VSEBINA – CONTENTS PREGLEDNI ^LANKI – REVIEW ARTICLES Applications of focused ion beam in material science Uporaba fokusiranega ionskega curka v znanosti materialov L. Repetto, G. Firpo, U. Valbusa . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143 Materials and Technology: historical overview Materiali in tehnologije: zgodovinski pregled N. Jamar, J. Jamar . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151 IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES Low energy-high flux nitridation of metal alloys: mechanisms, microstructures and high temperature oxidation behaviour Nitriranje kovinskih zlitin s fluksom z majhno energijo in veliko gostoto: mehanizmi, mikrostrukture in visokotemperaturno oksidacijsko vedenje F. Pedraza . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 Numerical and experimental analyses of the delamination of cross-ply laminates Numeri~na in eksperimentalna analiza delaminacije v kri`nih plo{~atih laminatih R. Zem~ík, V. La{ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171 Application of the theory of physical similarity for the filtration of metallic melts Uporaba teorije fizikalne podobnosti za opis filtriranja kovinske taline K. Stránský, J. Ba`an, J. Dobrovská, M. Balcar, P. Fila, L. Martínek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175 Priprava nanokompozita za biomedicinske aplikacije Preparation of nano-composites for biomedical applications S. ^ampelj, D. Makovec, L. [krlep, M. Drofenik . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 179 STROKOVNI ^LANKI – PROFESSIONAL ARTICLES The development of a chill mould for tool steels using numerical modelling Razvoj kokile za orodna jekla z uporabo numeri~nega modeliranja M. Balcar, R. @elezný, L. Sochor, P. Fila, L. Martínek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183 1. MEDNARODNA KONFERENCA O MATERIALIH IN TEHNOLOGIJAH POD POKROVITELJSTVOM IUVSTA IN FEMS 13. – 15. oktober, 2008, Portoro`, Slovenija 1st INTERNATIONAL CONFERENCE ON MATERIALS AND TECHNOLOGY SPONSORED BY IUVSTA AND FEMS 13–15 October, 2008, Portoro`, Slovenia . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 189 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 42(4)141–188(2008) MATER. TEHNOL. LETNIK VOLUME 42 [TEV. NO. 4 STR. P. 141–188 LJUBLJANA SLOVENIJA JUL.–AUG. 2008 L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE UPORABA FOKUSIRANEGA IONSKEGA CURKA V ZNANOSTI MATERIALOV Luca Repetto, Giuseppe Firpo, Ugo Valbusa Nanomed Lab, Centro di Biotecnologie Avanzate – CBA, L.go R. Benzi, 10, 16132 Genova, Italy and Università of Genova, Physics Department, Via Dodecaneso, 33 16146 Genova, Italy valbusafisica.unige.it Prejem rokopisa – received: 2007-10-08; sprejem za objavo – accepted for publication: 2008-06-24 The focused ion beam (FIB) microscope is a tool that has a widespread use in the field of material science because it is able to micromachining with high resolution imaging thus therefore enhancing a broad range of both fundamental and technological applications in material science. The FIB is based on a beam of Ga ions which sputter the sample enabling precise machining at the nanometer/micrometer scale. The FIB instruments received particular attentions in the 1980s when the semiconductor industry used it as offline equipment for mask or circuit repair, but only in the 1990s the FIB was used in research laboratory. Nowadays there are commercial instruments (Dual Beam FIB / SEM) that integrate the precision cross section power of a FIB with the high resolution imaging of an SEM creating a powerful cross section and imaging tool. The combined SEM capability allows for real time monitoring of the FIB cuts with a higher resolution. Key words: focused ion beam, ion source, dual beam instruments, ion-solid interactions, applications Mikroskop s fokusiranim snopom ionov (FIB) je orodje s {iroko uporabnostjo v znanosti materialov, ker omogo~a mikroobdelavo in opazovanje z veliko lo~ljivostjo in odpira {iroko podro~je temeljnih raziskav materialov in tehnolo{kih aplikacij. Podlaga FIB je curek Ga-ionov, s katerim se obstreljuje tar~a, kar omogo~a mikroobdelavo pri redu velikosti nanometer-mikrometer. Naprava FIB je vzbudila posebno pozornost v 80-ih letih, ko so jo uporabljali v industriji polprevodnikov za popravilo mask in tokokrogov, v 90-ih letih pa se je raz{irila tudi v raziskovalne laboratorije. Na voljo so komercialne naprave (Dual beam FIB), ki zdru`ujejo natan~nost FIB z visoko lo~ljivostjo SEM, kar ustvari u~inkovito orodje za obdelavo in opazovanje. Kombinirana naprava omogo~a v realnem ~asu opazovano FIB-obdelavo z veliko lo~ljivostjo. Klju~ne besede: izvor ionov, naprava z dvojnim curkom, interakcija ion-trdna snov, uporaba 1 INTRODUCTION In his famous talk on nanotechnology "There’s Plenty of Room at the Bottom" held at the 1959 meeting of the American Physical Society at Caltech 1, Richard Feynman considered "the problem of manipulating and controlling things on a small scale". As an example of this kind of manipulations, he considered the task of writing "the entire 24 volumes of the Encyclopaedia Britannica on the head of a pin" and for the purpose he imagined a machine which could afford this task: "we can reverse the lenses of the electron microscope in order to demagnify as well as magnify. A source of ions, sent through the microscope lenses in reverse, could be focused to a very small spot". Even if Feynman describes an interaction of the ions with the sample that is different from the one we know to happen (he postulates a deposition while we know that the result is a removal of atoms), all the same his words sounds like a very rough but essential picture of the focused ion beam (FIB). In this paper we describe how this very intuitive idea has become a real instrument (in Figure 1 we reproduce the result of Feynman’s "experiment" with a modern FIB). In particular we will consider the state of the art of dual-beam instruments, where the FIB is coupled to a scanning electron microscope (SEM) to realize one of the most powerful tools available in the field of nanotechnology. In Section 2 we will give an account of the main elements which build-up the FIB, namely an ion source and a set of lenses and scanning coils used to produce a finely focused ion spot which can be rastered and Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 143 UDK 620.1.08:537.5 ISSN 1580-2949 Review article/Pregledni ~lanek MTAEC9, 42(4)143(2008) Figure 1: Feynman’s experiment: "High school competition". The "o" letters in the smallest box have an internal diameter less than 400 nm. Slika 1: Feynmannov eksperiment: "Tekmovanje visoke {ole". Notra- nji premer majhnih ~rk o v majhnem okvirju je manj{i od 400 nm. pointed in the desired position to form images of the sample or mill specified shapes on it. In Section 3 we briefly describe the complexity of ion-solid interactions which for example allow the processes of sputtering and secondary electron emission (plus a plenty of other physical effects) used in the instrument operation. Finally, in Section 4 a sample of typical application will be described. The aim of this paper is not to give an exhaustive description of the physics and the technology of focused ion beams, but rather to provide a fast introduction with mention to recent applications. For complete descrip- tions of the topic Refs. 2–5. 