with a relatively high dislocation density in the same specimen. The electron diffraction patterns revealed a P. JUR^I ET AL.: MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS Materiali in tehnologije / Materials and technology 41 (2007) 6, 283–287 285 Figure 4: X-ray patterns from the slowly solidified material: red lines, Al; green lines, Al45Cr7 Slika 4: Rentgenska difrakcijska slika po~asi strjene zlitine: Al – rde- ~e ~rte, Al45Cr7 – zelene ~rte Figure 3: X-ray patterns from the slowly solidified material: red lines, Al; blue lines, Al13Cr2 Slika 3: Rentgenska difrakcijska slika po~asi strjene zlitine: Al – rde~e ~rte, Al13Cr2 – modre ~rte Figure 5: A typical primary particle in the specimen Al-1 Slika 5: Zna~ilen primarni delec v vzorcu Al-1 lattice distortion in comparison to the equilibrium situation. Both of these phenomena can be ascribed to the rapid solidification, which produced the supersatu- ration of the solid solution and the enhanced dislocation density in the matrix. The TEM micrograph in Figure 7 shows the rosette- like spherolite from the sample Al-2. The circular- shaped constituent with a diameter of 2 µm consists of the Al solid solution and an intermetallic phase, identified as the Al4Cr aluminide. Figure 8 shows the second constituent, which consists of many semi-globular particles with a size of several tens of nanometers, surrounded by dislocation clusters. These particles correspond very well to the high-temperature δ-phase with a stoichiometry of Al9Cr4. The matrix is formed in a similar way as in the previous specimen from the supersaturated Al solid solution. The last constituent of the microstructure of the investigated melt-spun ribbons is presented in Figure 9. It consists of vermicular Al7Cr precipitates embedded in the Al solid-solution matrix. As confirmed by the electron diffraction, the lattice spacings do not correspond exactly to the equilibrium Al7Cr phase; they are smaller, which suggests that the phase is also partly non-equilibrium. To understand the nature of the phases occurring in thin melt-spun ribbons, the Al–Cr equilibrium diagrams 11,12 must first be taken into consideration. Shunk 14 reported that -Al13Cr2 and -Al11Cr2 are the equilibrium P. JUR^I ET AL.: MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS 286 Materiali in tehnologije / Materials and technology 41 (2007) 6, 283–287 Figure 8: Semi-globular particles in the Al matrix in the specimen Al-2 Slika 8: Polglobularni delci v aluminijevi matici v vzorcu Al-2 Figure 6: The Al solid-solution matrix in the specimen Al-1 Slika 6: Matica trdne raztopine v vzorcu Al-1 Figure 7: Rosette-like spherolite in the specimen Al-2 Slika 7: Rozetasti sferoliti v vzorcu Al-2 phases for the pure Al–Cr system. However, at a higher concentration the Al11Cr2 phase is present at room temperature in the diagram 12, and only above 790 °C can this phase be expected to be an equilibrium aluminide in the alloy. However, it is expected to occur in the melt-spun ribbon, either as a direct consequence of the normal solidification or as a result of non-equili- brium freezing. There is no unequivocal information concerning the Al7Cr aluminide. In a diagram in 11, the Al7Cr phase is shown as an isoconcentric intermetallic compound. On the other hand, the phase is shown in 12 to occur in some narrow concentration range. The stoichiometry 13:2 is slightly smaller that 7:1; thus, -Al13Cr2 can probably be considered as a chromium-poor variant (Bertholide) of the Al7Cr aluminide. Other phases with a higher chromium content (Al4Cr, Al9Cr4) can be expected at room temperature only after slow solidification in the alloys with a high chromium content. At temperatures above 940 °C (Al4Cr) and 1030 °C (Al9Cr4) the occurrence of both phases is also shifted to a lower chromium content. During investigations of alloy processing, theories relating to the structure of the melt were suggested. For instance, it was determined that the structure of the solid alloy is in many cases also conserved to a limited extent in the liquid as a constituent with short-range order. The constituents with long-range orders of atoms were described as clusters on the basis of the "clusters theory" 15-17. It is important that these clusters often have the same compositions as the nearest solid phase. In the rapidly solidified alloy the clusters (or the phases with a similar chemical composition) can be conserved to room temperature, and this is the principal explanation for their occurrence in the ribbons. In addition, Selke 13 suggested that an "i-phase" occurred in the splat-quenched Al85-xCuxCr15 alloys; however, no information about its stoichiometry for the pure Al–Cr system was found so far. Therefore, it is practically impossible to estimate whether it corresponds to the identified phases in the Al–7%Cr alloy or not. 4 CONCLUSIONS The microstructure of the slowly solidified alloy Al–7% Cr consists of the matrix, primary globular or semi-globular dendrites and intermetallic phases. The matrix is an Al solid solution and the intermetallic phases are mainly the compound Al7Cr (Al13Cr2). The microstructure of RS ribbons consists of the matrix with a high dislocation density and of nano- crystalline phases of different size, shape and distribution. The matrix is a supersaturated solid solution in which many different phases occur. Some of them are semi-equilibrium, while the Al9Cr4 and Al4Cr phases are of a non-equilibrium origin. The RS ribbons processed by various conditions differ from each other mainly in terms of the quantity and the occurrence of non-equili- brium phases. 5 REFERENCES 1 Tewari, S. N.: J. Mater. Sci. Lett. 11 (1992), 1020–1022 2 Muller, B. A., Tanner, L. E., Perepezko, J. H.: Mater. Sci. Eng. A150 (1992), 123–132 3 Zhang, X., Atrens, A.: Mater. Sci. Eng. A159 (1992), 243–251 4 Era, H., Kishitake, K., Li, P.: Metall. Trans., 24A, (1993) 3, 751–756 5 Kuoji, M. et al.: J. Mater. Sci. 29 (1994), 1449–1454 6 Jur~i, P.: PhD Thesis, MtF STU Trnava, 1996 (In Czech) 7 Ehrstrom, J. C., Ponesu, A.: Mater. Sci. Engng., A186 (1994), 55–64 8 Premkumar, M. K., Lawley, A., Koczak, M. J.: Mater. Sci. Engng., A174 (1994), 127–139 9 Jones, H.: Mater. Sci. Engng., A375-377 (2004), 104–111 10 Lieblich, M. et al.: Mater. Sci. Techn., 12 (1996), 25–33 11 Web page: http://aluminium.matter.org.uk/content/html/eng/default. asp?catid=79&pageid=-884660481 12 Smithells Metal Reference Book, 8th Edition, Elsevier, 2004 13 Selke, H., Ryder, P. L.: Mater. Sci. Engng., A165 (1993), 81–87 14 Shunk, I. A.: Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York, 1969 15 Stewart, G. W., Benz, C. A.: Phys. Rev., 46 (1934), 703 16 Frenkel, Ja. I.: Kinetic theory of liquids, Leningrad, 1975 (In Russian) 17 Danilov, V. I., Rab~enko, I.V.: @ETF 7, 1937, 1153 (In Russian) 18 M. Hansen, K. Anderko: Constitution of binary diagrams, 1958 P. JUR^I ET AL.: MICROSTRUCTURAL EVALUATION OF RAPIDLY SOLIDIFIED Al–7Cr MELT SPUN RIBBONS Materiali in tehnologije / Materials and technology 41 (2007) 6, 283–287 287 Figure 9: The Al7Cr precipitates in the specimen Al-3 Slika 9: Izlo~ki Al7Cr v vzorcu Al-3 V. DUCMAN, T. KOPAR: THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES VPLIV RAZLI^NIH ODPADKOV NA IZHODNO SUROVINO ZA PROIZVODNJO OPE^NIH IZDELKOV Vilma Ducman, Tinkara Kopar Slovenian National Building and Civil Engineering Institute, Dimi~eva 12, 1000 Ljubljana, Slovenia vilma.ducmanzag.si Prejem rokopisa – received: 2007-07-09; sprejem za objavo – accepted for publication: 2007-10-22 The potential use of four different wastes in the clay-based industry has been studied. The selected wastes were: stone mud from the stone-processing industry, paper mud, sawdust, and sludge from the polishing process for silicate igneous rocks. Waste, depending on its composition, can be used as i) an opening agent, when it contains a large amount of silica or ii) a pore-forming agent, when it contains combustible materials. Mixtures of clay and different amounts of selected waste (up to the mass fraction of 50 %) were prepared. The influence of the addition of different wastes on drying and firing shrinkage was determined, as well as the water absorption, bulk density, bending and compressive strengths of fired samples. We found that the addition of paper mud and sawdust improves the drying process by reinforcing the clay body structure, which counteracts cracking. It also significantly improves the thermal insulation properties. Additions of silica stone mud, as well as granite polishing sludge act as opening agents, which decrease the deformability of green as well as fired specimens. Using waste in the production of clay-based products could represent a quite significant decrease in costs due to the replacement of basic raw materials with waste, and in some cases also the significant improvement of the quality of the final product. Key words: industrial waste, paper mud, sawdust, polishing sludge, silica stone mud, clay products V prispevku so predstavljene mo`nosti uporabe {tirih razli~nih industrijskih odpadkov v ope~ni industriji, in sicer kremenov mulj, papirni mulj, `agovina in odpadek, ki nastaja pri poliranju silikatnih magmatskih kamnin. Odpadke glede na sestavo lahko dodajamo kot dodatke, ki pove~ujejo odprto poroznost, kadar vsebujejo ve~je koli~ine kremena oziroma kadar vsebujejo organske komponente, ki zgorijo med `ganjem, kot dodatke za pove~evanje celotne poroznosti. Pripravili smo me{anice gline in izbranih odpadkov (do masnega dele`a 50 %). Dolo~ili smo vpliv teh dodatkov na skr~ke pri su{enju in `ganju ter na vpijanje vode, gostoto, upogibno in tla~no trdnost `ganih izdelkov. Ugotovili smo, da dodatek papirnega mulja in `agovine izbolj{a proces su{enja, saj u~vrsti strukturo ne`ganih proizvodov in tako prepre~uje nastanek razpok. Prav tako znatno izbolj{a toplotno-izolacijske karakteristike ope~nega izdelka. Kremenov mulj in odpadek pri poliranju granita sta dodatka, ki pove~ujeta odprto poroznost in zmanj{ujeta skr~ke pri su{enju in `ganju in s tem mo`nost nastanka deformacij `ganih proizvodov. Uporaba odpadkov v opekarski industriji lahko pomeni znatno zni`anje stro{kov zaradi delne zamenjave osnovne surovine z odpadnim materialom; prav tako lahko v dolo~enih primerih odpadki izbolj{ajo kakovost kon~nega proizvoda. Klju~ne besede: industrijski odpadki, papirni mulj, `agovina, polirni odpadki, kameni mulj, ope~ni proizvodi 1 INTRODUCTION Different industrial wastes could be quite success- fully used in the clay-based industry; some of them have a positive effect on the process and/or on the final properties, others are simply used to replace the basic raw material and lower the cost of the production. In this investigation four different industrial wastes were checked for their usability in the clay-based industry: silica stone mud from the stone-processing industry, paper mud, sawdust, and granite-like polishing sludge. The sawdust is a by-product of freshly felled timber and therefore contains no solvents, adhesives or other components. Papermaking sludge is a residue in the paper industry, which mainly consists of paper fibres, kaolin, lime and water. The fibrous structure of both wastes has a favourable influence on the stability of the freshly extruded green ware during drying and so counteracts cracking. Besides a favourable effect on the process parameters, sawdust and paper-making sludge are combustible and they can be used as pore-forming agents. Products with increased porosity have better thermal insulation properties1-5. Silica stone mud is a secondary material that remains after the screening of stone aggregate in quarries. It contains a large amount of very fine silica sand, feldspar and clay. Waste granite-like mud is a residue material of polishing operations in the natural stone industry. It contains very fine particles of silicate igneous rocks and a residue of the grinding abrasives. Both wastes could be used as a filler (opening agent) in the clay-based industry due to the large amount of very fine silica sand, which decreases the plasticity of the clay and reduces its shrinkage on drying and firing6-10. In order to study the influence of selected waste materials on the production and final properties of clay bricks, different mixtures of clay and waste were prepared (see Tables 1, 2, and 3 for the composition of the individual mixtures). Materiali in tehnologije / Materials and technology 41 (2007) 6, 289–293 289 UDK 691.421:628.54 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)289(2007) 2 EXPERIMENTAL 2.1 Raw materials and test mixtures 2.1.1 Brick-making clay A Clay taken from the production of masonry bricks from the eastern part of Slovenia was used. The clay can be classified as the chlorite-illitic type, with traces of montmorillonite and around the mass fraction w = 38 % of quartz. The grain size distribution is as follows (in mass fractions): 18.1 % > 20 µm; 36.0 % 2–20 µm, and 45.8 % < 2 µm. 2.1.2 Brick-making clay B Clay taken from the production of masonry bricks from the southern part of Slovenia was used. The clay can be classified as chlorite-illitic type, with around w = 23 % of quartz. The grain size distribution is as follows (in mass fractions): 22.5 % > 20 µm, 40.0 % 2–20 µm, and 37.5 % < 2 µm. 2.1.3 Papermaking sludge The papermaking sludge used was in the form of filter cake, with a water content of approximately w = 52 %. It consists of the following inorganic components: calcite, kaolinite and illite. The loss on ignition at 500 °C is w = 24 % and at 900 °C it is w = 48 %. 2.1.4 Sawdust The sawdust was chopped to pieces of around 1 mm: 11.7 % of the particles were ≥ 1 mm, 76.7 % of the particles were between 1 mm and 0.2 mm, and 11.6 % of the particles were < 0.2 mm. The water content was 19.7 %. Sawdust consists of 99 % combustible substances. The loss on ignition at 500 °C was 98 % and at 900 °C it was 98.6 %. 2.1.5 Silica stone mud The waste silica mud contained about 35 % of quartz, the rest is clay and feldspar. It had the following grain size distribution in mass percent: 16.5 % > 20 µm, 34.9 % 2–20 µm, and 48.6 % < 2 µm. Silica stone mud also contains the swelling mineral montmorillonit. 2.1.6 Granite-like polishing mud The waste granite polishing mud contained about 30 % quartz; the rest is clay, feldspars, carbonates and a residue of SiC polishing tools. It had the following grain size distribution in weight percentage: 8.4 % > 20 µm, 75.3 % 2–20 µm, and 16.3 % < 2 µm. 2.2 Shaping, drying and firing of the test specimens The test specimens were shaped in a laboratory de-airing extruder at a vacuum of about 80 %, i. e., 20 kPa. During extrusion, a proper amount of water was added to the mixtures to avoid surface cracks on the test V. DUCMAN, T. KOPAR: THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES 290 Materiali in tehnologije / Materials and technology 41 (2007) 6, 289–293 Table 2: Mixtures of clay and silica stone mud and average properties of laboratory-made test specimens Tabela 2: Me{anice gline in kremenovega mulja ter pripadajo~e lastnosti laboratorijsko pripravljenih vzorcev MIXTURE B1 B2 B3 B4 Clay content – clay B (w/%) 100 70 50 0 Silica stone mud (w/%) 0 30 50 100 Shaping Water content based on dry mass (w/%) 25.1 31.6 30.0 52.9 Water content based on wet mass (wt/%) 20.0 23.1 23.1 34.6 Shrinkage after drying (%) Measured along the prism length 8.4 10.5 11.6 15.4 Measured across the prism width 7.3 8.8 11.2 15.6 Firing at temperature (°C) (±15) 900 900 900 900 Shrinkage after firing (%) Measured along the prism length 1.4 0.9 1.2 1.6 Measured across the prism width 1.2 1.1 1.4 1.4 Body density after firing (kg/dm3) 1.96 1.90 1.86 1.79 Water absorption (%) 9.2 10.6 12.5 16.8 Clinker point Temp. of firing where water absorption amounts to w = 6 % 1017 1029 1042 1095 Bending strength (MPa) Measured on prisms 22.0 20.7 20.0 6.3* Compressive strength (MPa) Measured on prisms 62.5 52.5 50.7 /* *cracks visible before testing Table 1: Mixtures of clay, papermaking sludge and sawdust and the average properties of laboratory-made test specimens Tabela 1: Me{anice gline, papirnega mulja in `agovine ter pripada- jo~e lastnosti laboratorijsko pripravljenih vzorcev MIXTURE A1 A2 A3 A4 A5 A6 A7 Clay content – clay A (/%) 100 90 80 70 70 70 80 Sawdust (/%) 10 20 30 10 15 Papermaking sludge (/%) 30 20 5 Shaping Water content based on dry mass (w/%) 26.9 28.4 29.3 34.7 34.5 33.8 30.2 Water content based on wet mass (w/%) 21.2 22.1 22.6 25.7 25.7 25.2 23.2 Shrinkage after drying (%) Measured along the prism length 7.7 6.8 6.5 7.2 8.6 / 7.1 Measured across the prism width 6.4 6.5 6.4 7.5 9.6 / 6.7 Firing temperature (°C) (±15) 850 850 920 920 920 910 920 Body density after firing (kg/dm3) 1.81 *1.85 1.69 1.65 1.44 1.59 1.58 1.63 Water absorption (%) 16.7 19.6 21.5 30.5 24.9 25.3 22.2 Compressive strength (MPa) Prisms 23.9 17.4 19.0 10.7 29.3 23.0** 23.4 * fired at 920 °C, ** measured on cylinder / not determined specimens and to maintain a Pff number of 1.4 (±0.1) – see Tables 1, 2, and 3. The test specimens were dried for 7 d in ambient room conditions, followed by 24 h at 60 °C and 8 h at 100 °C in a dryer. The dried samples were then fired for 4 h in a laboratory kiln at selected temperatures using heating rates of 80 °C/h up to 400 °C, and 50 °C/h between 400 °C and the final temperature. These firing conditions are similar to those generally applied in the brick-making industry. 