2 THE FOCUSED ION BEAM The most important characteristic usually required to a FIB is the capability to mill very precise shapes in samples. This operation is performed by focusing in a very small spot the ions and moving this spot to the position where this localized sputtering is required. This is achieved through an ion source and an ion column where a set of lenses and scan coils are housed. The performances that can be obtained in this "primary" task are first of all related to the spot size, the intensity and the stability of the beam. In principle the smallest and most intense is the ion spot, the most precise and fastest will be the job. In practice, limitations in the maximum removal rate that can be achieved may occur as effect of unwanted phenomena like redeposition 4, but this will be considered later. 2.1 The ion source Like in optical systems, the fundamental ingredient to get a small and intense spot is a bright source, where the brightness is defined 7 as the differential current intensity d2I emitted by the surface element dA of the source into the solid angle dΩ, i. e. B = d2I/dAdΩ. The most efficient and practical way to satisfy the requirement of a high brightness is through the so called Liquid Metal Ion Sources (LMIS). LMIS are nowadays the common choice for general purpose instruments, where no special demand exists for the ion species and we will limit our description to instruments were this solution is adopted. In a liquid metal ion source, a field emitter (typically a tungsten needle with a tip radius of a few microns) is connected to a reservoir containing a liquid metal (or alloy). A heating system is provided for the reservoir if the chosen metal is not liquid at the temperature of operation. The metal, which can flow to the emitter, need to be a wetting liquid. In this case, if a voltage (the extraction voltage) of the order of 1 kV is applied by a nearby electrode the liquid assumes a conical shape which is the equilibrium configuration under the com- peting forces produced by the surface tension and the electrostatic field. In an ideal situation, where no further effect would occur, it has been shown 8 that an ideal cone with a half angle of 49.3° (called a "Taylor cone") would be created. In practice, the liquid cone is pulled by the electrostatic field until its end radius R reaches a value small enough to have an electric field that can start field evaporation 9 in proximity of the tip. In this situation any further reduction in the tip radius is inhibited and it has been shown 10 that the cone apex takes a rounded shape with R  1–10 nm. According this scheme, ion sources of Ga, Au, Si, Pd, B, P, As, Cu, Ge, Fe, In, U, Be, Cs, Li , Pb, etc. have been produced with brightness B  106 A/cm2 sr. Particularly relevant is the implementation of the Gallium LMIS, which is currently the most common choice in commercial instruments. The reasons for this choice are manifold: the low melting temperature (29.8 °C) is certainly fundamental, since it simplify the system and prevent or reduce chemical and physical interactions between the liquid metal and the field emitter. Gallium atomic mass (69.72) has an "intermediate" value bet- ween light and heavy elements; this makes gallium ions suitable for efficient sputtering with a wide choice of substrates (the maximum energy transfer in a scattering event occurs when the target and the projectile have the same mass). Its physical properties like surface tension, vapour pressure and vacuum properties maximize the source lifetime. Overall gun performances have been demonstrated to be excellent with respect characteristics like angular intensity and energy spread. 2.2 The column Although the mechanism of field evaporation is possible only for a very small R, the effective source size results to be much larger due to space-charge effects. In practice a virtual source size of the order of 50–100 nm has been estimated 11. To reach a probe size on the sample of the order of 10 nm or less the image of the source need to be demagnified. This and other similar tasks are performed by the ion column. We will not enter into the details of this part of the instrument since the argument is mostly technical, but we will limit to a functional description of its components. After being generated in the gun, ions are accelerated down the column by a voltage typically chosen in the range 5–50 kV. The optical system therein contained is usually composed by two electrostatic lenses. The choice of electrostatic lenses rather than electromagnetic ones like in SEM columns, is due to the fact that ions are much massive than electrons. This implies that if accelerated at the same energy, ions are much slower and thus the magnetic part of the Lorentz force is by far weaker. The first element of a two lens system is the condenser whose function is to "collect" the ions and to set a suitable divergence for the beam before it passes L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE 144 Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 the forming aperture. The apertures shape the probe and reduce the current. In this way a range of different probe currents can be selected by changing the aperture (and eventually the condenser voltage). Typically the current intensity can be chosen from a few picoamps to a few tenths of nanoamps. The lens that follows is the objec- tive lens that focus the beam on the sample. Further elements contained in the column are a deflection/astigmatism unit and a fast beam blanker used for raster and mill operation. 2.3 Dual-beam instruments As we will describe in Sec. 3, the ion-solid interaction is a quite a complicated process and the intuitive concept of a nano-mill for the FIB is valid only in a very rough approximation. In practice, the result of even very simple operations like milling a square hole in a flat sample can give unexpected results as shown in Figure 2. For "unique" samples this can be a question of great concern for the operator which is responsible for analysis that cannot be repeated. The best solution to this problem has been to provide the way to have real-time visual inspection on how the work is advancing. The technical implemen- tation of this solution are dual-beam instruments, where a SEM can operate simultaneously with the FIB. In standard dual-beam instruments the electronic column enter vertically the high vacuum chamber where the sample is placed, while the FIB column axis has some angle with respect to the SEM, typically around 50°. In Figure 3 the interior of a dual-beam instrument is shown, where the vertical electron column and the 54° tilted ion column can be seen. The capabilities of this workstation can be further increased to make it become a complete nano-factory by equipping the FIB with a Gas Injections System (GIS) and micro-manipulators and eventually the SEM with an analytical tool like an EDS (Energy Dispersive X-Ray Spectroscopy) system or a WDS (Wavelength Dispersive X-Ray Spectroscopy) system 12. The GIS is a mechanical arm inside the sample chamber ending with capillaries connected by valves to external reservoirs where different gas species can be contained. By placing these capillaries very close to the sample and opening one of the valves, the pressure can be locally increased without perturbing excessively the high vacuum condition needed for FIB and SEM operation. The gas injected in the chamber adsorbs on the sample surface. The molecules in this weak bond state can be broken by the ion beam impinging on the surface. Depending on the adsorbed gas, different processes can occur. If the gas is a precursor of some element or chemical compound, a part of the molecule becomes chemically bond to the substrate while the other part is removed by the pumping system. In this way, precise local deposition of metals like Au, Pt, W etc., or insulators like SiO2 can be realized. Other precursor can instead be injected to enhance or make selective the ion etching process. Micro-manipulators inside the sample chamber can be used to move portions of the sample extracted by using the FIB. A typical task is the preparation of lamellae for Transmission Electron Microscopy (TEM). In other cases they can be used like electrical probes for conductivity measurement during the work session. Analytical tools provide a way to establish the chemical composition of the sample. In conjunction with the FIB capability to realize sections, these systems literally acquire a third dimension making it possible a volumetric mapping of chemical species in the sample. L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 145 Figure 3: Interior of the sample chamber of a dual beam instrument. A: SEM, B: FIB, C: sample stage, D: GIS. Slika 3: Notranjost komore za vzorec pri napravi z dvojnim curkom. A: SEM, B: FIB, C: mizica z vzorcem, D: GIS Figure 2: Effects of redeposition: the milling process had been set to produce vertical walls, but the final result shows a "mild" slope on the side milled first. Slika 2: U~inek redepozicije: proces obdelave je bil nastavljen za izdelavo pokon~nih sten, kon~ni rezultat pa je majhen naklon stene, ki je bila izdelana najprej. 3 ION-SOLID INTERACTIONS It has already been mentioned that the most important operation that can be realized by a FIB is a spatially selective removal of atoms from the sample. From a general point of view, this process is called sputtering and has been studied in detail in a funda- mental paper by Sigmund 13. One of the initial observations reported by Sigmund is that it is unlikely for the first collision of the ion to produce a sputtered atom, since the transferred momen- tum (at least for normal incidence) has a component in the direction entering the surface. On the contrary sputtering is one of the results of the so called "collision cascade" initiated by the ion entering the sample. This ion undergoes a series of collisions in the target and each atom which acquires sufficient kinetic energy by this collision can initiate a new series of collisions. The physics of the collision cascade is complex and not fully understood. Moreover it cannot uniquely be defined since the open scattering channels can be different for different energy of the primary beam. We will limit our considerations to a very schematic classifi- cation. A fundamental division of the processes that take place during a collision cascade is between elastic and inelastic events. Elastic processes are responsible for displacement of lattice atoms, defects generation (amorphization) and sputtering. Inelastic events generate secondary electrons, X-rays, photoluminescence and phonons (heat). From the point of view of the incident ion, each event causes a transfer of energy to the solid (a speed reduction) and a deviation from the direction taken after the previous collision. At end (when and if the initial ion kinetic energy is not sufficient to make it move freely in the solid) the ion can stop in the sample (implantation) if the random deviations occurred did not bring it again to the surface with enough energy to escape in the vacuum (backscattering). The energy losses can occur in "nuclear channels" with elastic events (for typical FIB energies mostly by screened Coulomb scattering) and in "electronic channels" where inelastic interaction with lattice electrons produce excitation and ionization. Aside from sputtering, among the processes described, secondary electron generation is the most important. By utilizing the SEM detectors is in fact possible to generate images also with the FIB. The resulting images can show contrast mechanism (like the contrast for the local crystallographic orien- tation) that are not present when an electron primary beam is used. On the other hand, this imaging mode suffer from a continuous modification of the sample due to the contemporaneous occurrence of sputtering events. The processes described above are fundamental in nature. The actual phenomenology can be sometimes unexpected for the presence of further mechanisms. We describe briefly three of these mechanisms that can pose some limitation or at least make more complex the process of milling. 3.1 Redeposition Atoms removed by sputtering or backscattered ions can actually fail to "escape" from the solid and instead they can deposit on the surface of the sample very close to the milling site. The effect is a degradation in the quality of the milling operation that can easily reach unacceptable levels as shown in Figure 2. This effect is more pronounced when trying to mill high aspect ratio features and when using high ion fluxes. Possible solution are the reduction of the current and/or the dwell time at the expenses of a longer process time. A different solution is the introduction through the GIS of a reactive gas like F, which binds to the sputtered atoms and facilitate the removal by the pumping system. 