2.3 Test methods The particle size distribution of the tested clays was determined by sieving it down to 90 µm. Below 90 µm, the sedimentation method was applied using a Quanta- chrome Microscan II apparatus. The Quantachrome Microscan II apparatus was also used to determine the particle size distribution of the silica stone mud and the granite polishing mud. The mineral components of both clays and both stone wastes were determined by X-ray diffractometry using a Phillips Norelco apparatus. The linear shrinkage, ceramic body density, water absorption (by boiling test specimens in water for 2 h), and compressive strength were determined on fired samples of (160 × 50 × 25) mm prisms. The pore size and the pore size distribution were measured using Hg porosimetry. The maximum pressure when filling was 206,843 kPa,  130° and  485 · 10–5 N/cm. 3 RESULTS AND DISCUSSION 3.1 The addition of sawdust and papermaking sludge The properties related to shaping, drying and firing are listed in Table 1. The shaping parameter (i.e., the water content) shows that with the additions of pore-forming agents, more water should also be added to mixtures to avoid surface cracks during shaping. The shrinkage after drying is reduced with the addition of sawdust; most significantly for the specimen with the volume fraction of 20 % of sawdust, where the shrinkage after drying is 6.5 % in comparison to pure clay, where the shrinkage is 7.7 % (specimen A1). In contrast, the shrinkage after drying with the addition of papermaking sludge is increased to 8.6 % when the volume fraction of 30 % of the papermaking sludge is added. The decrease in the drying shrinkage is favourable because it reduces the danger of cracking during drying. In Figure 1 the firing analysis from the gradient kiln is presented for pure clay and the mixture A5 containing the mass fraction of 30 % of paper-making sludge. It is clear that the sludge addition increases the water absorption of the specimen, which is to be expected due to its pore-forming action. The shrinkage after firing is lower for the specimen with the sludge, which is a favourable effect since it contributes to the dimensional stability during firing. V. DUCMAN, T. KOPAR: THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES Materiali in tehnologije / Materials and technology 41 (2007) 6, 289–293 291 Table 3: Mixtures of clay and granite stone mud and average pro- perties of laboratory-made test specimens Table 3: Me{anice gline in odpadka od poliranja granite ter pripa- dajo~e lastnosti laboratorijsko pripravljenih vzorcev MIXTURE C1 C2 C3 C4 C5 Clay content – clay B (w/%) 100 95 90 80 70 Granite like stone mud (w/%) 0 5 10 20 30 Shaping Water content based on dry mass (w/%) 26.1 27.0 28.3 32.6 34.6 Water content based on wet mass (w/%) 20.7 21.3 22.1 24.6 25.7 Shrinkage after drying (%) Measured along the prism length 8.7 8.5 9.5 9.8 9.3 Measured across the prism width 7.5 6.6 7.6 8.7 7.8 Firing at temperature (°C) (±15) 915 915 915 915 915 Body density after drying (g/cm3) 2.05 2.02 1.98 1.88 1.80 Shrinkage after firing (%) Measured along the prism length 1.0 1.5 1.4 1.5 1.3 Measured across the prism width 1.1 1.5 1.4 1.7 1.4 Body density after firing (g/cm3) Measured on prisms 2.00 1.96 1.92 1.84 1.76 Water absorption after firing (%) Measured on prisms 8.8 9.6 10.7 12.8 16.6 Clinker point Temp. of firing where water absorption amounts to w = 6 % 1008 1022 1037 1045 1052 Bending strength (MPa) Measured on prisms 16.1 15.6 16.7 13.5 10.4 Compressive strength (MPa) Measured on prisms 86.2 77.6 78.5 71.7 62.4 T L in e a r s h ri n k a g e a n d w a te r a b s o rp ti o n /% emperature, /T °C Figure 1: Determination of the effect of firing temperature on the linear shrinkage and water absorption by firing in a gradient kiln for samples A1 and A5 Slika 1: Vpliv temperature `ganja na skr~ek in vpijanje vode (`ganje v gradientni pe~i) vzorcev A1 in A5 After firing the body density is significantly reduced with the addition of pore-forming agents, especially when sawdust is added. The density of sample A4 with 30 % of added sawdust is 1.44 kg/dm3, whereas the density of the clay without additives is 1.85 kg/dm3. The distribution of the porosity for samples A1, A4, and A5 is presented in Figure 2. In the case of clay without additives (mixture A1) the pore size distribution is uniform. The total porosity in the volume fraction is 32.1 % and over 95 % of the pores are smaller than 3 µm. The addition of sawdust (A4) creates larger pores, where 35 % of the pores are larger than 3 µm and the total porosity is 45.8 %. The addition of papermaking sludge (A5) influences the formation of finer pores, where with a total porosity of 42.3 % over 95 % of the pores are smaller than 3 µm. The reduction in density also influences the compressive strength: from 23.9 MPa for pure clay (fired at 850 °C) to 10.7 for the specimen with 30 % of sawdust. From comparisons of the compressive strengths for specimens A4 and A5 (30 % of sawdust and 30 % of papermaking sludge, respectively), as well as A6 and A7, both with the addition of pore-forming agents, it is clear that with the addition of papermaking sludge, an increase in the compressive strength is observed. This was ascribed to the presence of calcite in the papermaking sludge, as already described in the literature, where it was observed that the addition of 15 % of calcite increases the compressive strength of the clay body by up to 40 %, and at the same time it practically doubles the bending-tensile strength11. As previously described4, the optimal results regarding shrinkage, compressive strengths and body density after firing are obtained with a combination of both pore-forming agents. 3.2 The addition of silica stone mud The properties of the clay mixture with silica stone mud are presented in Table 2. Mixtures with silica stone mud that contain swelling minerals required quite a large content of water for proper shaping. The requirement for a larger amount of water is generally discouraging because this water should be removed in the drying process and this consequently increases the shrinkage after drying: from 8.4 % for pure clay to 15.4 % for pure silica stone mud. A high shrinkage introduces cracks into the green body, which was observed in specimen B4. This sample was prepared from pure silica stone mud, and the water content required for the extrusion amounted to 52.9 %. From Figure 3, which presents the firing analysis in a gradient kiln, it is clear that the addition of mud significantly influences the water absorption’s depen- dence on firing temperature, but only slightly influences the shrinkage after firing. The compressive strength of the fired specimen decreases with the silica stone mud additions, from 62.5 MPa for pure clay to 50.7 MPa when 50 % of silica stone mud is added. 3.3 The addition of granite-like stone mud The properties of the clay mixtures with granite-like polishing waste are presented in Table 3. Mixtures with V. DUCMAN, T. KOPAR: THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES 292 Materiali in tehnologije / Materials and technology 41 (2007) 6, 289–293 0 20 40 60 80 100 120 0,001 0,01 0,1 1 10 100 1000 P , /% ϕ ore diameter, /mmd C u m u la ti v e v o lu m e p ar t A1 A4 A5 Total porosity A1 32,1 % A4 45,8 % A5 42,3 % Figure 2: Pore size distribution for samples A1, A4, and A5 Slika 2: Porazdelitev velikosti por vzorcev A1, A4 in A5 -5,0 0,0 5,0 10,0 15,0 20,0 700 800 900 1000 1100 1200 1300 lin.shrinkage B1 water absorption B1 lin.shrinkage B2 water absorption B2 lin.shrinkage B3 water absorption B3 lin.shrinkage B4 water absorption B4 L in e a r s h ri n k a g e a n d w a te r a b s o rp ti o n /% Temperature, /T °C Figure 3: Determination of the effect of firing temperature on the linear shrinkage and water absorption by firing in a gradient kiln for samples of series B Slika 3: Vpliv temperature `ganja na skr~ek in vpijanje vode (`ganje v gradientni pe~i) vzorcev serije B -5,0 0,0 5,0 10,0 15,0 20,0 25,0 700 800 900 1000 1100 1200 lin.shrinkage C1 lin.shrinkage C2 water absorption C2 lin.shrinkage C3 water absorption C3 lin.shrinkage C4 water absorption C4 lin.shrinkage C5 water absorption C5 T L in e a r s h ri n k a g e a n d w a te r a b s o rp ti o n /% emperature, /T °C Figure 4: Determination of the effect of firing temperature on the linear shrinkage and water absorption by firing in a gradient kiln for samples of series C Slika 4: Vpliv temperature `ganja na skr~ek in vpijanje vode (`ganje v gradientni pe~i) vzorcev serije C waste additions require more water for proper shaping and consequently the linear shrinkage after drying increases from 8.7 % for the mixture without granite-like stone mud to 9.3 % for the mixture containing the mass fraction of 30 % mud. The effect is not pronounced for mixtures with up to 10 % of mud addition. The firing analysis in a gradient kiln showed that the addition of mud significantly increased the water absorption and that it reduced the shrinkage of the basic clay (Figure 4). The addition of mud increased the water absorption of the test specimens and decreased their body density and compressive strength after firing. The decrease in mechanical properties is more significant with a larger amount of mud, when for samples containing up to 10 % of mud addition only a slight decrease in mechanical properties can be observed. 4 CONCLUSIONS Many wastes, depending on their properties and the type of clay, can be successfully used in the brick-making industry. The clay designated as A contains traces of the swelling mineral montmorillonite, which contributes to the sensitivity to cracking of the products on drying. The addition of sawdust and papermaking sludge to brick-making clay favourably influences the process of shaping and drying due to the fibrous structure, which strengthens the green body and prevents the final products from cracking on drying. The addition of sawdust greatly increases the porosity of the fired body and therefore also significantly reduces the compressive strength. Papermaking sludge additions slightly increase the porosity and at the same time introduce finer pores, whose distribution is more homogeneous, thus lowering the compressive strength only slightly. At the same time, if paper sludge contains calcite, which is the case here, it contributes to an improvement in the mechanical properties of the fired clay products. With the optimal combination of paper-making sludge and sawdust (the amount of both wastes in the volume fraction is up to 30 % ) porous clay products can be achieved with almost the same mechanical properties as for the basic brick-making clay. The addition of silica stone mud that contains the swelling mineral montmorillonite requires a larger amount of water for shaping than the basic clay, and this amount of water will evaporate during the process of drying, which can introduce cracks into the green product. The use of such a stone mud is therefore limited. If silica stone mud was used in the clay-based industry anyway, great attention should be paid to the drying phase in order to avoid cracks. Granite-like polishing waste facilitates shaping, but it also makes the basic clay more sensitive to drying, especially when a larger amount is added. It also decreases the mechanical properties of the products. Both effects are not so pronounced when up to the mass fraction of 10 % of mud is added. The use of different wastes in the clay-based industry can have, in some cases, a positive impact on the final properties of clay products. In other cases it can be used as a substitute for basic raw materials, which can contribute to a significant saving in natural resources, and at the same time to a reduction in the amount of landfill Acknowledgements The research work described in this paper was supported by the companies Termit d.d. and Wiener- berger Opekarna Ormo`. Part of the work is the result of the European project POLISHCOVERINGS – De- velopment of an efficient and environmentally friendly polishing process for floor and wall coverings (Project number: CRAFT-1999-70904) 5 REFERENCES 1 K. Junge, Ziegel Industrie, 1 (1994), 35–38 2 D. Hauck, E. Jung, Ziegel Industrie Jahrbuch, 1991, 108–121 3 K. Junge, N. Pauls, Ziegel Industrie Jahrbuch, 1994, 90–96 4 V. Ducman, T. Kopar, Sawdust, Ind. ceram., 21 (2001) 2, 81–86 5 E. Rimpel, T. Schmedders, Ziegel Industrie Jahrbuch, 1996, 174–206 6 U. Hahn, Ziegel Industrie, 9 (1989), 458–464 7 M. Rickli, U. Eggenberger, T. Peters, Ch. Meyer, Th. Mumenthaler, Ziegel Industrie, 12 (1998), 818–827 8 F. Andreola, L. Barbieri, A. Corradi, I. Lancellotti, American Cer. Soc. Bulletin, 3,83 (2004), 9401–9408 9 V. Ducman, T. Kopar, Ind. ceram., 24 (2004) 1, 8–12 10 V. Ducman, T. Kopar, E. Sanchez, Ind. ceram., 25 (2005) 3, 164–169 11 M. Elwan, E. A. El-Alfi, H. El Didamony, Ind. Ceram., 21 (2001) 2, 87–90 V. DUCMAN, T. KOPAR: THE INFLUENCE OF DIFFERENT WASTE ADDITIONS TO CLAY-PRODUCT MIXTURES Materiali in tehnologije / Materials and technology 41 (2007) 6, 289–293 293 R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM DENTAL ALLOY JOINTS ELEKTROKEMIJSKE IN MEHANSKE LASTNOSTI RAZLI^NIH SPOJEV STELITNE DENTALNE ZLITINE Rok Zupan~i~1, Andra` Legat2, Nenad Funduk1 1Department of Prosthodontics, University of Ljubljana, Faculty of Medicine, Division of Dental Medicine, Ljubljana, Slovenia 2Slovenian National Building and Civil Engineering Institute, Ljubljana, Slovenia rok.zupancicmf.uni-lj.si Prejem rokopisa – received: 2007-09-19; sprejem za objavo – accepted for publication: 2007-10-18 In dentistry cobalt-chromium alloys are frequently used for partial denture frameworks. For fabrication of some complex frameworks, separate metal segments have to be joined. The longevity of these restorations is limited due to the mechanical or corrosive failure of the joints. The purpose of this study was to determine which joining method offers the best properties to cobalt-chromium alloy frameworks. Intact specimens, brazed and two types of laser-welded joints were compared for their electrochemical and mechanical characteristics. Electrochemical impedance spectroscopy and potentiodynamic polarization in two artificial saliva solutions were used to assess basic corrosion parameters and tensile strength of brazed and laser welded specimens was measured. The fracture surfaces and corrosion defects were examined in a scanning electron microscope. The average tensile strength of brazed joints was significantly greater than the tensile strength of both types of laser-welded joints. When laser welding was used, successful joining was limited to the peripheral aspects of the weld. The welding technique did not affect significantly the joint tensile strength. Electrochemical measurements indicated that primarily due to differences in pasivation ability, the corrosion resistance of the laser-welded joints was better than that of the brazed. Key words: brazing, laser welding, dental alloys, corrosion, strength V zobni protetiki pogosto izdelamo ogrodja fiksnih in snemnih proteti~nih izdelkov iz stelitnih (kobalt-kromovih) zlitin. Pri kompleksnej{ih konstrukcijah je treba spojiti posamezne dele ogrodja. Trajnost proteti~nega izdelka pogosto omejujejo mehanske in korozijske po{kodbe spojev. Namen te raziskave je bil ugotoviti, s katerim na~inom spajanja dobimo korozijsko in mehansko najbolj odporne spoje. Primerjali smo lotanje in dva na~ina laserskega varjenja. Osnovne korozijske parametre smo ugotovili z elektrokemijsko potenciodinamsko polarizacijo in elektrokemijsko impedan~no spektroskopijo v dveh razli~nih raztopinah umetne sline. Pri lotanih in lasersko varjenih vzorcih smo izmerili natezno trdnost. Lomne ploskve in korozijske po{kodbe smo pregledali z vrsti~nim elektronskim mikroskopom. Povpre~na natezna trdnost lotanih spojev je bila zna~ilno vi~ja od trdnosti obeh skupin lasersko varjenih spojev. Z laserskim varjenjem smo uspe{no spojili le povr{insko plast vzorcev. Na~in laserskega varjenja ni zna~ilno vplival na natezno trdnost spojev. Elektrokemijske lastnosti spojev ka`ejo, da so, predvsem zaradi razlike v sposobnosti pasivacije, lasersko varjeni spoji korozijsko obstojnej{i od lotanih. Klju~ne besede: lotanje, lasersko varjenje, dentalne zlitine, korozija, trdnost 1 INTRODUCTION Cobalt-chromium (Co-Cr) alloys are frequently used for fixed and removable partial denture frameworks.1–3 In the fabrication of some complex frameworks and in repairs or additions, separate metal segments of the framework have to be joined.2 The longevity of these restorations is limited due to the mechanical or corrosive failure of the joints.4–6 Several joining techniques are available but brazing and laser welding are most commonly used. Brazing is a process in which a molten filler metal wets and fills the gap between the parent metal surfaces.3 The filler metal has a lower melting point than the parent metal.3 In welding, the parent metals fuse and form the joint with or without a filler alloy.2 Conventional heat sources tend to produce excessive thermal damage to prosthodontic restorations and are therefore not used in dentistry.2 Laser welding has recently gained popularity, mostly because it is simpler and less time consuming than brazing.7 For brazing Co-Cr alloys, a Co-Cr alloy with a lower melting point or a noble alloy serves as a filler.5–8 Gold-based filler alloys are often used because their melting points are well below those of Co-Cr alloys.5–8 Research has shown that these joints have poor mechanical properties and corrosion resistance5,7,9 and exhibit a significantly lower tensile strength than laser welded joints.7,9 Angelini et al 5 compared the corrosion resistance of Co-Cr alloy brazing done with a gold-based filler and a Co-Cr filler and concluded that Co-Cr filler is more appropriate than gold-based filler. Co-Cr dental alloys have an excellent corrosion resistance, which is provided by a thin adherent layer of chromium-based oxides on the surface.5,6,10 Considerable defects caused by corrosion tend to appear primarily in the joints. Beside specific problems related to crevice forming, joining of 2 metals with different corrosion potentials can form a galvanic cell. 11 In such a case, the less noble metal acts as a main anode, and it could exhibit a relatively high dissolution rate.12,13 Such problems with brazed joints, caused by the dissimilar composition of the filler and the parent metals are reported.14–16 Corrosion not only results in a poor esthetic outcome, but can also compromise physical properties Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 295 UDK 669'25'26:611.314 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)295(2007) and induce biological irritation in form of an allergic reaction, lichen planus, or some other soft tissue inflammation.5,12,17,18 Corrosion properties are commonly assessed by various types of electrochemical measurements.13,19 Electrochemical potentiodynamic polarization (EPP) techniques are often used, providing general information about the corrosion resistance and susceptibility, such as the general corrosion rate, the range of passivation, and the break-down potential.12,20 These results should be carefully considered since with EPP the information is not obtained in stationary conditions.20 In rather passive systems with relatively low corrosion rates (application of noble metals or systems with protective coatings) more reliable information can usually be gained from electrochemical impedance spectroscopy (EIS) measure- ments.12,20 EIS applies sinusoidal voltage signal of relatively small amplitudes (usually a few tens of mV) and the conditions of the electrodes are only slightly disturbed.12,21 Beside the general corrosion properties of an investigated system, specific information about underlying electrochemical mechanisms can also be obtained from the measured impedance spectra.20,21 These electrochemical methods have been success- fully implemented in several investigations of various corrosion problems in dentistry.21–25 Electrolyte solu- tions, such as artificial saliva, are often used as the corrosion medium because their electrochemical properties are similar to those of the natural saliva.26 The level of corrosion resistance of brazed and laser-welded Co-Cr dental alloy joints has so far been determined merely with microscopic assessment of corrosion defects. Electrochemical properties of different joints have not yet been compared quantitatively. Most studies investigating the mechanical resistance of various joints showed that laser welds have higher tensile strength than soldered or brazed joints using noble solders.7,9,26,27 However, laser welds are more prone to fatigue damage than brazed joints.9 Short laser pulses which heat metal beyond the melting point are used for laser welding. The amount of energy released in each laser pulse is controlled by setting the welding parameters (voltage, pulse duration, and focus diameter).4 Spots where laser pulses are applied cool rapidly and the welding depth is sometimes relatively shallow in comparison with the diameter of the welded object.28,29 More powerful laser pulses not only deepen the weld penetration, but also increase poro- sity.2,28,29 In an attempt to overcome this problem, different joint designs have been proposed.2 If adjacent joint-forming surfaces are ground so as to form the shape of the letter X, they can be laser-welded starting from the center and the joint is built towards the surface of the object.2 In the process, metal is added to the joint by a filler wire with a composition equal, or very similar to the parent metal. The aim of this study was to determine which method of joining Co-Cr alloy framework segments produces joints with the best strength and corrosion resistance. The joining methods used were brazing with a Co-Cr filler and laser welding with 2 different joint designs. 2 MATERIAL AND METHODS 52 cylindrical specimens with a diameter of 2 mm and the length of 35 mm were cast in a Co-Cr alloy (Remanium GM 380; Dentaurum), following the proce- dures recommended by the alloy manufacturer. The alloy composition is shown in Table I. 16 specimens were selected for electrochemical measurement, and divided into 4 groups of 4 specimens each. In the intact group, the specimens were left as cast. The specimens of the remaining 3 groups were sectioned at the center, perpen- dicular to their long-axis, using a 0.6 mm separating disk and subsequently re-joined by brazing (brazing group) and laser welding, using an X- or I-shaped joint design (X laser and I laser groups, respectively). 36 specimens were selected for tensile strength testing, and divided into 3 groups in which specimen pairs (n = 6) were to be joined by brazing or laser welding, using an X- or I-shaped joint design. To achieve a standard gap for the brazing group, a 0.3 mm metal shim was placed between the two halves of specimens. The shim was removed once the speci- mens were positioned and invested in the phosphate- bonded investment. Flux was applied, and a Co-Cr solder was used as the filler metal. Its composition is shown in Table I. The assembly was pre-heated to 750 °C and torch brazed at 1180 °C, following the proce- dures recommended by the manufacturer of the solder. For laser welding the I-shaped joint design, the joint surfaces of specimens were straight and placed in tight contact with each other. An Nd:YAG laser with the wavelength of 1064 nm was used with the following settings: voltage of 290 V, pulse duration of 10 ms, and weld spot diameter of 0.7 mm. A weld spot overlap of approximately 75 % was used, so that the joint was R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... 296 Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 Table I: Composition (w/%) of Remanium GM 380 alloy, Co-Cr solder, and filler wire Tabela I: Sestava zlitine Remanium GM 380, Co-Cr lota in varilne `ice v masnih dele`ih (w/%) Co Cr Mo Mn Ni C Fe Si N B Nb GM 380 alloy 64.6 29 4.5 <1 – <1 – <1 <1 – – Co-Cr solder 61 28.5 3.5 – – <1 1.5 4 – 1 – filler wire – 22.1 9.1 – 63.8 – 1 – – – 3 formed with 25 pulses per specimen. These laser con- ditions were chosen to simulate typical laboratory procedures for welding of Co-Cr dental alloys.2,7,27 With the X-shaped laser welding design the joint surfaces were ground to form the shape of the letter X if viewed from the side. During grinding, the surfaces were cooled with 75 % ethanol. The tip was rounded with a hand instrument to facilitate the alignment of the halves. Laser welding was performed using pulses of lower energy (settings: 255 V, 4.5 ms, 0.9 mm) because no deep weld penetration was required for this joint design. A Co-Cr filler wire was used to complete the joints. The composition of the wire is shown in Table I. Specimens were polished using conventional laboratory procedures2 for removable partial denture frameworks with silicone polishers and polishing paste for Co-Cr alloys. The final joint diameter varied slightly due to the custom finish- ing. The corrosion parameters were assessed by electrochemical potentiodynamic polarization (EPP) and electrochemical impedance spectroscopy (EIS). Since it was expected that the investigated electrochemical systems would exhibit relatively low corrosion activities, the polarization potential in EPP measurements were changed in a wide interval from cathodic to anodic region (–0.5 V vs Ecorr to +2 V vs Ecorr). From measured potentiodynamic curves, the corrosion current densities and break-down potentials were determined. In EIS measurements, the sinusoidal voltage signal with an amplitude of 10 mV in the frequency interval between 10 mHz and 5 kHz was applied. Significant parameters, as the total impedance (IZ0I) and the polarization charge-transfer resistance (RCT), were estimated from the measured spectra. The working electrode was always the specimen with a saturated calomel electrode serving as a reference electrode and a graphite electrode as a counter electrode. Fusayama type artificial saliva 30 was used as the corrosion medium. Its pH value is 4.65 and it has the following composition: NaCl 0.4 g/L, KCl 0.4 g/L, CaCl2 ·2H2O 0.795 g/L, NaH2PO4 0.69 g/L and urea 1 g/L. A potentiostat (PC3/750; Gamry Instruments Inc) with the appropriate software (CMS 100 and CMS 300) was used for the experiment. The joined specimens were tested in tension. The total length of the specimens was 7 cm, which allowed for secure fixation in the testing machine (Z 030; Zwick GmbH & Co). The joint diameter was measured with each specimen using an electronic caliper having an accuracy of 0.01 mm. The determined value was used to compute the tensile strength of the joints by the formula = F/πr2, where is the tensile strength, F is the load at fracture and r is the half-diameter of the joint.3 Tensile tests were performed under constant extension rate of 0.008 s–1 (relative extension) according to the standard EN 10002-1.31 A 30kN load cell was used (KAP-TC, class 0.05; Zwick GmbH & Co) with the software (testXpert V10.11) provided by the manufacturer. The data were statistically analyzed using a 1-way analysis of variance (ANOVA) and Scheffé post hoc tests. Differences between groups were regarded as significant at = 0.05. After mechanical and electro- chemical testing the fracture surfaces and corrosion defects were analyzed by scanning electron microscopy (JSM–5500; JEOL), whereas the longitudinal sections were studied under a metallographic optical microscope (Neophot 22, Carl Zeiss AG). 3 RESULTS Potentiodynamic curves obtained for intact alloy and different joints are shown in Figure 1. The curves of intact alloy and laser-welded joints are similar as all of them contain rather steep and nearly straight sections up to approximately 800 mV versus saturated calomel electrode. Very low dependence of electrochemical current on the applied potential in this region indicates a nearly passive state. Beyond these potentials, the so-called break-down potentials, there are distinct transitions to active corrosion, where measured currents start to increase rapidly. The potentiodynamic curve corresponding to the brazed joint shows that the measured current increased continuously and no distinct passive region was observed. The differences in break-down potentials between groups were statistically analyzed and found to be significant (P = 0.036), with the brazed joint having much lower average break-down potential (544 mV) than other specimens (pooled R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 297 Table II: Mean values and standard deviations (SD) of corrosion potentials (Ecorr), break-down potentials (Ebd), corrosion current densities (Icor), total impedances at lowest frequency (IZ0I) and estimated polarization charge-transfer resistances (RCT) Tabela II: Povpre~ne vrednosti in standardne deviacije (SD) korozijskih potencialov (Ecorr), poru{itvenih potencialov (Ebd), gostote korozijskega toka (Icor), impedance pri najni`ji frekvenci (IZ0I) in ocene polarizacijske upornosti (RCT) mean Ecorr (mV) SD Ecorr (mV) mean Ebd (mV) SD Ebd (mV) mean Icorr (µA/cm2) SD Icorr µA/cm2 mean IZ0I (kΩ) SD IZ0I (kΩ) mean RCT (kΩ) SD RCT (kΩ) intact −462.3 111.7 758.3 73.6 0.10 0.05 182.9 38.0 75.1 10.8 brazing −409.9 32.0 543.8 298.0 0.07 0.04 80.9 37.3 7.1 5.1 I laser −560.2 9.7 875.5 32.4 0.22 0.14 140.1 30.3 21.5 4.0 X laser −555.0 29.5 885.0 84.4 0.16 0.06 173.6 16.2 29.5 6.0 average 840 mV). All corrosion current densities were very low (from 10–7 to 10–6 A/cm2). Exact values of corrosion current densities were calculated by the extrapolation of the Tafel slopes19 however, they are generally represented as the intersections of anodic and cathodic curves (Figure 1). Estimated values of the corrosion potentials, corrosion current density, and break-down potentials are shown in Table II. The electrochemical impedance spectra generally confirmed the observation drawn from the electro- chemical potentiodynamic measurements. Impedance spectra are presented as Nyquist plots in Figure 2. Total impedance at the low frequency range is generally related to the corrosion resistance. The impedances at the lowest frequency IZ0I and estimated values of the polarization charge-transfer resistance RCT are shown in Table II. The results obtained from EPP measurements and those from EIS measurements were in agreement. Both of them clearly expressed higher corrosion resistance of the laser-welded joints compared to the corrosion resistance of the brazed joints. SEM analysis of the specimens confirmed the results of the electrochemical measurements. The most pronounced corrosion damage was at the brazed joints, located primarily at the parent metal. Corrosion damage of the intact alloy and the laser-welded joints was minor and mostly located at certain defects on surfaces, such as inclusions or scratches. The mean value (standard deviation) for the tensile strength of brazed joints was 792 (238.5) MPa. This was significantly (P = 0.004) greater than the tensile strength of both types of laser-welded joints. The mean values for I-shaped and X-shaped joint designs were 404 (76.7) MPa and 405 (120.4) MPa, respectively. These are shown in Figure 3. All specimens, regardless of the joining technique, fractured in the joints. The strength of the laser-welded joints did not depend significantly on the joint design used. SEM examination, however, revealed differences in the effective cross-sections joined. The fracture surfaces of brazed joints were relatively smooth and they exhibited a fine grained partially ductile nature (Figure 4). Fracture surfaces of the I-shaped laser-welded joints showed that only peripheral aspects of these specimens were successfully joined, since under the surface there were some voids R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... 298 Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 Figure 4: Fracture surface of brazed joint at X 400 magnification Slika 4: Lomna ploskev lotanega vzorca pri 400-kratni pove~avi Figure 2: Nyquist plots for different joints and intact alloy Slika 2: Nyquistovi diagrami za razli~ne spoje in intaktno zlitino Figure 3: Means and standard deviations of tensile strength of the different joint types Slika 3: Povpre~ne vrednosti in standardne deviacije natezne trdnosti razli~nih spojev Figure 1: Potentiodynamic curves for different joints and intact alloy Slika 1: Potenciodinamske krivulje za razli~ne spoje in intaktno zlitino and the central area remained unwelded. Even after X-shaped laser welding, the central area remained partly unwelded, but the joined effective cross-section was larger than in the I-shaped laser-welded specimens. However, there was no significant difference in the tensile strength of the two groups. In the laser welded joints the fracture surfaces were coarse grained and brittle (Figure 5). 