3.2 Channelling Channelling occurs in crystalline material and is an apparent inhomogeneity of the sputtering yield across a chemically homogenous surface (Figure 5). It is due to the lower atom density along low index directions, resulting in a lower probability for the incident ion to hit a target ion. 3.3 Auto-organization Off-normal incidence of the primary beam can make evident the effects of the instability generated by the dependence on the local curvature of the sputtering yield 14–18. Since the erosion rate in depressions is larger than on surface mounds, any surface deviation from flatness tend to be amplified. The presence of a competing smoothing force due to surface atom diffusion produces typical structures showing long-range correlations. Easy to observe are wavelike structures usually indicated as ripples. In conventional milling processes the generation of these structures does not have particular relevance as long as height corrugations of the order of 10 nm can be neglected. 4 APPLICATIONS The FIB has reached a relatively large diffusion thanks to its application in microelectronic industry started in the 1980s. The capability to remove or add atoms in selected sites with submicron precision makes the FIB an instrument which can hardly be replaced in applications like failure analysis, mask and integrated circuit (IC) repair. A classical example is the inspection of a buried IC, where conventional techniques employ- ing mechanical tools to reveal the hidden parts can introduce artefacts that cannot be tolerated if resolution in the nanometre range is required. L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE 146 Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 As the diffusion in industry increased, in the 1990s, FIB systems started to be acquired by research labs: one fundamental application became the preparation of TEM cross-section lamellae. In this case several advantages can be indicated with respect to traditional techniques like the use of a ultramicrotome. First of all FIB prepa- ration is site-specific: the lamella can be extracted from a selected location in the sample. This possibility is fundamental for structured samples as in the case of biological specimens. Then, in general, a fewer artefacts are introduced above all when dealing with samples showing a non-homogeneous hardness across the section or with very soft materials. A further well-established application of the FIB is direct three-dimensional micro and nano-machining. These capabilities have used for MEMS (Micro-Electro- Mechanical Systems) and photonics structures reali- zations both in processes where only the FIB is used for rapid prototyping and in situation where finer details are added to classical lithography works. Since all these topics are widely covered in the literature, in the remaining part of this paper we will consider relatively new applications like the production of solid-state nanopores, and the study of ion-induce self-organization processes. They can serve as examples for two different ways to go beyond established limits of state-of-the art FIB instruments. 4.1 Production of solid-state nanopores Nanopores produced in solid state membranes have been proposed as the key element for a new class of devices deputed to fast DNA sequencing or (in a more general case) characterization (see 19 and references therein contained). In the basic set-up, an insulating membrane where a hole with size in the 1–10 nm range has been drilled, separates in two part a reservoir containing an ionic solution. By placing an electrode in each part and by establishing a voltage bias between the two regions, a ionic current starts to flow across the pore. If DNA molecules are inserted in the negatively biased region, they will tend to go through the pore and during the translocation a variation in the ionic current is expected. In principle, for pores small enough, it should be possible associate the instantaneous current variation to the single base which is occupying the pore in that instant. Variant of this scheme have been proposed, but all of them require an insulating membrane with a pore that is comparable in size with single-stranded DNA, i. e. a few nanometres in diameter. Classical lithographic techniques do not have sufficient resolution to be used in nanopores production, and, in this case, even the FIB can not perform the task directly. It has been reported in several papers a minimum reproducible size for FIB drilled pores in typical Si3N4 or SiO2 membranes (100 nm thick) of about 30 nm. Techniques like nano-sculpting 20 and high energy electron irradiation 21–23 have been proposed and an effective shrinkage to the desired size has been demonstrated. In the first case, a SiO2 membrane, where pores with a diameter of 50 nm have been realized by using a FIB, is irradiated with a broad 3 keV Ar+ beam. Under this irradiation, a gradual closure occurs as can be deduced from real time measurement of the transmitted ion current and final size less than 1 nm can be achieved. The process is explained through a model where adatoms (surface-diffusing mobile species), created by the incident beam, diffuse to the pore and fill it in. In the second case pores of similar initial size are produced by electron-beam lithography and plasma etching in a SiO2 membrane. Subsequent exposure to 300 keV electrons in a TEM reduces the size in a controllable way, with a precision of 0.2 nm (resizing occurs while imaging and this allows to follow the process). The claimed mechanism is a fluidization of the SiO2 under L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 147 Figure 4: Sequence showing shrinkage of a solid state nanopore in Si3N4 under electron irradiation. In the last tab an ex-situ acquired TEM micrograph shows a final size smaller than 10 nm. Slika 4: Sekvence ka`ejo kr~enje nanopore v Si3N4 pri obsevanju z elektroni. Zadnji TEM-posnetek prikazuje, da je kon~na velikost manj{a od 10 nm. the high energy electron irradiation and a pore shrinkage under the action of surface tension. Recently a new techniques that can entirely be realized inside a dual-beam instrument has been proposed and the results appear to be interesting. In this case the pore size is reduced during SEM imaging and the shrinking rate can be monitored simply following the image evolution. The resolution of this process is limited by the SEM capabilities to values larger than 1 nm. In Figure 4 a size reduction sequence is shown. The final size is evaluated ex-situ in a TEM. 4.2 Nanostructuring by self-organization processes The self-organizations processes briefly described in Sec. 3 have been often indicated as a possible high throughput method for nanostructuring surfaces. Recently it has been shown that similar structures can be created also by using the FIB and moreover the long range correlation of this structures can be increased if they are produced on a surface where an ordered template has previously been milled. Although the high throughput characteristic of this technique is lost when the process is realized with a FIB (due to the relative low currents used with respect to broad-beam sources) nevertheless some advantages are obtained. First of all, the possibility to use higher local fluxes allows the exploration of regimes which hardly can be accessed with conventional guns. Then, the self-organization process can be induced with different condition in adjacent regions producing structures of higher complexity. Finally, dual-beam instruments offer the capability of following in a real-time mode the process. This can give the necessary feedback to overcome the lack of knowledge of all the variables entering the self-organization process that influences the capability to predict the final morphology of the surface. From the point of view of the FIB user, the advantage of this technique is in term of increased resolution in producing particular structures as shown in Figure 6. Moreover for particular structures the process time can be greatly reduced. 5 CONCLUSION We have reviewed the main characteristics of Focused Ion Beam technology with particular attention to dual-beam platforms. The applications that we have mentioned demonstrate the widespread diffusion that this kind of instruments have reached, diffusion that goes beyond the fields where the use was initiated, i. e. microelectronic industry. We have finally considered two new research fields where the FIBs can both play a fundamental role and find new ways to go beyond its current limits. Acknowledgment We acknowledge the Italian Ministry for Research (MUR) which funded the NANOMED Project and allowed the acquisition of a state-of-the-art dual-beam instrument. We also acknowledge Fondazione Carige for funding the cleanroom where the FIB/SEM is located. 6 REFERENCES 1 First published in Engineering and Science magazine, XXIII (1960) 5 Available from World Wide Web: http://www.its.caltech.edu/ feynman/plenty.html L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE 148 Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 Figure 6: Pillar in silicon produced by off-normal incidence with a FIB. In the inset a particular at higher magnification clearly shows that the dimensions of the pillars is beyond FIB resolution. Slika 6: Stolpci na siliciju, dose`eni z nenormalno vpadnostjo v FIB. Slika pri veliki pove~avi v okvirju prikazuje, da je dimenzija stolpcev pod lo~ljivostjo FIB. Figure 5: Ripple formation on a silicate bio-structure (spicule). The ion beam is orthogonal to the sample holder, but due to the cylindrical shape of the sample, its incidence angle increases with the radius. Ripple rotation for angles closer to grazing incidence appear, as expected from Bradley and Harper theory 14. Slika 5: Nastajanje brazd na silikatni biostrukturi (spicule). Curek ionov je pravokoten na nosilec vzorca, zaradi valjaste oblike vzorca pa vpadni kot raste s premerom. Pojavi se zasuk brazd za kot blizu drsnega, kot napoveduje Bradley-Harperova teorija. 2 J. Orloff, M. Utlaut, L. Swanson, "High Resolution Focused Ion Beams: FIB and Its Applications", Kluwer, New York 2003 3 J. Meingailis, J. Vac. Sci. Technol. B 5 (1987), 468 4 J. Orloff, Rev. Sci. Instrum. 64 (1993), 1105 5 Introduction to Focused Ion Beams – Instrumentation, Theory, Tech- niques and Practice, L. Giannuzzi, F. Stevie (eds), Springer, Boston 2005 6 T. Ishitani, T. Ohnishi, J. Vac. Sci. Technol. A 9 (1991), 3084 7 M Born, E. Wolf, Principles of Optics, Pergamon, Oxford 1991, 6th (corrected) ed. 8 G. I. Taylor, Proc. R. Sot. London A 280 (1964), 383 9 E. W. Müller Phys. Rev. 102 (1956), 618 10 D. R. Kingham, L. W. Swanson, Vacuum 34 (1984), 941 11 J. W. Ward, J. Vat. Sci. Technol. B 3 (1985), 207 12 J. Goldstein, D. Newbury, D. Joy, C. Lyman, P. Echlin, E. Lifshin, L. Sawyer, J. Michael, Scanning Electron Microscopy and X-Ray Microanalysis, Springer (2003), 3rd ed. 13 P. Sigmund, Phys. Rev. 184 (1969), 383 14 R. M. Bradley, J. M. E. Harper, J. Vat. Sci. Technol. A 6 (1988), 2390 15 J. D. Erlebacher, M. J. Aziz, E. Chason, M. Sinclair, J. Florio, Phys. Rev. Lett. 82 (1999), 2330 16 S. Habenicht, Phys. Rev. B 63 (2001), 125419 17 M. A. Makeev, R. Cuerno, A. L. Barabasi, Nucl. Instr and Meth. B 197 (2002), 185 18 U. Valbusa, C. Boragno, F. Buatier de Mongeot, J. Phys. Condens. Matter 14 (2002), 8153 19 C. Dekker, Nature Nanotechnology 2 (2007), 209 20 Li J., Stein D., McMullan C., Branton D., Aziz M. J., Golovchenko J., Nature 412 (2001), 166 21 A. J. Storm, J. H. Chen, X. S. Ling, H.W. Zandbergen, C. Dekker, Nature Materials, 2 (2003), 537 22 H. Chang, S. M. Iqbal, E. A. Stach, A. H. King, N. J. Zaluzec, R. Bashir, Appl. Phys. Lett. 88 (2006), 103109 23 W. M. Zhang, Y. G. Wang, J. Li, J. M. Xue, H. Ji, Q. Ouyang, J. Xu, Y. Zhang, Appl. Phys. Lett. 90 (2007), 163102 L. REPETTO ET AL.: APPLICATIONS OF FOCUSED ION BEAM IN MATERIAL SCIENCE Materiali in tehnologije / Materials and technology 42 (2008) 4, 143–149 149 N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW MATERIALI IN TEHNOLOGIJE: ZGODOVINSKI PREGLED Nina Jamar1, Jana Jamar2 1Lipce 10, 4273 Blejska Dobrava, Slovenia 2Uredni{tvo serijske publikacije Materiali in tehnologije, In{titut za kovinske materiale in tehnologije, Lepi pot 11, 1000 Ljubljana, Slovenia nina.jamartelemach.net Prejem rokopisa – received: 2008-01-10; sprejem za objavo – accepted for publication: 2008-07-14 This article describes the evolution of the scientific journal Materials and Technology (Materiali in tehnologije) from its predecessors, the Iron and Steel Journal (@elezarski zbornik) and Metals Alloys Technologies (Kovine zlitine tehnologije). We present the statistical data for the journal relating to the years 2000–2006. The article looks at the electronic form of the journal and the influence of online publishing on the recognition and citing of Materials and Technology around the world. We show how the process of publishing the scientific journal takes into account ISO 9000. The article concludes with a description of how Materials and Technology will develop in the future. Key words: scientific serial publication, Materials and Technology, historical overview, statistical data, electronic publishing, ISO 9000 ^lanek opisuje razvoj znanstvene serijske publikacije Materiali in tehnologije (Materials and Technology) od njenih predhodnikov, @elezarski zbornik (Iron and Steel Journal) in Kovine zlitine tehnologije (Metals Alloys Technologies), do dana{njih dni. Predstavljeni so statisti~ni podatki za obdobje od leta 2000 do leta 2006. Opisan je razvoj elektronske oblike serijske publikacije in vpliv elektronskega izdajanja na mednarodno prepoznavnost in citiranje serijske publikacije