4 DISCUSSION The average tensile strength of the laser-welded joints was significantly lower than that of the brazed joints, mainly due to the smaller cross-section of the welded joints and partly due to the relatively strong brazed joints. For the purpose of comparing brazing with laser welding of Co-Cr alloys, most authors used a noble filler metal.5–7 Average tensile strengths of these bra- zings ranged from 357 MPa to 519 MPa.7,9 In this study, the average tensile strength of brazed specimens was 792 MPa. The average tensile strength of laser-welded Co-Cr dental alloy joints investigated thus far has ranged from 480 MPa to 751 MPa,7,28, 29 which exceeds the average strength of laser welds in this study (404 MPa and 405 MPa). An important reason for the relative weakness of laser-welded joints in this study is a small effective cross-section of specimens that was actually joined. This is a problem associated with low weld penetration depth.28,29 In laser welding using the I-shaped joint design, laser pulses were not powerful enough to reach the central parts of the specimens, although peripherally the metal was overheated with a resulting porosity. Optimizing laser parameters could, to some extent, improve the quality of these welds.29 The composition of the alloy greatly affects its weldability and in this respect, the carbon content is critical.28 For laser welding, the manufacturer recommends an alloy with no carbon (Remanium 900). Yet, according to clinical experience, this alloy is not stiff enough for partial denture frameworks and was therefore not used in this study. Nevertheless, laser welding of Co-Cr denture frameworks has been used with considerable success in clinical practice. A plausible explanation is that during mastication prostheses are never subjected to isolated tensile loads. In bending, most of the load is placed on the peripheral parts of the framework; one side is subjected to tension and the other one to compression, while central portions are less involved. Since the examination of fracture surfaces in this study showed that the laser welding technique was much more effective in the peripheral than in the central parts of the specimens, it is possible that the technique would prove satisfactory in most clinical situations. The results of electrochemical measurements, which are in agreement with previously published data,6,12,17,18 showed excellent corrosion resistance of the intact Co-Cr dental alloy. On the basis of previous qualitative observations, it was expected that laser-welded joints would exhibit better resistance to corrosion than brazed joints. This was confirmed by the results of both electro- chemical methods used. With the potentiodynamic curves of the welded joints, a characteristic passive region was observed, whereas in the case of the brazed joints, measured current increased rapidly and continuously. The more noble potential of the filler additionally polarized the parent metal and consequently almost no passivation took place. Similar conclusions could be drawn from the measured impedance spectra, where all characteristic impedance values of the brazed joints were lower than those of the intact alloy and the welded joints. However, the corrosion processes, which were investigated by EPP and EIS, occurred at the surface of the test specimen and thus were not influenced by bulk defects. Compared to brazed joints, laser-welded joints had more porosity and similar defects likely to initiate corrosion in a clinical situation. From EPP and EIS measurements it was not possible to determine weather the corrosion process was localized or generalized, which is a limitation of this study. SEM analysis revealed localized phenomena, such as pitting near the joints and pronounced corrosion in some defects on surfaces. In case of localized corrosion, and over longer periods of time, the process could become autocatalytic and more pronounced than in the study. The laser welding process could, to some extent, be improved by increasing the weld penetration depth. Preparation of the areas to be welded, such as marking them black with a black felt-tipped pen or airborne- particle-abrading can reduce laser beam reflection and could probably improve the welding efficiency.2 In situations where surfaces to be laser-welded can be R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 299 Figure 5: Fracture surface of laser welded joint at X 400 magnifica- tion Slika 5: Lomna ploskev lasersko varjenega vzorca pri 400-kratni pove~avi designed in advance, large joint surfaces might ensure sufficient strength and limited thickness might enable complete joining with minimum porosity. These questions remain to be addressed in the future. 5 CONCLUSION Laser-welded Co-Cr alloy joints exhibit excellent corrosion resistance, but their tensile strength is limited due to the shallow weld penetration. Using laser welding mainly the peripheral parts of the joints can be successfully welded. Brazed joints are less resistant to corrosion but have significantly higher tensile strength. 6 REFERENCES 1 Wataha JC. Alloys for prosthodontic restorations. J Prosthet Dent 87 (2002), 351–63 2 Wulfes H. Precision milling and partial denture constructions. Bremen: Academia Dental; 2003. p. 259–60; 115–9; 108–13 3 Craig RG. Dental materials: properties and manipulation. 7th ed. St Louis. Mosby; 2000. p. 228–38; 20–3 4 Henriques GE, Consani S, Rollo JM, Andrade e Silva F. Soldering and remelting influence on fatigue strength of cobalt-chromium alloys. J Prosthet Dent. 78 (1997), 146–52 5 Angelini E, Pezzoli M, Rosalbino F, Zucchi F. Influence of corrosion on brazed joints’ strength. J Dent 19 (1991), 56–61 6 Luthy H, Marinello CP, Reclaru L, Sharer P. Corrosion conside- rations in the brazing repair of cobalt based partial dentures. J Prosthet Dent 75 (1996), 515–24 7 NaBadalung DP, Nicholls JI. Laser welding of a cobalt-chromium removable partial denture alloy. J Prosthet Dent 79 (1998), 285–90 8 Dominici JT, Sobczak KP, Mitchell RJ. A comparison of infrared- and torch-soldering of Au-Pd and Co-Cr metal-ceramic alloys using a high-fusing solder. J Prosthodont 4 (1995), 101–10 9 Wiskott HW, Macheret F, Bussy F, Besler UC. Mechanical and elemental characterization of solder joints and welds using a gold-palladium alloy. J Prosthet Dent 77 (1997), 607–16 10 Dong H, Nagamatsu Y, Chen KK, Tajima K, Kakigawa H, Shi S. Corrosion behavior of dental alloys in various types of electrolyzed water. Dent Mater J 22 (2003), 482–93 11 Landolt D. Introduction to surface reactions: Electrochemical basis of corrosion in Corrosion mechanics in theory and practice. Vol 17. 2nd ed. New York: Marcel Dekker Inc; 2002. p 1–19 12 Kelly RG, Shoesmith DW, Buchheit RG. Electrochemical techniques in corrosion science and engineering. New York: Marcel Dekker Inc; 2003. p. 55–124; 125–150; 205–57 13 McDonald MM. Corrosion of braze joints. In: Metals Handbook 9th Edition, Vol 13, Corrosion, Metals Park: ASM International, 1987. p. 876–886 14 Shigeto N, Yanagihara T, Hamada T, Budtz-Jorgensen E. Corrosion properties of soldered joints. Part 1. Electrochemical action of dental solder and dental nickel-chromuim alloy. J Prosthet Dent 62 (1989), 512–515 15 Cortada M, Giner LL, Costa S, Gil FJ, Rodriguez D, Planell JA. Galvanic corrosion behavior of titanium implants coupled to dental alloys. J Mater Sci Mater Med 11 (2000), 287–93 16 Al-Ali S, Oshida Y, Andres CJ, Barco MT, Brown DT, Hovijitra S, et al. Effects of coupling methods on galvanic corrosion behavior of commercially pure titanium with dental precious alloys. Bio-med Mat Eng 15 (2005), 307–16 17 Wataha JC. Biocompatibility of dental casting alloys. J Prosthet Dent 87 (2000), 223–34 18 Geurtsen W. Biocompatibility of dental casting alloys. Crit Rev Oral Biol Med 13 (2002), 71–84 19 Martin MD, Broughton S, Drangsholt M. Oral lichen planus and dental materials: a case-control study. Contact Dermatitis 48 (2003), 331–6 20 Mansfeld F, Shih H, Greene H, Tsai CH. Analysis of EIS data for common corrosion processes. In: Electrochemical Impedance: Analysis and Interpretation, Scully JR, Silverman DC, Kendig MW, editors. Philadelphia: American Society for Testing and Materials; 1993. p. 37–53 21 Sun D, Monaghan P, Brantley WA, Johnston WM. Electrochemical impedance spectroscopy study of high-palladium dental alloys. Part I: Behavior at open-circuit potential. J Mater Sci Mater Med 13 (2002), 435–42 22 Cai Z, Vermilyea SG, Brantley WA. In vitro corrosion resistance of high-palladium dental casting alloys. Dent Mater 15 (1999), 202–10 23 Mueller HJ, Hirthe RW. Electrochemical characterization and immersion corrosion of a consolidated silver dental biomaterial. Biomaterials 22 (2001), 2635–46 24 Mueller HJ. In vitro tarnish and corrosion of a consolidated silver material for direct filling applications. Dent Mater 17 (2001), 60–70 25 Acciari HA, Guastaldi AC, Brett CMA. Corrosion of the component phases presents in high copper dental amalgams. Application of electrochemical impedance spectroscopy and electrochemical noise analysis. Corr Sci 47 (2005), 547–653 26 Duffo GS Quezada Castillo E. Development of an artificial saliva solution for studying the corrosion behavior of dental alloys. Corrosion 60 (2004), 594-602 27 Jemt T, Henry P, Linden B, Naert I, Weber H, Bergstrom C. A com- parison of laser-welded titanium and conventional cast frameworks supported by implants in the partially edentulous jaw: a 3-year prospective multicenter study. Int J Prosthodont 13 (2000), 282–8 28 Bertrand C, Le Petitcorps Y, Albingre L, Dupuis V. The laser welding technique applied to the nonprecious dental alloys procedure and results. Br Dent J 190 (2001), 255–7 29 Bertrand C, le Petitcorps Y, Albingre L, Dupuis V. Optimization of operator and physical parameters for laser welding of dental materials. Br Dent J 196 (2004), 413–8 30 Fusayama T, Katayori T, Nomoto S. Corrosion of gold and amalgam placed in contact with each other. J Dent Res 42 (1963), 1183–97 31 European Committee for Standardization. Standard EN 10002 – 1: Metallic materials; tensile testing; part 1: method of test. Brussels, 1996 R. ZUPAN^I^ ET AL.: ELECTROCHEMICAL AND MECHANICAL PROPERTIES OF COBALT-CHROMIUM ... 300 Materiali in tehnologije / Materials and technology 41 (2007) 6, 295–300 T. KVACKAJ ET AL.: DEVELOPMENT OF MICROSTRUCTURE OF STEEL FOR THERMAL POWER GENERATION DEVELOPMENT OF MICROSTRUCTURE OF STEEL FOR THERMAL POWER GENERATION RAZVOJ MIKROSTRUKTURE JEKEL ZA TERMI^NO GENERACIJO ENERGIJE Kvackaj Tibor, Kuskulic T., Fujda M., Pokorny I., Weiss M.1, Bevilaqua T. 1 Technical University in Kosice, Slovakia 1@P-Podbrezová, a. s. tibor.kvackajtuke.sk Prejem rokopisa – received: 2006-10-05; sprejem za objavo – accepted for publication: 2007-10-18 The evolution of microstructure during the reheating and cooling of steel for thermal power generation was investigated. On the basis of the microstructure produced during cooling a CCT diagram is proposed. Key words: steel, thermal power generation, microstructure, CCT diagram Raziskan je bil razvoj mikrostrukture pri segrevanju in ohlajanju jekel za toplotno generacijo energije. Na podlagi mikrostrukture, ki je nastala pri ohlajanju, je bil predlo`en CCT-diagram. Klju~ne besede: jeklo, toplotna generacija energije, mikrostrukture, CCT diagram 1 INTRODUCTION Natural gas is now an important source of energy. However, especially in developing countries, lignite and anthracite will also be used for the generation of elec- trical energy, and power stations utilizing these energy resources for steam production will also play an important role in the future. The minimization of the capital costs required to build new steam-power stations leads to the extensive use of ferrite-martensite steels, or martensite steels for all the main components in boilers and turbines 1. For modern power stations the operating conditions are defined as follows: • 540 oC/18 MPa • 600 oC /30 MPa • 720 oC /37.5 MPa - expected for the future The applied materials were analyzed for their creep rupture strength under the following conditions: • short-term creep behavior up to t = 10 000 h • long-term creep test up to t = 100 000 h In Figure 12 the chemical analysis and the creep resistance is given for some classic and newly developed steels. The 1CrMoV steel has been used for a very long time and is considered as a classic steel. The steels with an increased content of Cr up to 9–12 % with addition of Nb and B are considered as newly developed types. The newly developed steels can be, in comparison to classic materials, applied in operating conditions with increased steam temperatures of 30-70 °C. The steels for Materiali in tehnologije / Materials and technology 41 (2007) 6, 301–303 301 Table 13: Chemical composition and properties of two steels Tabela 13: Kemi~na sestava in lastnosti dveh jekel w C Cr Mo W Ni V Nb N B/µg/g Rp 0,2/ MPa FATT50/ °C GX12CrMoVNbN9 1 0,12 9 1,0 – 0,4 0,2 0,06 0,05 80 633 +46 GX12CrMoWVNbN 10 1 1 0,12 10 1,0 1,0 0,7 0,2 0,06 0,05 10 730 +40 Figure 12: Creep rupture strength of the new 9–10 % Cr rotor materials applied in Europe Slika 1: Odpornost proti lezenju pri novih 9–10 % Cr jeklih, ki se uporabljajo v Evropi UDK 669.14:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 41(6)301(2007) steam-turbine rotors can be divided into three groups according to their chemical composition: • 10 % Cr + Mo • 10 % Cr + W+ Mo • 9 % Cr + Mo + B These steels are used for the manufacturing of rotors with diameters up to 1200 mm. The basic properties for two steel grades are given in Table 1 3. Besides the steels based on CrMoV or CrMoVW, steels based on CrMoNiV or CrMoNiVW have also been developed. The research work is currently oriented to improve steel technology with the aim to decrease the content of unwanted elements in the steel and to master the forging technology and heat-treating processes of these high-purity steels. The aim is to obtain good chemical and structural homogenity of the material, as well as forgings with minimum of defects 4. The properties of steel depend on the steam characte- ristics, since the turbines can operate in conditions of: • HP – high pressure • IP – intermediate pressure • LP – low pressure The specifications of the material should be defined by the following characteristics: 1. Static strength – failure strength 2. Creep rupture strength, high-temperature strength 3. Toughness – fracture toughness 4. Fatigue properties – low-cycle fatigue – high-cycle fatigue 5. Crack growth rate – static-creep (CG) – alternative – fatigue 6. Corrosion resistance – local corrosion – corrosion under pressure – corrosion fatigue 7. Erosion resistance 2 MATERIAL AND EXPERIMENTS For the experiments the CrNiMoV steel, which is the equivalent to STN 41 6537, with the chemical composition in Table 2, was used. The experiments were aimed at: – the evaluation of the influence of temperature and time on the austenite grain size. – the influence of cooling rate on the formation of the microstructure and the ARA diagram. The experimental methods were: – light microscopy – differential dilatometry 3 RESULTS AND DISCUSSION The diagrams showing the influence of reheating temperature and reheating time on the austenite grain size change are in Figure 2 and Figure 3. From the de- pendencies in these two figures it was concluded that: – the holding time at 1000 °C does not influence the austenite grain size. Thus it is possible to classify this temperature as "low-sensitive" to austenite grain size change. – the reheating temperature of 1050 °C had a great influence on the austenite grain size, while, for holding times of 15 min and 30 min the difference in the austenite grain size is on average 8 µm, but for holding times of 45 min and 60 min this difference increases on average to 45,5 µm. It is possible to classify this temperature as sensitive to austenite grain size change. – for reheating temperatures of (1100, 1150, 1200) °C, there is a slow increase in the austenite grain size for all the reheating times. It is possible to classify this T. KVACKAJ ET AL.: DEVELOPMENT OF MICROSTRUCTURE OF STEEL FOR THERMAL POWER GENERATION 302 Materiali in tehnologije / Materials and technology 41 (2007) 6, 301–303 Table 2: Chemical analysis of the experimental steel Tabela 2: Kemi~na sestava jekla za raziskavo C Mn Si P S Cr Ni Cu Mo V Al As Sn Sb Ca H N O w/% w/(µg/g) 0.29 0.04 <0.01 0.003 0.003 1.57 2.84 0.010 0.39 0.11 0.004 11 8 <5 20 0.5 44 25 0 20 40 60 80 100 120 140 160 0 10 20 30 40 50 60 70 Time, t /min D ia m e te r o f A G S : D γ /µ m 1 1 0 0 5 0 0 0 ° ° C C 1100 °C 1150 °C 1200 °C Figure 3: Dependence of AGS on reheating time Slika 3: Odvisnost AGS od trajanja segrevanja 0 20 40 60 80 100 120 140 160 950 1000 1050 1100 1150 1200 1250 T /°C D γ / µ m 3 15 min 0 min 45 min 60 min Figure 2: Dependence of AGS on reheating time Slika 2: Odvisnost velikosti AGS avstenitnih zrn od temperature segrevanja temperature interval as low-sensitive to austenite grain size change. At the reheating temperature of 1050 °C a faster growth of austenite grain size is observed by increasing the reheating time. From 30 min to 45 min there is a loss of the hindering effect of carbide and nitride particles on the migration of austenite grain boundaries. The solubility of the VC and VN precipitates for given contents of vanadium and carbon is calculated from Equations (1) and (2): lg (w(V)4/3 · w(C)) = 7,06 – 10 800/T (1) lg (w(V) · w(N)) = 3,02 – 7 840/T (2) The solubility is reached at 940 °C for vanadium carbide and 970 °C for vanadium nitride. It is clear that the precipitates of vanadium carbide and nitride do not hinder the growth of austenite grains during the reheating time at 1050 °C. The influence of cooling rate on the formation of the final microstructures can be evaluated by considering the dilatometry curves, the hardness and the microstructures after different cooling rates from a constant reheating temperature. In this way the transformation CCT diagram can be obtained Figure 4. The obtained diagram shows that in the examined steel a real-time ferritic transformation of austenite does not occur and that only the formation of martensite, martensite + bainite, or bainite takes place. The investigated steel is thus a high-through-hardening steel. 