Materiali in tehnologije. Predstavljen je potek izdajateljske dejavnosti serijske publikacije Materiali in tehnologije po ISO 9000. ^lanek se kon~uje z opisom, kako naj bi se serijska publikacija Materiali in tehnologije razvijala v prihodnje. Klju~ne besede: znanstvena serijska publikacija, Materiali in tehnologije, zgodovinski pregled, statisti~ni podatki, elektronsko zalo`ni{tvo, ISO 9000 1 @ELEZARSKI ZBORNIK – IRON AND STEEL JOURNAL The periodic publication the Iron and Steel Journal (@elezarski Zbornik) (ISSN 0372-8633) started as a joint project of the Slovenian Ironworks company and the Metallurgical Institute Ljubljana, an independent institution and also the central R&D facility for the company. At that time, in decisions about the development of technology and products, theoretical knowledge was applied to an increasing extent in addition to experience, and at the management level it was considered useful to provide, for people working in R&D, the possibility to exchange ideas and R&D results. At almost the same time, the organisation of the first of annual conferences to provide a platform for oral presentations and discussions about R&D started. The journal was published from 1967 to 1991 as a quarterly, with only three numbers in the first volume. Published articles were classified according to the Universal Decimal Classification and later also using the ASM/SLA Classification. The abstracts were published in English and German, and from 1968 also in Russian. By 1975 the journal started to print authors’ abstracts in Slovenian, German, English and Russian. Starting with Volume 2, the Annual Chronological Index was published in the last number of each volume or in the first issue of the next volume. 2 KOVINE ZLITINE TEHNOLOGIJE – METALS ALLOYS TECHNOLOGIES In 1992 the Iron and Steel Journal (@elezarski zbornik) changed its name to Metals Alloys Techno- logies (Kovine Zlitine Tehnologije) (ISSN 1318-0010). The name was changed to better cover the topics of the published articles, which had broadened from topics related to iron and steel to other metallic alloys as well as inorganic materials, polymers, and materials that are used in vacuum technology. A similar evolution took place in some western countries, e.g., in Great Britain the name of the journal Acta metallurgica was changed to Acta materialia, and in Germany the name Zeitschrift für Metallkunde was changed to Zeitshcrift für Material- kunde. Four issues per volume of Metals Alloys Technologies were published up to 1995, and six issues per volume were published from 1996 on. The manuscripts were submitted for publication from authors from Slovenia and abroad. The Annual Index, which included the chronological, authors’ and subject index, was added to the last issue of the year. At the request of an industrial society, in 1997 a special number was published. The role of the publisher was assumed by the Institute of Metals and Technology, and several institutions and industrial companies were brought in as associate publishers to provide a broader base. The Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 151 UDK 050:669.1:6(497.4) ISSN 1580-2949 Review article/Pregledni ~lanek MTAEC9, 42(4)151(2008) associate publishers were the companies ACRONI Jese- nice, IMPOL Slovenska Bistrica, Slovenia steelworks, Metal Ravne, Talum Kidri~evo, the National Institute of Chemistry, Institute "Jo`ef Stefan", the Faculty of Mechanical Engineering, and the Slovenian Society of Tribology. The Ministry for Science and Technology of the Republic of Slovenia (now the Ministry of Higher Education, Science and Technology of the Republic of Slovenia) funded some of the publishing expenses. The content was original scientific and professional contri- butions, review articles and expanded texts based on communications presented at the annual Conference on Materials and Technology. All the articles were published according to international ISO standards (ISO 8, ISO 18, ISO 214, ISO 215, ISO 690, ISO 690-2, ISO 832, ISO 999, ISO 2145, ISO 3297, ISO 5122, ISO 8459/1-5) and in line with the Instructions of the Slovenian Research Agency (SIST ISO 4, SISI ISO 8, SIST ISO 215, SIST ISO 214, SIST ISO 18, SIST ISO 690-2, SISI ISO 999, SIST ISO 2145, SIST ISO 5122). These instructions required ISSN, standard terminology, international measures and units, abstract, keywords, the beginning of the article on the odd page, and also the data on the author and the periodical publication had to be presented on the same page. Titles, abstracts and keywords were, for all articles, published in Slovenian and English. Since 1998 the date of receipt of the manuscript and the date of acceptance of the article for publication have also been published. At the suggestion of the Ministry for Science and Technology of the Republic of Slovenia, in 1998 the editorship of the journal started to categorize and arrange the published articles according to the recommendations of the typology for guiding bibliographies and the entry of data in COBISS Slovenia (Cooperative Online Bibliographic Systems and Services). Between 1996 and 1999, 436 entries in COBISS for the articles were made for articles published in previous years. In bibliometric analysis and a comparison of the Iron and Steel Journal (@elezarski Zbornik) and Metals Alloys Technologies (Kovine Zlitine Tehnologije) 1996/97 several parameters were used for the compa- rison, for example: • the number of authors of each article and the number of institutions from which the authors came, • the breadth of the contents base, • the shift from a professional publication to a scientific periodical publication. The results of the analysis showed a significant improvement of the journal Metals Alloys Technologies (Kovine Zlitine Technologije) in terms of originality, quality and the presentation of topics, as well as the regularity of publishing. The improvement was due to three factors: better knowledge and experience of the domestic authors, the increasing share of authors from abroad, and the publishing of topics on a wide range of materials as well as phenomena and processes.1 3 MATERIALS AND TECHNOLOGY – MATERIALI IN TEHNOLOGIJE In 2000 a new name, Materials and Technology (Materiali in tehnologije) (ISSN 1580-2949), was given to the journal and a new editorial board was selected, although the editors were not changed. The content was broadened to cover articles dealing with topics from a wide range of materials, such as metals and alloys, polymers, inorganic and vacuum materials, their testing and characterisation, development, manufacturing, processing and application for engineering structures. Also, articles on new developments in composites were published. The journal became the leading periodical publication in the field