4 CONCLUSION From the results of the experimental work on a CrNiMoV steel it is possible to draw the following conclusions: – for the reheating temperature T = 1000 °C, the effect of reheating time is negligible and the austenite grain size remains in the interval  = 49–57 µm. This reheating temperature ensures the fine-grained austenitic structure, which is a good starting point for achieving a fine-grained secondary structure, also in cases without plastic deformation after reheating. This is also valid for 15 min or 30 min reheating at 1050 °C when the diameter  = 52–61 µm is obtained. – for the reheating temperature of 1050 °C and reheating time t = 45 min and 60 min an increased sensitivity to austenite grain growth is found and the size of  = 103–109 µm is obtained. Similar characteristics are also observed for the reheating conditions T = 1100–1200 °C and t = 15 min or 60 min when the size of  = 106–140 µm is obtained. If the reheating conditions are chosen considering the mentioned intervals, a controlled forging and con- trolled cooling regime will be required for achieving a fine-grained secondary structure. – the CCT diagram obtained shows that the investi- gated steel is self-hardening. Acknowledgement This research was carried out within the scope of the EUREKA E!3192 ENSTEEL project. 5 REFERENCES 1 Kvackaj, T.: Entry opponency of project EUREKA E!3192 ENSTEEL 2 Sheng, S., Kern, T. U.: High strenght cast and forged materials for application in steam turbine design, In: PARSONS 2000 Advanced materials for 21st century turbines and power plant, 3–7 July 2000, Cambridge 3 Mayer, K. H., Kern, T. U., Staubli, M., Tolksdorf, E.: Long-term investigation of specimens of 24 production components manufactured from advanced martensitic 10 % Cr steels for 600 °C steam turbines, In: PARSONS 2000 Advanced materials for 21st century turbines and power plant, 3–7 July 2000, Cambridge 4 Scarlin, R. B.: Improved materials for high efficiency steam turbines, In: Advances in turbine materials, design and manufacturing, 4-6 November 1997, Newcastle upon Tyne T. KVACKAJ ET AL.: DEVELOPMENT OF MICROSTRUCTURE OF STEEL FOR THERMAL POWER GENERATION Materiali in tehnologije / Materials and technology 41 (2007) 6, 301–303 303 Figure 4: CCT diagram Slika 4: CCT diagram MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH, PORTORO@, 8.–10. OKTOBER, 2007 YOUNG SCIENTISTS AWARDS, 15th CONFERENCE ON MATERIALS AND TECHNOLOGY, PORTORO@, 8–10 OCTOBER, 2007 MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH JURIJ GONTAREV VALJI, d. o. o., [tore, @elezarska cesta 3, 3220 [tore Po zaklju~ku dodiplomskega {tudija materialov na Fakulteti NTF Univerze v Ljubljani sem se zaposlil v livarni valjev Valji, d. o. o., [tore. Kot mladi raziskova- lec sem vklju~en v bilateralne projekte med Valji, d. o. o., in IMT v Ljubljani kot tudi v projekte, ki jih sofinancira MG in MVZT. IDENTIFIKACIJA KARBIDOV V ZLITINI PLA[^A VALJA ZA TOPLO VALJANJE JEKLENE PLO^EVINE Jurij Gontarev1, Mirko Dober{ek2 1VALJI, d. o. o., [tore, @elezarska cesta 3, 3220 [tore 2In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana Proizvodnji program livarne je ulivanje in toplotna obdelava razli~nih vrst lito`eleznih valjev. Livarna v [torah ima 100-letno tradicijo. Izdeluje valje za toplo valjanje jekla (dolgi in plo{~ati program), valje za za~etna ogrodja toplih valjarn v barvni metalurgiji (alu- minij, baker, cink), votle valje za papirno in gumarsko industrijo, valje za mineralo{ko industrijo. Izdeluje enoslojne valje s klasi~no gravitacijsko tehnologijo ter ve~slojne valje s tehnologijo centrifugalnega litja. V zadnjem ~asu je poudarek na razvoju srednje in mo~no legiranih litin z indefinitno izlo~enim grafitom. Za delovne plasti teh valjev se uporabljajo litine z nadevtek- ti~no sestavo od 0,8 % do 2,3 % C. Litine so legirane s karbidotvorci (Cr, W, Mo, V), zato se v njihovih mikro- strukturah pojavlja vrsta razli~ni enostavnih in kom- pleksnih karbidov. Ker sem v letu 2006 kot MR na konferenci o mate- rialih in tehnologijah v Portoro`u predstavil centrifu- galno litje valjev, sem letos poro~al o identifikaciji karbidov v mikrostrukturah centrifugalno lite litine z masno vsebnostjo ogljika med 0,8 % in 1,0 %. V livarni [tore izdelujemo ve~slojne centrifugalno lite valje z oznako kvalitete S-HSS. Delovna povr{ina (pla{~) valjev je izdelana iz srednje legirane nadevtek- ti~ne jeklene litine. V legiranih jeklenih litinah se pojavi vrsta karbidov tipa: M3C, M23C6, M7C3, M6C, M2C in MC, ki so po navadi {e kompleksne sestave. V labora- torijsko izdelanih zlitinah in "in vivo" pla{~a valja smo identificirali karbide. Zaradi enostavnej{e identifikacije karbidov v mikrostrukturah smo izvedli jedkalne {tudije z razli~nimi jedkali (Picral, Muracami in Groesbeck). V strjevalnih mikrostrukturah teh zlitin je prostorninski dele` do 10 % vezanega evtektika s karbidi tipa M7C3 in M6C. S selektivnim jedkanjem in mikrokemijsko analizo s SEM, EDS in WDS smo dolo~ili karbide, ki se tudi morfolo{ko razlikujejo. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 305 Slika 1: Dvojni vezani evtektik (SEI pos.) V litih mikrostrukturah zlitine S-HSS so mo~no nehomogena primarna kristalna zrna avstenita, ki se pri danih pogojih ohlajanja transformirajo v martenzit. V meddendritnih prostorih sta kristalizirala dva tipa evtektikov:  + M7C3 in  + M6C (slika 1). Z metodo SEM-EDS smo z analizo sekundarnih elektronov potrdili, da sta v mikrostrukturi dva tipa evtektikov z razli~nima karbidoma. Opa`eni drobni karbidi v matrici so kromovi z raztopljenim vanadijem in molibdenom podobne vsebnosti kot v karbidih evtek- ti~nega zloga. Menimo, da je prostorninski dele` evtek- tika, bogatega s karbidom (Fe3Mo3)C, cca 15 % vsega evtektika v mikrostrukturi. Ugotovili smo, da se silicij delno topi v karbidih, bogatih z molibdenom, medtem ko ga v kromovih karbi- dih ni. Mangan se delno topi v kromovih karbidih, medtem ko ga v karbidih, bogatih z molibdenom, nismo ugotovili. Po mejah kristalnih zrn avstenita smo na{li drobna zrna kromovih karbidov tipa M7C3 in M23C6, ki so primarnega in sekundarnega izvora. Opazili smo, da so karbidi tipa M6C bolj kompaktno vezani z matrico, medtem ko smo na mejah M7C3/ opazili razpoke. Zanimivo je, da smo v mikrostrukturah ugotvili vana- dijeve karbide V4C3 kljub nizki vsebnosti vanadija v zlitini. Ti karbidi so nalo`eni v evtekti~nih obmo~jih mikrostrukture, kar nakazuje njihovo strjevanje dale~ pod temperaturo njihovega tali{~a, o ~emer poro~ajo tudi nekateri tuji avtorji. Z jedkalno {tudijo smo ugotovili, da lahko na eno- staven na~in lo~imo nekatere vrste karbidov, kar je pomembno za karakterizacijo mikrostrukture v nepo- sredni proizvodnji (hitro in poceni), medtem ko so zahtevnej{e mikrokemijske analize s SEM-EDS, -WDS potrdile prisotnost teh vrst karbidov in identificirale dvojni vezani evtektik. CARBIDE IDENTIFICATION IN ALLOY FOR ROLLS MADE FOR HOT STRIP MILL WORK ROLL Jurij Gontarev1, Mirko Dober{ek2 1VALJI, d. o. o., [tore, @elezarska cesta 3, 3220 [tore 2Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana In foundry [tore we produce centrifugally casted multilayer rolls. Working layer (coat) of rolls is made from medium-alloyed hypereutectoid steel casting. We identified present carbides in laboratory made alloys. Solidified microstructures contain about 10 vol.% of bonded eutectic with carbides type M7C3 and M6C. With selective etching and microchemical analysis on SEM, EDS and WDS we determined present carbides that differentiate morphologically. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 306 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Kristina @agar Naslov: Mlin{e 32, 1411 Izlake Rojena 15. 8. 1981 v Trbovljah Moje raziskovalno delo se navezuje na sintezo perov- skitnih nanopal~k z metodo elektroforezne depozicije (EPD) solov v pore polikarbonatnih membran. Po oprav- ljeni depoziciji je treba vzorec kalcinirati in odstraniti nosilec (polikarbonatno membrano). Za uspe{no sintezo nanopal~k je treba pripraviti stabilne sole in optimizirati pogoje elektrodepozicije. Za karakterizacijo nanopal~k uporabljamo rentgensko difrakcijo (XRD), termi~ni analizi (DTA, TGA), vrsti~ni in presevni elektronski mikroskop (SEM in TEM). SINTEZA IN KARAKTERIZACIJA PEROVSKITNIH NANOPAL^K @agar Kristina, [turm S., ^eh Miran Institut "Jo`ef Stefan", Jamova 39, 1000 Ljubljana Z elektroforetsko depozicijo sol-gela v nano dimen- zijske pore polikarbonatne membrane smo sintizirali perovskitne nanopal~ke (BaTiO3, SrTiO3 in CaTiO3). Rast nanopal~k je potekala na delovni elektrodi iz aluminija, na katero je bila pritrjena polikarbonatna membrana. Kot protielektrodo smo uporabljali mre`asto platinasto elektrodo. Kot nosilec za nanos depozita med elektroforetsko depozicijo smo uporabljali polikarbo- natne membrane (PC) z debelino od 10 µm do 25 µm in s premerom por 200 nm. Elektroforetsko depozicijo smo izvajali 30 min pri razli~nih napetostih, in sicer pri 2 V, 5 V in 30 V. Po depoziciji je sledila kalcinacija vzorcev pri visokih temperaturah. Med toplotno obdelavo je PC- membrana zgorela, vzorec pa je kalciniral in se zgostil. Vzorce smo nato karakterizirali z naslednjimi metodami: z rentgensko pra{kovno difrakcijo (XRD), diferen~no termi~no analizo (DTA), vrsti~no (SEM) in s presevno (TEM) elektronsko mikroskopijo. Z rentgensko pra{kov- no difrakcijsko analizo smo potrdili kristalini~nost perovskitov. SEM-preiskave so potrdile, da imajo perovskitne nanopal~ke, ki rastejo v porah PC-membran enak premer po celotni dol`ini. Ravno tako smo opazili nastanek plasti perovskita na povr{ini membran, ki je nastal pod vplivom vi{je delovne napetosti. S TEM smo opazovali strukturo in premer nanopal~k, ki je bil med 50 nm in 180 nm (slika 1). Slika 2 prikazuje pove~ano regijo nanopal~ke, ki je polikristalini~na z velikostjo zrn med 25 nm in 50 nm. Primerjava med eksperimentalnim in izra~unanim difrakcijskim vzorcem je pokazala, da so zrna kubi~nega BaTiO3. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 307 Slika 1: TEM slika BaTiO3-nanopal~ke Slika 2: TEM-slika polikristalini~ne BaTiO3-nanopal~e z odgovarja- jo~im elektronskim difrakcijskim vzorcem SYNTHESIS AND CHARACTERIZATION OF PEROVSKITE NANORODS @agar Kristina, [turm S., ^eh Miran Jo`ef Stefan Institut, Jamova 39, 1000 Ljubljana In our work we present the synthesis of perovskite nanorods (BaTiO3, SrTiO3 and CaTiO3) by sol-gel electrophoretic deposition into template membranes. For that we used several processing methods: sol-gel processing, electrophoretic deposition and template- based growth. The growth of the nanorods occurred at a working electrode of aluminum, on which we attached a polycarbonate membrane. Pt mesh was used as the counter electrode. The track-etched hydrophilic polycarbonate (PC) membranes were used as template membranes, with pore diameters of 200 nm and a thickness of 10 µm. For the electrophoretic growth the potentials of 2 V, 5 V and 30 V were applied between the electrodes, and this was maintained for 30 min. After the deposition, the samples were annealed at elevated temperatures. This heating procedure was done in order to burn off the polycarbonate membrane and to make the nanorods dense and crystalline. The samples were then characterized by X-ray diffraction (XRD), differential thermal analysis (DTA), scanning (SEM) and transmission (TEM) electron microscopy. Crystalline perovskite phases were confirmed by XRD analysis. By using SEM we found that the perovskite nanorods grown in a PC membrane have a uniform diameter throughout their entire length. We also observed that higher potential leads to a layered formation of perovskite on the membrane surface after the pores are filled. By using TEM and electron diffraction analysis, we found that these nanorods are dense and polycrystalline, with the grain size ranging between 25 nm up to 50 nm. The size of nanorods is approximately 10 µm with a diameter of an individual nanorod in the range of 100 nm to 180 nm. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 308 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Bla` Brulc SINTEZA POLI(ß-BENZIL L-ASPARTATA) IN NJEGOVIH KOPOLIMEROV Z L,L-LAKTIDOM Bla` Brulc, Maja Gri~ar, Ida Poljan{ek, Ema @agar, Majda @igon Kemijski in{titut, Hajdrihova 19, SI-1001 Ljubljana, Slovenija Biorazgradljive polimere definiramo kot take, ki se razgrajujejo v naravnem okolju oziroma ki jih katabo- lizirajo mikroorganizmi (bakterije, glive...) ali pa se encimsko razgradijo in vivo. ^e so njihovi razgradni produkti okolju ne{kodljivi (v idealnem primeru ogljikov dioksid in voda) ali se celo koristno izrabijo v svojem okolju (npr. mle~na kislina, aminokisline...), pa govo- rimo o biozdru`ljivih polimerih. Omenjeni polimeri se v zadnjih 25 letih vse bolj uporabljajo kot nosilci za nadzorovano spro{~anje zdravilnih u~inkovin. To podro~je je eno od najhitreje razvijajo~ih se podro~ij biokemijske znanosti. Sistemi za kontrolirano spro{~anje zdravilnih u~inkovin imajo {tevilne prednosti pred konvencionalnimi sistemi doziranja zdravilnih u~inkovin, kot so na primer: izbolj- {ana u~inkovitost, zmanj{ana toksi~nost, zdravljenje, prijetnej{e za pacienta (manj{e {tevilo administracij zdravila, manj izraziti stranski u~inki). V sintezi poliaminokislin so se kot primerni reagenti v preteklosti izkazali t. i. N-karboksianhidridi amino- kislin (4-substituirani oksadolidin-2,5-dioni). Najpogosteje uporabljeni iniciatorji polimerizacije N-karboksianhidridov so primarni in terciarni amini ter alkoksidi. Mehanizem iniciacije je mo~no odvisen od narave uporabljenega iniciatorja. Elektrofilni mesti 2 in 5 sta po eni strani ob~utljivi za napad nukleofilnih inicia- torjev (primarni amini, alkoholi), {ibka kislost laktam- skega protona na mestu 3 pa je razlog, da poteka iniciacija z bolj bazi~nimi in nenukleofilnimi iniciatorji v smeri mehanizma "aktiviranega monomera". Sintezo benzilno za{~itenega N-karboksianhidrida L-asparaginske kisline smo izvedli z reakcijo ß–benzil L-aspartata s trifosgenom kot acilirnim reagentom. N-Karboksianhidridi so zelo ob~utljivi predvsem za vlago pa tudi na prisotne ne~isto~e, zato jih je treba ~istiti (ve~kratna kristalizacija), obvezna pa je tudi hramba v argonovi atmosferi pri nizki temperaturi. ^eprav je bil `e prekristaliziran produkt reakcije elementno ~ist, je bila za nadaljnje sintezne korake potrebna dodatna prekristalizacija, saj lahko v nasled- njem koraku z neo~i{~enim N-karboksianhidridom sintetiziramo le ni`jemolekularne homopolimere. Iz ve~krat prekristaliziranega N-karboksianhidrida L-asparaginske kisline smo pri sobni ali nekoliko zvi{ani temperaturi v N,N-dimetilformamidu z dobrim izkorist- kom (> 95 %) sintetizirali poli(ß-benzil L-aspartate) s pov- pre~jem molskih mas redov velikosti 103–104 g mol–1. Pripravili smo produkte z zelo ozko porazdelitvijo molskih mas (PDI = 1,00–1,09) Povpre~ja molskih mas polimerov so bila dolo~ena z metodo izklju~itvene kromatografije z detektorjem na sipanje svetlobe. Molska masa produktov ni bila odvisna od temperature, pri kateri smo izvajali reakcije (25–40 °C), mo~no pa je odvisna od uporabljenega iniciatorja polimerizacije (trietilamin vs. n-pentilamin) ter ~istosti izhodnega monomera. Tako smo z uporabo trietilamina ob enakih reakcijskih pogojih iz trikrat prekristaliziranega mono- mera v primerjavi z le enkrat kristaliziranim pripravili poli(ß-benzil L-aspartat) s trikrat ve~jo molsko maso (9,2 × 103 g mol–1) Reakcija z n-pentilaminom prav tako pote~e do prakti~no monodisperznih polimerov (PDI = 1,00), a so povpre~ja molskih mas za ~as reakcije 24 h (nepopolna konverzija) nekoliko ni`ja. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 309 Da bi se izognili potencialnim stranskim reakcijam in ob~utljivosti monomera za vlago, smo vse reakcije izvajali v suhi argonovi atmosferi. Reakcijo kopolimerizacije z laktidom smo najprej izvajali v masi pri temperaturi 140 °C. A pri tej tem- peraturi prihaja do hidrolize benzilnih estrskih skupin na asparaginskih enotah. Iz tako spro{~enih karboksilnih skupin na glavni verigi pa lahko rastejo stranske polilaktidne verige. Rezultat so kopolimeri s pripaja- njem, katerih strukture ni mogo~e nadzorovati, saj pri teh reakcijskih pogojih obseg hidrolize estrskih skupin ni kontroliran. Zato smo nadaljnje reakcije izvajali v razli~nih topilih (THF, dioksan, N,N-dimetilformamid) pri temperaturi 50 °C, saj smo z NMR spektroskopijo ugotovili, da so v teh topilih pri omenjenih temperaturah benzilne estrske skupine kvantitativno obstojne vsaj 24 h. V prej{nji stopnji pripravljeni poli(ß-benzil L-aspar- tat) nam v tem koraku hkrati rabi kot reagent in skupaj s predestiliranim kositrovim(II) bis(2-etilheksanoatom) kot koiniciator polimerizacije z odpiranjem obro~a lak- tida. Tak na~in sinteze omogo~a pripravo nizkopoli- disperznih blokkopolimerov asparaginske kisline in laktida, v prihodnje pa bo potrebna optimizacija reakcij- skih parametrov sinteze. SYNTHESIS OF POLY(ß-BENZYL L-ASPARTATE) AND ITS COPOLYMERS WITH L,L-LACTIDE Bla` Brulc, Maja Gri~ar, Ida Poljan{ek, Ema @agar, Majda @igon National Institute of Chemistry, Hajdrihova 19, SI-1001 Ljubljana, Slovenia In our work poly(ß-benzyl L-aspartate)s and their block copolymers with L,L-lactide with varying molar mass averages and low polydispersity indices (PDI = 1.00–1.09) were prepared. NMR and FT-IR spectro- scopy was used to elucidate the products’ chemical composition, and size-exclusion chromatography coupled to multi-angle laser photometer (SEC-MALLS) was used for the determination of the absolute molar mass averages of the products. Polymerizations of aspartic acid N-carboxyan- hydrides were carried out in dry N,N-dimethylformamide at room temperature and at slightly elevated tempe- ratures (up to 40 °C) in a dry argon atmosphere using triethylamine or n-pentylamine as the initiatior. The polymerization mechanism is strongly dependent on the nature of the initiator used, both mechanistic pathways giving practically monodisperse products. Homo- polymers synthesized in the described way were further used as macrocoinitiators in the ring-opening poly- merization of L,L-lactide using stannous(II) octoate as the catalyst. NMR spectra and SEC-MALLS chromatograms showed that copolymers synthesized in this manner were linear, but those prepared at higher temperatures exhibited some degree of branching due to partial hydrolysis of pendant benzylic ester groups. Deprotected carboxylic groups act as reactive sites for either grafted polyaspartate or lactide side chains. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 310 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Nata{a Drnov{ek DVOPLASTNA PREVLEKA NA ZLITINI Ti6Al4VZA BIOMEDICINSKO UPORABO Nata{a Drnov{ek1, Nina Daneu1, Sa{a Novak1, Katja Rade1, Janez Kova~2 1Odsek za nanostrukturne materiale, IJS, Ljubljana, Slovenija 2Odsek za tehnologijo povr{in in optoelektroniko, IJS, Ljubljana, Slovenija Titanove zlitine se zaradi svojih dobrih lastnosti, kot so relativno nizek elasti~ni modul, visoka trdnost, koro- zijska odpornost in biokompatibilnost, {iroko uporabljajo za kostne vsadke. Kjub dobrim lastnostim pa klini~ni re- zultati ka`ejo, da se dolo~en dele` vsadkov razmaje, kar je delno pripisano nezadovoljivi pritrditvi vsadka. Zato sta za trajne vsadke potrebni dobra pritrditev in stabilnost, torej dobro vra{~anje kostnega tkiva v vsadek, kar lahko dose`emo z modificiranjem povr{ine zlitine. Oseointegracijo vsadka in tkiva je mo`no izbolj{ati z bioaktivnimi prevlekami, ki omogo~ijo neposredno povezavo vsadka s kostjo. Tak{ni bioaktivni prevleki sta hidroksiapatit in biosteklo. Pri na{ih raziskavah smo uporabili bioaktivno steklo. Na vsadek smo najprej nanesli tanko prevleko iz TiO2, ki izbolj{a adhezijo bioaktivne prevleke na podlago iz Ti-zlitine, prav tako izbolj{a biokompatibilnost, korozijsko odpornost ter prepre~i difuzijo kovinskih atomov v telo. V tem delu smo plast TiO2 na povr{ini Ti6Al4V pripravili s hidrotermalno obdelavo. Da bi dobili pribli`no 100 nm debelo oksidno plast TiO2-strukture anataz, ki je bolj bioaktiven od rutila, smo spreminjali pogoje obdelave, kot so temperatura, tlak, ~as in pH suspenzije. Preu~ili smo tudi, kako dodatna toplotna obdelava, ki je potrebna za sintranje zgornje bioaktivne prevleke, vpliva na `e s hidrotermalno obdelavo nastalo plast TiO2. Z analizami XPS, AES, SEM in TEM smo preverili debelino, strukturo in proces nastanka oksidne plasti. 100 nm debelo oksidno plast iz TiO2 anataza (slika 1) smo dobili pri dovolj visokem pH suspenzije (8) z dodatkom ionov Ti4+ po 24 h hidrotermalne obdelave pri 150 °C. Dodatna toplotna obdelava je pove~ala debelino oksidne plasti, vendar so visoke temperature (750 °C) povzro~ile, da je oksid za~el odpadati od podlage. Na TiO2-plast smo nato z elektroforetsko depozicijo nanesli {e prevleko iz bioaktivnega stekla in `e po eni minuti smo dobili 150 µm debelo plast, ki se je po MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 311 Slika 2: Prevleka iz biostekla na zlitini Ti6Al4V po sintranju Slika 1: a) TEM-posnetek TiO2-plasti na Ti6Al4V in b) elektronska difrakcija te plasti, ki potrjuje, da je oksid anataz sintranju dobro dr`ala podlage iz Ti-zlitine in ni imela razpok (slika 2). A DOUBLE-LAYER COATING ON A Ti6Al4V ALLOY FOR BIOMEDICAL APPLICATIONS Nata{a Drnov{ek1, Nina Daneu1, Sa{a Novak1, Katja Rade1, Janez Kova~2 1Jo`ef Stefan Institute, Department for Nanostructured Materials, Ljubljana, Slovenia 2Jo`ef Stefan Institute, Department of thin films and surfaces, Ljubljana, Slovenija Titanium alloys are widely used as implant materials because of their good properties, such as a relatively low elastic modulus, a high specific strength, good corrosion resistance and superior biocompatibility. Clinical results, however, reveal a certain percentage of loosening of hip prostheses, partly ascribed to insufficient fixation. Hence, in order to improve the fixation and stability of an implant, a surface modification of the alloy is needed. Implant-tissue osseointegration can be enhanced by bioactive coatings, such as hydroxyapatite or bioactive glass, which are able to provide direct bonding of an implant to the host bone. According to literature data, a thin TiO2 layer at the surface of the Ti-alloy improves the corrosion resistance, prevents the diffusion of metal ions in the body as well as providing better adhesion of the bioactive coating to the alloy substrate. The main goal of our studies is to improve the bone-implant bonding ability and to enhance the osseointegration of the implant by forming a functionally graded bioactive coating on the Ti-alloy substrate. In this particular work, the TiO2 layer was formed on the surface of Ti6Al4V by a hydrothermal treatment of the alloy under various conditions. The thickness of the TiO2 layer was further increased by thermal treatment. Then, a bioactive layer was deposited over the TiO2 layer using a dip-coating method. The microstructural properties of the coating were studied using SEM, TEM and XPS analyses. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 312 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Jakob König, univ. dipl. ing. kem., rojen 23. 8. 1979 na Jesenicah. Zaposlitev: Institut "Jo`ef Stefan", Odsek za raziskave sodobnih materialov V {tudijskem letu 2007/08 {tudent 4. letnika podiplom- skega {tudija na Mednarodni podiplomski {oli Jo`efa Stefana Podro~je dela: materiali, elektronska keramika Zaradi pove~ane skrbi za varovanje zdravja ljudi in ohranitev okolja se je v zadnjih letih mo~no pove~alo {tevilo raziskav alternativnih materialov. Tako se na podro~ju feroelektri~nih materialov razvijajo materiali, ki ne vsebujejo svinca ("lead-free materials"). Mednje spada tudi Na0,5Bi0,5TiO3, katerega lastnosti so odvisne od zunanjih pogojev (npr. od zunanjega elektri~nega po- lja ali od aksialne tla~ne obremenitve). Ta odvisnost nam daje mo`nost prilagajanja lastnosti materiala ("tunable materials") in s tem odziva komponente oz. naprave. V na{em delu raziskujemo vpliv aksialne tla~ne obreme- nitve na dielektri~ne lastnosti materialov na osnovi Na0,5Bi0,5TiO3. POVE^ANJE VPLIVA AKSIALNE TLA^NE OBREMENITVE NA DIELEKTRI^NE LAST- NOSTI Na0,5Bi0,5TiO3 Z DODAJANJEM NaTaO3 Jakob König, Bo{tjan Jan~ar, Danilo Suvorov Odsek za raziskave sodobnih materialov, Institut "Jo`ef Stefan", Ljubljana [tudije vpliva zunanjih mehanski sil na dielektri~ne lastnosti Na0,5Bi0,5TiO3 so pokazale, da se vrednost dielektri~ne konstante pod vplivom aksialne tla~ne obremenitve zni`uje. To zni`anje je najve~je v obmo~ju maksimuma dielektri~ne konstante pri 320 °C, medtem ko je pri sobni temperaturi ta vpliv veliko manj{i. To je posledica velikega mehanskega koercitivnega polja, ki ga lahko pribli`no ocenimo iz elektri~nega koercitivnega polja (73 kV/cm). Cilj na{ega dela je bilo pove~anje vpliva tla~ne obremenitve na dielektri~ne lastnosti Na0,5Bi0,5TiO3 pri sobni temperaturi, kar smo dosegli z dodatkom ustreznega materiala (NaTaO3). Z dodatkom NaTaO3 se vpliv tla~ne obremenitve na dielektri~nost pove~a, saj dodatek NaTaO3 pomakne dielektri~ni maksimum k ni`jim temperaturam (slika 1) in zni`a elektri~no koercitivno polje. Vpliv aksialne tla~ne obremenitve je najve~ji pri vzorcu s 15 % NaTaO3, kjer je relativna sprememba dielektri~nosti 14 % (slika 2). Dielektri~ne lastnosti niso reverzibilne s prenehanjem tla~ne obremenitve, vendar v primeru, ko vzorec segrejemo na 600 °C za 30 min, ponovno dobimo za~etne dielektri~ne lastnosti materiala. INCREASING THE EFFECT OF AXIAL PRESSURE ON THE PERMITTIVITY OF Na0.5Bi0.5TiO3 BY ADDING NaTaO3 Jakob König, Bo{tjan Jan~ar, Danilo Suvorov Advanced Materials Department, Jo`ef Stefan Institute, Jamova 39, Ljubljana Na0.5Bi0.5TiO3 is one of the most studied lead-free ferroelectric materials. Studies of the influence of MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 313 Slika 2: Vrednost dielektri~ne konstante brez obremenitve in pri tlaku 200 MPa za razli~ne dodatke NaTaO3 (merjeno pri sobni temperaturi). Slika 1: Premikanje dielektri~nega maksimuma k ni`jim temperatu- ram (10 % dodatka 270 °C, 30 % dodatka 0 °C). mechanical forces on the electrical properties of Na0.5Bi0.5TiO3 have shown that the dielectric constant decreases under an applied axial pressure. The reduction of the permittivity is the largest in the region of the dielectric maximum and is much smaller at room temperature, which is connected with the large mechanical coercive field, closely linked with the electrical coercive field (73 kV/cm). The pressure’s influence on permittivity at room temperature can be increased by an appropriate choice of modifying material. In our study we investigated the influence of NaTaO3 additions on the properties of Na0.5Bi0.5TiO3. The axial pressure dependence of permittivity at room temperature of the materials from the Na0.5Bi0.5TiO3-NaTaO3 solid-solution system will be presented. With the addition of NaTaO3 the pressure dependence of the permittivity increases as the dielectric anomalies are shifted towards lower temperatures and the coercive field is lowered. The maximum of the pressure dependence of the permittivity was achieved in a sample with 15 % of NaTaO3. The dielectric properties of the materials are time dependent, and are not fully reversible when the pressure is removed. However, after annealing the samples at 600 °C for 30 min the initial dielectric properties are recovered. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 314 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Katja Rade [TUDIJ POLIMETAKRILNE KISLINE V PRISOTNOSTI RAZLI^NIH KATIONOV V VODNEM MEDIJU Katja Rade1, Ksenija Kogej2, Sa{a Novak1 1Institut "Jo`ef Stefan", Odsek za nanostrukturne materiale 2FKKT, katedra za fizikalno kemijo Polimetakrilna kislina (PMA) je najpreprostej{a polimerna kislina, ki je izpostavljena konformacijskemu prehodu, posledi~no pa je dobra modelna spojina za razumevanje npr. denaturacije molekul proteinov. Zaradi svoje sposobnosti ohranjanja kompaktne oblike pri nizkih pH-vrednostih (kot je v `elodcu) in prehodu v obliko naklju~nega klob~i~a pri vi{jih pH-vrednostih (debelo ~revo) je uporabna tudi za ciljano dostavo zdravilnih u~inkovin, ki so ob~utljive za kislo okolje. Za uporabo PMA v te namene je potrebno podrobno poznanje vedenja spojine. Zaenkrat {e ni veliko znanega o konformacijskem prehodu spojine v prisotnosti katio- nov razli~nih valenc. S potenciometri~nimi, kalorime- tri~nimi in fluorimetri~nimi titracijami smo opazovali obna{anje PMA v prisotnosti ionov Li+, Mg2+ in La3+ (v 0,1 M raztopinah kloridov). ^eprav je proces s termo- dinamskega stali{~a neugoden, razvijanje molekule pote~e v vseh primerih, vrednost Gibbsove proste ental- pije pa se z nara{~ajo~o valenco prisotnega kationa manj{a. V navzo~nosti lantanovih ionov zaradi visokih nabojev verjetno nastane tudi medmolekulska asociacija. Atakti~no PMA smo preizkusili tudi kot dispergant za suspenzijo prahu aluminijevega oksida (Al2O3). Izka`e se, da je PMA ustrezen dodatek, ki zaradi svoje polimerne narave in negativno nabitih karboksilatnih ostankov deluje kot elektrosteri~ni stabilizator, vendar se njegovo vedenje razlikuje, ~e titracijo izvedemo iz kislega v bazi~no pH-obmo~je kot nasprotno. Pri nizkih pH-vrednostih se PMA ve`e na pozitivna mesta na delcih Al2O3, s tem nevtralizira neto naboj in posledi~no zni`a vrednost zetapotenciala (ZP). Pri vi{jih pH-vred- nostih se negativno nabita PMA ve`e na preostala pozitivna mesta na ve~inoma negativno nabitih delcih Al2O3 in s tem {e pove~a absolutno vrednost ZP. V obmo~ju zelo visokih pH se ZP zaradi prisotnosti ionov Na+ in Al3+, ki interagirajo z negativno nabitimi karboksilatnimi ostanki, ponovno zmanj{a in je celo ni`ji kot v primeru suspenzije brez dodatkov. Poleg tega k zmanj{anju ZP izdatno pripomore pove~ana prevodnost, ki je posledica dodajanja mo~nih elektrolitov (HCl, NaOH); hipotezo smo preverili z dodajanjem nevtral- nega elektrolita (NaCl) suspenziji Al2O3. Z na{im delom smo `eleli pokazati, da bolj poglob- ljen princip in opazovanje vedenja dispergantov lahko prispeva dodatne informacije in s tem prispeva k uspe{nej{i pripravi suspenzij z `elenimi lastnostmi. STUDY OF POLYMETHACRYLIC ACID IN PRESENCE OF VARIOUS CATIONS IN AQUEOUS MEDIA Polymethacrylic acid (PMA) is an interesting compound for various reasons. It is the simplest poly- meric acid that is subjected to conformational transition MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 315 Slika 2: Vedenje suspenzije Al2O3 ob dodatku PMA pri titraciji v kislo ali bazi~no ter pri dodatku nevtralnega elektrolitaSlika 1: Polimetakrilna kislina v kompaktni in iztegnjeni obliki and so it appears very useful in understanding of conformational transitions of proteins. Because of its ability to stay in compact form at lower pH values (stomach) and transition into extended conformation at higher pH values (gall bladder) it can also be used in targeted drug delivery system. In order to use PMA in such ways it is very important to know its behaviour in details. There were no studies yet made about its conformational transition in the presence of cations of different valencies. We made comparison of PMA behaviour in the presence of Li+, Mg2+ and La3+ ions (all in 0.1 M solution of chloride salt) by means of potentiometric, calorimetric and fluorimetric titration. The unfolding occurs although it is thermodynamically unfavourable. The Gibbs free energy reduces from the case without added electrolyte to solutions with added lanthanum ions. In the latter case there is also a strong indication of intermolecular association because of triple charge of lanthanum ions and high negative charge on PMA. Atactic PMA was also used as a deflocculant in case of submicrometer sized aluminium oxide. It appears that PMA can be successfully used (because of its polymeric nature it acts as electrosteric stabilizer) but shows different behaviour if titrated from acidic to basic pH than vice versa. At lower pH PMA binds to positive sites onto alumina particles, neutralizes the net charge and therefore decreases average zeta potential (ZP). At higher pH the negatively charged PMA binds onto the positive sites on particles’ surface which is mainly negatively charged and increases the absolute value of ZP. In comparison to the case with no added PMA The ZP is again decreased at very high pH values because of the presence of Na+ and Al3+ ions in the solution (these interact with negatively charged carboxylic groups). MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 316 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 Ines Bra~ko, univ. dipl. in`. kem., rojena 24. 4. 