of materials in Slovenia. The topics published justified the change of the journal’s name, and this change attracted authors from other scientific disciplines, especially solid-state physics. The future aims of the journal are to achieve a higher international recognition and an appropriate citation index. Only by a continuous improvement in quality, which is related to the striving for greater originality and quality of the published works and the attraction of new, younger authors, especially from foreign countries, will an appropriate indexing in the Science Citation Index be achieved. This will attract new authors, Slovenian and foreign, who are to a large extent currently submitting their papers for publication in foreign journals. Six issues were published per volume, and double issues have not been published in the past two years. The sixth and last issue in the volume also includes the Annual Index: chronological, authors and subject index. During the 35th year of publication of the journal Materials and Technology – Materiali in Tehnologije in 2001, a bibliometric-bibliographic comparison of the journals Materials and Technology (2000) and Materials Science and Technology (2000) was done. It should be emphasized that Materials science and Technology is an international journal with the impact factor. The results have shown that there are no significant differences between the journals. The only noticeable difference is in the number of cited (and listed) sources. But when we are talking about the sources of references and about the age of the cited sources there are no significant differences. The finding that there are no significant differences between journals opens a lot of questions concerning international and local journals.2 The articles from periodical publication Materiali in tehnologije / Materials and Technology were till 2007/1 indexed in nine international secondary publications and databases: • Metals Abstracts, Engineered Materials Abstracts, Business Alert Abstracts (Steels, Nonferrous, Poly- mers, Ceramics, Composites), Chemical Abstracts, Aluminium Industry Abstracts, Referativnyj `urnal Metallurgija, Metadex, Inside Conferences, and DOMA, and since 2007/1 also in: N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW 152 Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 • DOAJ (Directory of Open Access Journals), GOOGLE SCHOLAR and SCIRUS. We are pleased to inform that Materiali in Tehno- logije has been selected for coverage in Thomson Reuters products and custom information services. Beginning with vol. 41 (1) 2007, this publication is indexed and abstracted in the following: – Science Citation Index Expanded (also known as SciSearch®) – Materials Science Citation Index – JournalCitations Reports / Science Edition In the future Materiali in Tehnologije may be evaluated and included in additional Thomson Reuters products and information services to meet the needs of the scientific and scholarly research community. The process of publishing Materials and Technology is in accordance with ISO 9000 (Diagram 1). Table 1: Typology of articles (1967–2006) Tabela 1: Tipologija ~lankov (1967–2006) Typology of articles Number % Review scientific articles 121 6.66 Original scientific articles 984 54.19 Professional articles 600 33.03 Technical news 86 4.74 Other 25 1.38 Together 1816 100.00 During the 40th year of publication, in 2007, a Special Issue of the journal Materials and Technology – Mate- riali in Tehnologije "The Bibliography of the articles 1967–2006" was published. The post office issued a stamp "40 years of publish- ing of the periodical publication Materials and Techno- logy" to mark the anniversary of the journal. The anniversary was also marked by the entering in COBISS (www.cobiss.si) all articles (1816) published in the journal and its predecessors in the years 1967 to 2007. 3.1 Statistical data on Materials and Technology – Materiali in Tehnologije: 2000–2006 Table 2: Number of issues, pages and articles Tabela 2: Fizi~ni obseg Issue Pages Articles 2000 / 1–6 452 79 2001 / 1–6 500 72 2002 / 1–6 492 71 2003 / 1–6 420 67 2004 / 1–6 + special number 491 67 2005 / 1-6 334 30 2006 / 1-6 354 44 Total 3043 430 Table 3: Share of scientific versus professional articles Tabela 3: Razmerje med objavljenimi znanstvenimi in strokovnimi ~lanki Year Professional % Scientific % 2000 38.00 62.00 2001 33.00 67.00 2002 21.00 79.00 2003 36.00 64.00 2004 28.35 71.65 2005 26.70 73.30 2006 13.65 86.35 Table 4: Language of publishing Tabela 4: Jezik ~lankov Year Slovenian % English % 2000 84.80 15.20 2001 72.20 27.80 2002 69.00 31.00 2003 58.00 42.00 2004 53.75 46.25 2005 31.70 68.30 2006 20.45 79.55 Table 5: Authors Tabela 5: Analiza po avtorjih Year All authors Articles Authors /article 2000 222 79 3 2001 222 72 3 2002 216 71 3 2003 194 67 3 2004 243 67 4 2005 109 30 4 2006 147 44 3 Table 6: Nationality of authors Tabela 6: Mednarodnost avtorjev Year Slovenia % Foreign countries % 2000 85.10 14.90 2001 85.10 14.80 2002 80.10 19.90 2003 73.20 26.80 2004 62.15 37.85 2005 69.70 30.30 2006 52.40 47.60 Table 7: Subject area Tabela 7: Vsebinski pregled ~lankov Scientific field 2000 2001 2002 2003 2004 2005 2006 Metallic materials 51.90 68.10 71.85 67.15 71.70 70.00 68.20 Inorganic materials 31.60 11.10 12.65 7.45 11.95 23.30 9.10 Vacuum technology 8.90 12.50 7.00 10.45 7.45 0.00 4.55 Polymers 6.30 6.90 7.05 3.00 3.00 3.35 13.60 Building materials 0.00 0.00 0.00 7.45 5.95 3.35 4.55 Information science 1.26 0.00 1.40 0.00 0.00 0.00 0.00 Research policy 0.00 1.40 0.00 0.00 0.00 0.00 0.00 Research and development 0.00 0.00 0.00 1.50 0.00 0.00 0.00 Standardization 0.00 0.00 0.00 1.50 0.00 0.00 0.00 Methodology 0.00 0.00 0.00 1.50 0.00 0.00 0.00 N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 153 N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW 154 Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 Diagram 1: The process of publishing Materials and Technology in accordance with ISO 9000 Diagram 1: Potek izdajateljske dejavnosti serijske publikacije Materiali in tehnologije po ISO 9000 Table 8: References quoted Tabela 8: Citiranje literature Number of articles with references quoted (%) Number of refe- rences 0 1–4 5–9 10–14 15–19 20–29 30 andmore Year 2000 1.27 15.19 39.24 26.58 8.86 7.59 1.27 2001 0.00 13.88 40.28 23.61 6.95 12.50 2.78 2002 0.00 11.30 33.80 28.15 12.70 11.25 2.80 2003 1.50 14.90 41.80 26.90 3.00 5.95 5.95 2004 1.50 10.45 43.25 22.35 14.95 6.00 1.50 2005 0.00 3.35 50.00 16.65 13.35 6.65 10.00 2006 0.00 18.18 34.10 22.72 13.63 4.54 6.81 Table 9: Average citations per article Tabela 9: Povpre~no {tevilo citatov na ~lanek Year Citations Articles Articles without citations Citations / article 2000 797 79 1 10.10 2001 807 72 0 11.20 2002 848 71 0 11.95 2003 