1977 v Mariboru Zaposlena na Institutu "Jo`ef Stefan", Odsek za raziska- ve sodobnih materialov V {tudijskem letu 2007/08 {tudentka 2. letnika doktor- skega {tudija na Mednarodni podiplomski {oli Jo`efa Stefana Podro~je dela: 1D-nanostrukturni materiali V zadnjih desetih letih je bilo veliko pozornosti namenjene raziskavam in sintezi 1D-nanostrukturnih materialov (nanocevke, nanopal~ke, nano`i~ke). Zaradi unikatnih elektri~nih, mehanskih, magnetnih in opti~nih lastnosti, ki so posledica velikosti (majhnosti) in veli- kega razmerja med povr{ino in volumnom teh nano- struktur, so ti materiali zanimivi za {tevilne aplikacije v nanonapravah prihodnosti. Hidrotermalna sinteza je obetajo~a tehnika za sintezo razli~nih perovskitov z 1D- nanostrukturo. Vendar pa je o reakcijskih mehanizmih, ki potekajo med tak{no sintezo in so klju~nega pomena za pripravo `elenih produktov, malo znano. V na{i {tudiji uporabljamo kot modelni sistem kalcijev titanat – CaTiO3. RAZUMEVANJE NASTANKA NANOSTRUKTURNEGA PEROVSKITA CaTiO3 POD HIDROTERMALNIMI POGOJI Ines Bra~ko1, Bo{tjan Jan~ar1, Sa{o [turm2, Danilo Suvorov1 1Odsek za raziskave sodobnih materialov, Institut "Jo`ef Stefan", Ljubljana, Slovenija 2Odsek za nanostrukturne materiale, Institut "Jo`ef Stefan", Ljubljana, Slovenija V raziskavi smo za pripravo nanostrukturnega kalcijevega titanata pod hidrotermalnimi pogoji uporabili vodno raztopino kalcijevega acetata in titan(IV)izo- propoksida v mo~no alkalnem mediju. Za karakterizacijo produktov smo uporabili razli~ne tehnike, povezane s presevno elektronsko mikroskopijo (TEM). V {tudiji predstavljamo: • nanostrukturni kalcijev titanat v obliki plastovitih monokristalini~nih dobro kristaliziranih kristalov (slika spodaj); • vmesne nanostrukturne faze: nanocevke, amorfne nanodelce, tanke plastovite nanostrukture in nano- `i~ke (slika spodaj). Z uporabo spektroskopije merjenja izgub energije elektronov (EELS) in z analizo finih struktur Ti-L2,3 in ionizacijskih robov O-K (ELNES) smo ugotovili, da so nanocevke po sestavi in strukturi blizu rutilni fazi TiO2. EELS analiza je nadalje pokazala, da je vgrajevanje Ca mo`no v drugem sklopu vmesnih faz (amorfni nanodelci, plastovite nanostrukture in nano`i~ke), ki se sicer razlikujejo po sestavi, morfologiji in kristalini~nosti. Menimo, da je mo`no te vmesne faze Ca–Ti uporabiti kot matrico za sintezo 1D-nanostrukturnega perovskita CaTiO3 in da je s kontroliranjem nastanka metastabilnih faz mogo~e kontrolirati morfologijo nanostrukturnega CaTiO3. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 317 UNDERSTANDING THE FORMATION OF NANOSTRUCTURED PEROVSKITE CaTiO3 UNDER HYDROTHERMAL CONDITIONS Ines Bra~ko1, Bo{tjan Jan~ar1, Sa{o [turm2, Danilo Suvorov1 1Advanced Materials Department, Jo`ef Stefan Institute, Ljubljana, Slovenija 2Nanostructured Materials Department, Jo`ef Stefan Institute, Ljubljana, Slovenija One dimensional nanostructures such as nanotubes and nanowires are attracting great interest of research community due to their potential applications, which arise from their unique optical, mechanical and electrical properties. Over the past decade many techniques have been developed and applied for the synthesis of one dimensional nanostructured materials. Employment of hydrothermal method for preparation of nanostructured materials enables one to synthesize different nano- structured perovskites. Several successful syntheses of perovskite nanostructured materials and also suggestions to apply hydrothermal condition for preparation of nanowires of any tetragonal perovskite have been reported in the literature. However, there is still little known about the reaction mechanisms present in the hydrothermal synthesis. In order to understand the formation mechanisms that occur during synthesis of perovskite-based nanowires via hydrothermal route we studied formation of nano- structured CaTiO3 from Titanium(IV) isopropoxide and Calcium acetate aqueous solution. Low temperature reactions were performed in a Teflon autoclave in a highly alkaline environment. In our study, we present different intermediate phases taken from various reaction steps during hydrothermal synthesis, which differ in morphology, crystal structure and composition. High- Resolution Transmission Microscopy (HRTEM), Energy Dispersive X-ray Spectroscopy (EDS) and Electron Energy-Loss Spectroscopy (EELS) analyses revealed presence of at least four phases. Morphologically the most interesting phases important for formation of perovskite one dimensional nanostructures were single-crystalline sheets and nanowires. Nanowires with Ca:Ti ratio 1:3 and diameter of about 10 nm were 100 nm long. Single-crystalline sheets were 35nm thick and dimensions of about 500 nm × 500 nm with the Ca:Ti ratio approximately 0,5. Further analysis of EEL spectrum, introduction of the X-ray Powder Diffraction (XRD) and Differential Thermal Analysis (DTA) revealed that the crystal structure of single-crystalline sheets can be interpreted as a layered structure similar to either kassite – CaTi2O5(OH)2 or cafetite – CaTi2O5(H2O). We believe that these phases can be employed as a template precursor for synthesis of one-dimensional nanostructured CaTiO3 based perovskites and that by controlling the formation of the metastable phases with high aspect ratios under hydrothermal conditions the morphology of CaTiO3 nanostructures can be controlled. MLADI RAZISKOVALCI – NAGRAJENCI 15. KONFERENCE O MATERIALIH IN TEHNOLOGIJAH 318 Materiali in tehnologije / Materials and technology 41 (2007) 6, 305–318 DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES DOKTORSKA DELA – DOCTOR'S DEGREES Na Naravoslovnotehni{ki fakulteti Univerze v Ljubljani je dne 9. 11.2007 pred komisijo v sestavi: izr. prof. dr. Jakob Likar kot predsednik in ~lani: red. prof. dr. Ladislav Kosec, red. prof. dr. Franc Vodopivec, red. prof. dr. Savo Spai}, red. prof. dr. Radomir Turk mag. Franc Tehovnik, univ. dipl. in`. metal. zagovarjal doktorsko disertacijo z na- slovom: Preoblikovalna sposobnost avstenitnih nerjavnih jekel Workability of austenitic stainless steels Doktorska disertacija je bila izdelana pod mentorstvom red. prof. dr. Ladislava Kosca ter somentorstvom red. prof. dr. Franca Vodopivca. PREOBLIKOVALNA SPOSOBNOST AVSTENITNIH NERJAVNIH JEKEL UDK: 669.14.018.8:621.771 POVZETEK Delo obravnava pregled strokovne literature o strjevalni mikrostrukturi in preoblikovanju v vro~em avstenitnih nerjavnih jekel, legiranih z molibdenom. Strjevanje se za~ne v odvisnosti od vsebnosti glavnih legirnih elementov, s primarno kristalizacijo -ferita ali avstenita. Preoblikovalnost jekla s strjevalno mikrostrukturo v vro~em je bolj{a pri primarni kristalizaciji -ferita, ker ta prepre~i koncentracijo ne~isto~ v preostali meddendritni talini. Mehanizma meh~anja med preoblikovanjem v vro~em avstenitnih nerjavnih jekel sta dinami~na in stati~na rekristalizacija. Vpliv pogojev deformacije na procese spro{~anja deformacijske energije v povezavi s kemijsko sestavo jekla in vsebnostjo -ferita smo ugotovili na podlagi nateznih preskusov v vro~em, ki pa niso omogo~ili ocene, kako se v za~etnih valjarni{kih prehodih vede jeklo s strjevalno mikrostrukturo. Avstenitna nerjavna jekla do volumenskega dele`a -ferita 12,5 % v strjevalni mikrostrukturi pri upogibni deformaciji do 28 %, niso izpostavljena nastanku razpok na povr{ini in na robovih presku{ancev. -ferit ni zmanj{al deformabilnosti jekla v temperaturnem obmo~ju med 1050 °C in 1250 °C, ne glede, ali je bil pri ni`ji temperaturi razporejen v obliki zvezne ali deloma prekinjene mre`e okoli zrn avstenita, in pri vi{ji DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES Materiali in tehnologije / Materials and technology 41 (2007) 6, 319–320 319 WORKABILITY OF AUSTENITIC STAINLESS STEELS UDC: 669.14.018.8:621.771 ABSTRACT In the first part of this thesis a survey of the solidification structure and hot workability of austenitic stainless steels alloyed with molybdenum is presented. The solidification occurs, depending on the content of the main alloying elements, with the primary solidification of -ferrite or austenite. The hot workability of the as-solidified microstructure is better in the case of the primary solidification of -ferrite, because it prevents the concentration of impurities in the residual interdendritic melt. The mechanisms of softening during the hot working of austenitic stainless steels are dynamic and static recrystalization. In the experimental work the influence of deformation on the processes of the relaxation of deformation energy, related to the composition of the steel and the quantity of -ferrite, was determined with hot tensile tests. These tests did not allow to evaluate the behaviour of as-cast steel during the first rolling passes. Austenitic stainless steels with a solidification microstructure of up to 12.5 % of -ferrite were not sensitive to hot cracking of the surfaces or edges during hot bending tests with 28 % of deformation, in the temperature range from 1050 °C to 1250 °C. At lower temperatures the -ferrite was in the form of a continuous or discontinuous layer at the austenite grain temperaturi, ko je ferit v ovalnih delcih, ki so lo~eni s {iroko plastjo avstenita. Nerjavno jeklo z masnim dele`em svinca 0,0082 % s strjevalno mikrostrukturo je ob~utljivo za nastanek razpok po interdendritnih povr{inah, na katerih je ve~je {tevilo zrn svinca, ki jih je na meje odrinila napredujo~a fronta strjevanja. Z vro~im valjanjem klinastih preizku{ancev smo ugotovili temperaturo in stopnjo deformacije, pri kateri je jeklo ob~utljivo za nastanek valjarni{kih napak. Med vro~im valjanjem je -ferit enako duktilen kot avstenit. Pri temperaturah preoblikovanja 1100 °C in 1150 ° in pri okoli 30-odstotni stopnji deformacije se na robovih vzorcev pojavijo pre~ne razpoke, v notranjosti vzorcev pa mikrorazpoke med avstenitom in -feritom, ki so posledica izcejanja `vepla. Klju~ne besede: avstenitna nerjavna jekla, konti- nuirno litje, strjevalna mikrostruktura, -ferit, inter- kristalne razpoke, preoblikovalnost v vro~em, vpliv ne~isto~, mikrorazpoke DOKTORSKA, MAGISTRSKA IN DIPLOMSKA DELA – DOCTOR'S, MASTER'S AND DIPLOMA DEGREES 320 Materiali in tehnologije / Materials and technology 41 (2007) 6, 319–320 boundaries, while at higher temperatures the -ferrite was in the form of rounded particles surrounded by the austenite matrix. The as-cast stainless steel with 0.0082 wt. % of lead was sensitive to interdendritic crack formation as the number of insoluble lead particles was increased at the dendritic surface, when the particles were pushed by the solidification front. With hot rolling of a wedge sample the temperature and the deformation step when the steel became sensitive to hot cracks were determined. Edge cross cracks appeared in the temperature range between 1100 °C and 1150 °C at around a 30 % step of deformation. Inside the hot-rolled samples microcracks were observed at the austenite and -ferrite boundaries, where the segregation of sulphur was observed, also. Key words: austenitic stainless steel, continuous casting, solidification microstructure, -ferrite, grain- boundary cracks, hot workability, influence of impurities LETNO KAZALO – INDEX Letnik / Volume 41 2007 ISSN 1580-2949 © Materiali in tehnologije IMT Ljubljana, Lepi pot 11, 1000 Ljubljana, Slovenija M EHNOLOGIJEIN ATER IALI M A T E R I A L S A N D T E C H N O L O G Y MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY VSEBINA / CONTENTS LETNIK / VOLUME 41, 2007/1, 2, 3, 4, 5, 6 2007/1 Materiali in tehnologije – 40 let Materials and technology – 40 years F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Prvi urednik revije – Jo`a Arh F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6 Laudation in honour of professor dr. Franc Vodopivec on the occasion of his 75th birthday M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 75 let Franca Vodopivca L. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Marin Gabrov{ek – ob jubileju F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 Zgodovina serijske publikacije Materiali in tehnologije / Materials and technology Historical overview of the scientific journal Materiali in tehnologije / Materials and technology N. Jamar, J. Jamar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 Theoretical calculation of the lubrication-layer thickness during metal drawing Teoreti~ni izra~un debeline plasti maziva pri vle~enju kovin D. ]ur~ija, I. Mamuzi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 The notch effect on the fatigue strength of 51CrV4Mo spring steel VPliv zareze na trajno nihajno trdnost vzmetnega jekla 51CrV4Mo B. [u{tar{i~, B. Sen~i~, B. Arzen{ek, P. Jodin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29 An integrity analysis of washing-machine holders Analiza celovitosti nosilca kadi v pralnem stroju N. Gubeljak, M. Mejac, J. Predan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35 The effect of a material’s heterogeneity on the stress and strain distribution in the vicinity of a crack front Vpliv heterogenosti materiala na porazdelitev napetosti in deformacije v bli`ini konice razpoke D. Kozak, N. Gubeljak, J. Vojvodi~ Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41 The numerical solution of strain wave propagation in elastical helical spring Numeri~na re{itev propagacije deformacijskega vala v elasti~ni spiralni vzmeti S. Ayadi, E. Hadj-Taïeb, G. Pluvinage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47 [tudij notranje oksidacije v naoglji~enih hitrostrjenih trakovih Cu The study of the internal oxidation in internally carbonised Cu ribbons R. Rudolf, L. Kosec, I. An`el, L. Gusel, M. Poharc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53 2007/2 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 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85 322 MATERIALI IN TEHNOLOGIJE 41 (2007) 6 LETNO KAZALO – INDEX 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 2007/3 Structural steels with micrometer grain size: a survey Konstrukcijska jekla z mikrometrskimi kristalnimi zrni: pregled F. Vodopivec, D. Kmeti~, F. Tehovnik, J. Vojvodi~-Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111 The implementation of an online mathematical model of billet reheating in an OFU furnace Implementacija simulacijskega modela za spremljanje ogrevanja gredic v OFU-pe~i A. Jakli~, F. Vode, T. Marolt, B. Kumer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119 The behaviour of coarse-grain HAZ steel with small defects during cyclic loading Vedenje jekla grobozrnatega TVP z napakami pri cikli~ni obremenitvi V. Gliha, T. Vuherer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 125 The effect of cold work on the sensitisation of austenitic stainless steels Vpliv hladne deformacije na pove~anje ob~utljivosti nerjavnih jekel M. Dománková, M. Peter, M. Roman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131 Fatigue-crack propagation near a threshold region in the framework of two-parameter fracture mechanics Dvoparametrska lomno mehanska analiza hitrosti utrujenostne razpoke blizu praga propagacije S. Seitl, P. Huta . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 135 Modelling of the solidification process and the chemical heterogeneity of a 26NiCrMoV115 steel ingot Modeliranje procesa strjevanja in kemi~ne heterogenosti ingota iz jekla 26NiCrMoV115 M. Balcar, R. @elezný, L. Martínek, P. Fila, J. Ba`an . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 Frekven~na odvisnost rezidualnega trenja viskoznostnega vakuumskega merilnika z lebde~o kroglico Frequency dependence of spinning rotor gauge residual drag J. [etina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145 A wet-steam pipeline fracture Prelom cevovoda za vla`no paro R. Celin, D. Kmeti~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151 2007/4 A fatigue characterization of honeycomb sandwich panels with a defect Utrujenostna karakterizacija satastih sendvi~nih panelov z napako B. Keskes, Y. Menger, A. Abbadi, J. Gilgert, N. Bouaouadja, Z. Azari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 Fatigue properties of a high-strength-steel welded joint Utrujenostne lastnosti zvara visokotrdnega jekla Z. Burzi}, V. Grabulov, S. Sedmak, A. Sedmak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 163 Mehanske lastnosti zvara iz jekla maraging po izlo~evalnem `arjenju Mechanical properties of maraging steel welds after aging heat treatment D. Klob~ar, J. Tu{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 167 An experimental verification of numerical models for the fracture and fatigue of welded structures Eksperimentalna verifikacija numeri^nih modelov za prelom in utrujenost zvarjenih struktur S. Sedmak, A. Sedmak, M. Arsi}, J. Vojvodi~ Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173 Izra~un parametrov Weibullove porazdelitve za oceno upogibne trdnosti valovitih stre{nih plo{~ Computation of the parameters of the Weibull distribution for estimating the bending strength of corrugated roofing sheets M. Ambro`i~, K. Vidovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 179 Investigation of the influence of the melt slag regime in a ladle furnace on the cleanliness of the steel Raziskava vpliva re`ima `lindre v ponov~ni pe~i na ~istost jekla Z. Adolf, I. Husar, P. Suchánek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 185 LETNO KAZALO – INDEX MATERIALI IN TEHNOLOGIJE 41 (2007) 6 323 The influence of illite-kaolinite clays’ mineral content on the products’ shrinkage during drying and firing Vpliv vsebnosti glin ilinit-kaolinit na kr~enje pri su{enju in `ganju M. Krgovi}, N. Marstijepovi}, M. Ivanovi}, R. Zejak, M. Kne`evi}, S. \urkovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189 The application of spheroidal graphite cast iron in Bosnia and Herzegovina Uporaba nodularne grafitne litine v Bosni in Hercegovini D. Pihura, M. Oru~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193 2007/5 [estdeset let prof. dr. Vasilija Pre{erna Laudation in honour of Professor Dr. Vasilij Pre{ern on the occasion of his 60th birthday M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 The oxidation and reduction of chromium during the elaboration of stainless steels in an electric arc furnace Oksidacija in redukcija kroma iz `lindre med izdelavo nerjavnih jekel v elektrooblo~ni pe~i B. Arh, F. Tehovnik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 203 A new topology for the trajectories of the meniscus during continuous steel casting Nova topologija trajektorij meniskusa pri neprekinjenem litju jekla I. B. Risteski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213 Multiscale modelling of short cracks in random polycrystalline aggregates Ve~nivojsko modeliranje kratkih razpok v naklju~nih ve~kristalnih skupkih L. Cizelj, I. Simonovski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227 Changes to the fracture behaviour of medium-alloyed ledeburitic tool steel after plasma nitriding Spremembe v na~inu preloma srednje legiranega ledeburitnega jekla zaradi plazemskega nitriranja P. Jur~i, F. Hnilica, J. Cejp . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 231 The fracture and fatigue of surface-treated tetragonal zirconia (Y-TZP) dental ceramics Prelom in utrujenost povr{insko obdelane tetragonalne (Y-TZP) dentalne keramike T. Kosma~, ^. Oblak, P. Jevnikar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 237 Povr{ina zlitine Cu-Sn-Zn-Pb po obsevanju z ultravijoli~nim du{ikovim laserjem Surface of Cu-Sn-Zn-Pb alloy irradiated with ultraviolet nitrogen laser F. Zupani~, T. Bon~ina, D. Pipi}, V. Hen~ - Bartoli} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243 A preliminary S-N curve for the typical stiffened-plate panels of shipbuilding structures Preliminarna krivulja S-N za toge plo{~ate panele za ladjedelni{ke strukture L. Gusha, S. Lufi, M. Gjonaj . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 249 2007/6 A new version of the theory of ductility and creep under cyclic loading Nova verzija teorije o duktilnosti in lezenju pri cikli~ni obremenitvi L. B. Getsov, M. G. Kabelevskiy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Thermoelectrical properties of a monocrystalline Al64Cu23Fe13 quasicrystal Termoelektri~ne lastnosti monokristalnega kvazikristala Al64Cu23Fe13 I. Smiljani}, A. Bilu{i}, @. Bihar, J. Lukatela, B. Leonti}, J. Dolin{ek, A. Smontara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 265 Faze v kvazikristalni zlitini Al64,4Cu22,5Fe13,1 Phases in a quasicrystalline alloy Al64,4Cu23,5Fe13,1 T. Bon~ina, B. Markoli, I. An`el, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 271 Hydrogen absorption by Ti–Zr–Ni-based alloys Absorpcija vodika v zlitinah Ti–Zr–Ni I. [kulj, A. Kocjan, P. J. McGuiness, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 279 Microstructural evaluation of rapidly solidified Al–7Cr melt spun ribbons Ovrednotenje mikrostrukture hitrostrjenih trakov Al-7Cr P. Jur~i, M. Dománková, M. Hudáková, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283 The influence of different waste additions to clay-product mixtures Vpliv razli~nih odpadkov na izhodno surovino za proizvodnjo ope~nih izdelkov V. Ducman, T. Kopar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289 Electrochemical and mechanical properties of cobalt-chromium dental alloy joints Elektrokemijske in mehanske lastnosti razli~nih spojev stelitne dentalne zlitine R. Zupan~i~, A. Legat, N. Funduk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Development of microstructure of steel for thermal power generation Razvoj mikrostrukture jekel za termi~no generacijo energije Kvackaj T., Kuskulic T., Fujda M., Pokorny I., Weiss M., Bevilaqua T. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 301 LETNO KAZALO – INDEX 324 MATERIALI IN TEHNOLOGIJE 41 (2007) 6 MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY AVTORSKO KAZALO / AUTHOR INDEX LETNIK / VOLUME 41, 2007, A–@ A Abbadi A. 157 Adolf Z. 185 Ambro`i~ M. 179 An`el I. 53, 271 Arh B. 203 Arsi} M. 173 Arzen{ek B. 29 Ayadi S. 47 Azari Z. 157 B Ba`an J. 99, 139 Balcar M. 139 Bele M. 103 Bevilaqua T. 301 Bihar @. 265 Bilu{i} A. 265 Bon~ina T. 243, 271 Bonfoh N. 77 Bouaouadja N. 157 Burzi} Z. 163 C Cejp J. 231 Celin R. 151 Cizelj L. 227 ] ]ur~ija D. 21 ^ ^ampelj S. 103 D Dobina N. 67 Dolin{ek J. 265 Dománková M. 131, 283 Drofenik M. 103 Ducman V. 289 \ \urkovi} S. 189 F Fila P. 139 Fujda M. 301 Funduk N. 295 G Getsov L. B. 67, 257 Gilgert J. 157 Gjonaj M. 249 Gliha V. 125 Grabulov V. 163 Grga~ P. 81 Gubeljak N. 35, 41 Gusel L. 53 Gusha L. 249 H Hadj-Taïeb E. 47 Hen~ - Bartoli} V. 243 Hnilica F. 231 Hudáková M. 81, 283 Husar I. 185 Huta P. 135 I Ivanovi} 189 M. J Jakli~ A. 119 Jamar J. 13 Jamar N. 13 Jamnik J. 103 Jenko M. 7, 199 Jevnikar P. 237 Jodin P. 29 Jozi} D. 91 Jur~i P. 231, 283 K Kabelevskiy M. G. 257 Keskes B. 157 Klob~ar D. 167 Kmeti~ D. 111, 151 Kne`evi} 189 M. Kocjan A. 279 Kopar T. 289 Kosec L. 8, 53 Kosma~ T. 237 Kozak D. 41 Krgovi} M. 189 Kumer B. 119 Kunc R. 73 Kuskulic T. 301 Kusý M. 81 Kvackaj T. 301 L Labor M. 95 Legat A. 295 Leonti} B. 265 Lipinski P. 77 Lufi S. 249 Lukatela J. 265 M Ma~ek J. 85 Makovec D. 103 Mamuzi} I. 21 Marin{ek M. 85 Markoli B. 271 Marolt T. 119 Marstijepovi} N. 189 Martinac V. 95 Martínek L. 139 McGuiness P. J. 279 Mejac M. 35 Menger Y. 157 Miro{evi}-Anzulovi} M. 95 O Oblak ^. 237 Oru~ 193 M. P Peter M. 131 Petric N. 95 Pihura 193 D. Pipi} D. 243 Pluvinage G. 47 Poharc M. 53 Pokorny I. 301 Prebil I. 73 Predan J. 35 R Radovanovi} I. 91 Razpotnik T. 85 Risteski I. B. 213 LETNO KAZALO – INDEX MATERIALI IN TEHNOLOGIJE 41 (2007) 6 325 Roman M. 131 Rudolf R. 53 Rybnikov A. 67 S Sedlická V. 81 Sedmak A. 163, 173 Sedmak S. 163, 173 Seitl S. 135 Sen~i~ B. 29 Simonovski I. 227 Smiljani} I. 265 Smontara A. 265 Stránský K. 99 Suchánek P. 185 [ [etina J. 145 [kulj I. 279 [u{tar{i~ B. 29, 279, 283 T Tehovnik F. 111, 203 Tiem S. 77 Tu{ek J. 167 V Vidovi~ K. 179 Vode F. 119 Vodopivec F. 3, 6, 9, 111 Vojvodi~ Tuma J. 41, 111, 173 Vuherer T. 125 W Weiss M. 301 Z Zejak R. 189 Zeli} J. 91 Zupan K. 85 Zupan~i~ R. 295 Zupani~ F. 243, 271 @ @elezný R. 139 @erovnik A. 73 @vokelj M. 73 LETNO KAZALO – INDEX 326 MATERIALI IN TEHNOLOGIJE 41 (2007) 6 MATERIALI IN TEHNOLOGIJE / MATERIALS AND TECHNOLOGY VSEBINSKO KAZALO / SUBJECT INDEX LETNIK / VOLUME 41, 2007 Jubileji – Jubilee Materiali in tehnologije – 40 let Materials and technology – 40 years F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Prvi urednik revije – Jo`a Arh F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6 Laudation in honour of professor dr. Franc Vodopivec on the occasion of his 75th birthday M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 75 let Franca Vodopivca L. Kosec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Marin Gabrov{ek – ob jubileju F. Vodopivec . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 [estdeset let prof. dr. Vasilija Pre{erna Laudation in honour of Professor Dr. Vasilij Pre{ern on the occasion of his 60th birthday M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 199 Kovinski materiali – Metallic materials Theoretical calculation of the lubrication-layer thickness during metal drawing Teoreti~ni izra~un debeline plasti maziva pri vle~enju kovin D. ]ur~ija, I. Mamuzi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 The notch effect on the fatigue strength of 51CrV4Mo spring steel VPliv zareze na trajno nihajno trdnost vzmetnega jekla 51CrV4Mo B. [u{tar{i~, B. Sen~i~, B. Arzen{ek, P. Jodin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29 An integrity analysis of washing-machine holders Analiza celovitosti nosilca kadi v pralnem stroju N. Gubeljak, M. Mejac, J. Predan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35 The effect of a material’s heterogeneity on the stress and strain distribution in the vicinity of a crack front Vpliv heterogenosti materiala na porazdelitev napetosti in deformacije v bli`ini konice razpoke D. Kozak, N. Gubeljak, J. Vojvodi~ Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41 The numerical solution of strain wave propagation in elastical helical spring Numeri~na re{itev propagacije deformacijskega vala v elasti~ni spiralni vzmeti S. Ayadi, E. Hadj-Taïeb, G. Pluvinage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47 [tudij notranje oksidacije v naoglji~enih hitrostrjenih trakovih Cu The study of the internal oxidation in internally carbonised Cu ribbons R. Rudolf, L. Kosec, I. An`el, L. Gusel, M. Poharc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53 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 LETNO KAZALO – INDEX MATERIALI IN TEHNOLOGIJE 41 (2007) 6 327 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 Structural steels with micrometer grain size: a survey Konstrukcijska jekla z mikrometrskimi kristalnimi zrni: pregled F. Vodopivec, D. Kmeti~, F. Tehovnik, J. Vojvodi~-Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111 The implementation of an online mathematical model of billet reheating in an OFU furnace Implementacija simulacijskega modela za spremljanje ogrevanja gredic v OFU-pe~i A. Jakli~, F. Vode, T. Marolt, B. Kumer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119 The behaviour of coarse-grain HAZ steel with small defects during cyclic loading Vedenje jekla grobozrnatega TVP z napakami pri cikli~ni obremenitvi V. Gliha, T. Vuherer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 125 The effect of cold work on the sensitisation of austenitic stainless steels Vpliv hladne deformacije na pove~anje ob~utljivosti nerjavnih jekel M. Dománková, M. Peter, M. Roman . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131 Fatigue-crack propagation near a threshold region in the framework of two-parameter fracture mechanics Dvoparametrska lomno mehanska analiza hitrosti utrujenostne razpoke blizu praga propagacije S. Seitl, P. Huta . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 135 Modelling of the solidification process and the chemical heterogeneity of a 26NiCrMoV115 steel ingot Modeliranje procesa strjevanja in kemi~ne heterogenosti ingota iz jekla 26NiCrMoV115 M. Balcar, R. @elezný, L. Martínek, P. Fila, J. Ba`an . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 A wet-steam pipeline fracture Prelom cevovoda za vla`no paro R. Celin, D. Kmeti~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151 A fatigue characterization of honeycomb sandwich panels with a defect Utrujenostna karakterizacija satastih sendvi~nih panelov z napako B. Keskes, Y. Menger, A. Abbadi, J. Gilgert, N. Bouaouadja, Z. Azari . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 Fatigue properties of a high-strength-steel welded joint Utrujenostne lastnosti zvara visokotrdnega jekla Z. Burzi}, V. Grabulov, S. Sedmak, A. Sedmak . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 163 Mehanske lastnosti zvara iz jekla maraging po izlo~evalnem `arjenju Mechanical properties of maraging steel welds after aging heat treatment D. Klob~ar, J. Tu{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 167 An experimental verification of numerical models for the fracture and fatigue of welded structures Eksperimentalna verifikacija numeri^nih modelov za prelom in utrujenost zvarjenih struktur S. Sedmak, A. Sedmak, M. Arsi}, J. Vojvodi~ Tuma . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 173 Investigation of the influence of the melt slag regime in a ladle furnace on the cleanliness of the steel Raziskava vpliva re`ima `lindre v ponov~ni pe~i na ~istost jekla Z. Adolf, I. Husar, P. Suchánek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 185 The application of spheroidal graphite cast iron in Bosnia and Herzegovina Uporaba nodularne grafitne litine v Bosni in Hercegovini D. Pihura, M. Oru~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 193 The oxidation and reduction of chromium during the elaboration of stainless steels in an electric arc furnace Oksidacija in redukcija kroma iz `lindre med izdelavo nerjavnih jekel v elektrooblo~ni pe~i B. Arh, F. Tehovnik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 203 A new topology for the trajectories of the meniscus during continuous steel casting Nova topologija trajektorij meniskusa pri neprekinjenem litju jekla I. B. Risteski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213 Multiscale modelling of short cracks in random polycrystalline aggregates Ve~nivojsko modeliranje kratkih razpok v naklju~nih ve~kristalnih skupkih L. Cizelj, I. Simonovski . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227 LETNO KAZALO – INDEX 328 MATERIALI IN TEHNOLOGIJE 41 (2007) 6 Changes to the fracture behaviour of medium-alloyed ledeburitic tool steel after plasma nitriding Spremembe v na~inu preloma srednje legiranega ledeburitnega jekla zaradi plazemskega nitriranja P. Jur~i, F. Hnilica, J. Cejp . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 231 Povr{ina zlitine Cu-Sn-Zn-Pb po obsevanju z ultravijoli~nim du{ikovim laserjem Surface of Cu-Sn-Zn-Pb alloy irradiated with ultraviolet nitrogen laser F. Zupani~, T. Bon~ina, D. Pipi}, V. Hen~ - Bartoli} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243 A preliminary S-N curve for the typical stiffened-plate panels of shipbuilding structures Preliminarna krivulja S-N za toge plo{~ate panele za ladjedelni{ke strukture L. Gusha, S. Lufi, M. Gjonaj . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 249 A new version of the theory of ductility and creep under cyclic loading Nova verzija teorije o duktilnosti in lezenju pri cikli~ni obremenitvi L. B. Getsov, M. G. Kabelevskiy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 257 Thermoelectrical properties of a monocrystalline Al64Cu23Fe13 quasicrystal Termoelektri~ne lastnosti monokristalnega kvazikristala Al64Cu23Fe13 I. Smiljani}, A. Bilu{i}, @. Bihar, J. Lukatela, B. Leonti}, J. Dolin{ek3, A. Smontara . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 265 Faze v kvazikristalni zlitini Al64,4Cu22,5Fe13,1 Phases in a quasicrystalline alloy Al64,4Cu23,5Fe13,1 T. Bon~ina, B. Markoli, I. An`el, F. Zupani~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 271 Hydrogen absorption by Ti–Zr–Ni-based alloys Absorpcija vodika v zlitinah Ti–Zr–Ni I. [kulj, A. Kocjan, P. J. McGuiness, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 279 Microstructural evaluation of rapidly solidified Al–7Cr melt spun ribbons Ovrednotenje mikrostrukture hitrostrjenih trakov Al-7Cr P. Jur~i, M. Dománková, M. Hudáková, B. [u{tar{i~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283 Electrochemical and mechanical properties of cobalt-chromium dental alloy joints Elektrokemijske in mehanske lastnosti razli~nih spojev stelitne dentalne zlitine R. Zupan~i~, A. Legat, N. Funduk . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 295 Development of microstructure of steel for thermal power generation Razvoj mikrostrukture jekel za termi~no generacijo energije Kvackaj T., Kuskulic T., Fujda M., Pokorny I., Weiss M., Bevilaqua T. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 301 Anorganski materiali – Inorganic materials 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 The fracture and fatigue of surface-treated tetragonal zirconia (Y-TZP) dental ceramics Prelom in utrujenost povr{insko obdelane tetragonalne (Y-TZP) dentalne keramike T. Kosma~, ^. Oblak, P. Jevnikar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 237 Vakuumska tehnika – Vacuum technique Frekven~na odvisnost rezidualnega trenja viskoznostnega vakuumskega merilnika z lebde~o kroglico Frequency dependence of spinning rotor gauge residual drag J. [etina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145 Gradbeni materiali – Materials in civil engineering Izra~un parametrov Weibullove porazdelitve za oceno upogibne trdnosti valovitih stre{nih plo{~ Computation of the parameters of the Weibull distribution for estimating the bending strength of corrugated roofing sheets M. Ambro`i~, K. Vidovi~ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 179 The influence of illite-kaolinite clays’ mineral content on the products’ shrinkage during drying and firing Vpliv vsebnosti glin ilinit-kaolinit na kr~enje pri su{enju in `ganju M. Krgovi}, N. Marstijepovi}, M. Ivanovi}, R. Zejak, M. Kne`evi}, S. \urkovi} . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189 The influence of different waste additions to clay-product mixtures Vpliv razli~nih odpadkov na izhodno surovino za proizvodnjo ope~nih izdelkov V. Ducman, T. Kopar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289 LETNO KAZALO – INDEX MATERIALI IN TEHNOLOGIJE 41 (2007) 6 329 Informatika – Informatics Zgodovina serijske publikacije Materiali in tehnologije / Materials and technology Historical overview of the scientific journal Materiali in tehnologije / Materials and technology N. Jamar, J. Jamar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 LETNO KAZALO – INDEX 330 MATERIALI IN TEHNOLOGIJE 41 (2007) 6