775 67 1 11.55 2004 734 67 0 10.95 2005 530 30 0 18.30 2006 576 44 0 13.10 Based on the statistical data, the following findings should be emphasized (Tables 2–9): • the number of articles decreased from 79 articles in 2000 to 44 articles in 2006. This is in part due to the decrease in the R&D activity after the independence of Slovenia and in part due to the devaluation of articles published in Slovenian journals, • the number of scientific articles increased in relation to professional articles from 62 % in 2000 to 86 % in 2006, an indication of the improved average origi- nality and scientific value of the presented topics, • the number of articles published in English has increased very considerably, from 15.20 % in 2000 to 79.55 % in 2006, as a direct consequence of the official underrating of the use of Slovenian for publishing articles on topics related to technical aspects, technology and natural sciences, • the share of international authors increased very sub- stantially, from 14.90 % in 2000 to 47.60 % in 2006, • the average number of references quoted per article published has increased slightly, from 10.10 in 2000 to 13.10 in 2006. 3.2 Access in electronic form In electronic form the journals Metals Alloys Technology – Kovine Zlitine Tehnologije and Materials and Technology – Materiali in Tehnologije were initially accessible at the following URL: http://www.imt.si/ materiali-tehnologije (ISSN 1580-3414). The idea of the digitalisation of the journal Metals Alloys Technologies came in 1995. The decision was made for the journal to have a web site for the basic presentation of the journal. The web site was accessed at http://www.ctk.uni-lj.si/kovine/. In the next two years, 1997 and 1998, the newly developed HTML 3.2 and HTML 4 allowed new solutions, while Microsoft Office 97 made possible the very simple transformation of documents from the Office environment to hypertext. As a result, the idea of the electronic publishing of articles in full-text form was revived. The growth of electronic serials with access to the articles in full-text form allowed a comparison between different file types in which the full texts were accessible at web sites. The most interesting for the comparison were the journals on natural sciences and technology because of them having solutions for the graphical elements in the articles, i.e., graphs and figures, and for other difficult text parts, like mathematical derivations and chemical formulae. The web site of the journal was supplemented with a simple search tool for searching on the basis of indexes and hypertext assemblies of articles (data on authors, abstracts in Slovenian and English, key words in Slo- venian and English). With the change of the title Metals Alloys Technologies to Materials and Technology, the web site was given a new form and content. The new URL of the electronic version of the journal is http://www.imt.si/ materiali-tehnologije. The web site presented a sort of web portal for the field of natural sciences and technology. The aim of the web site was to achieve the highest data accessibility for a wide circle of users. The web site also enables an interactive attitude between readers and the editorship. With the monitoring of new technologies in electronic publishing as well as with further education, more useful connections with electronic forums and with notifications for the users about the new ways of accessing the data from this field of science (monitoring of the development and accessibi- lity of new databases and standards, notifications about conferences, the possibility for further education, and the monitoring of legislation) it is possible that the web site will be friendly and useful for researchers, students and other.3 The advantages of publishing the journal in electro- nic form are: • ordering of articles from a simple database allows subject enquiries over the content of the articles, • simplification of the contacts between the authors and the editorship, • permanent and simple access to the full text of the articles regardless of the place and time, • the electronic form can be accessed before the printing of the journal, • wider impact and visibility in the international community, N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 155 • different search options (by indexes, abstracts, keywords), • possibilities for connections between quoted and other related articles, • quick accessibility of authors to notifications of the editorship: to instructions for submission of manu- scripts and other notifications, • access and downloading of the full text of the article, • traceability of the article (history of the article at one point), • the electronic form is cheaper than the printed form. 3.3 Materials and Technology – the future • Materials and Technology – Materiali in Tehnologije is the leading periodical publication in the field of materials and composites in Slovenia, and strengthening this position should be the main goal for the Editorial Board in the future. • The number of articles in English should increase further. However, the activity of maintaining Slo- venian as a developed cultural language should not be neglected. For this reason, as the leading periodical publication, in the future articles in Slovenian should also be published. A compromise should be found between the majority of articles published in English, mostly original scientific works, and some of the articles published in Slo- venian, probably mostly with authors from industry. The bilingual publishing of all articles depends on the extent of funding. • The journal covers the fields where figures, micro- structures and diagrams are very important, for this reason the editors should conserve the quality of printing at the current level or even improve it. • The editors should attract more distinguished authors especially from English-speaking countries. • The number of citations of the journal in prestigious periodical publications should be increased. • The electronic form of the periodical publication of the journal must maintain its level of quality. • The electronic version of the journal should exist as a web portal, where the user, in addition to the articles in full-text form, should find information about other services from the scientific fields covered (the cata- logue of web connections, news, interesting infor- mation and data of interest for scientists). • The journal should actively take part in the develop- ment of a potential information service for Slovenian professional and scientific periodical publication. 4 REFERENCES 1 N. Jamar, M. Ba{, P. Ju`ni~, Mater. Tehnol., 34 (2000) 1/2, 7–14 2 P. Ju`ni~, N. Jamar, Mater. Tehnol., 36 (2002) 3/4, 169–177 3 M. Pu{nik, The role of special libraries by the acceleration of the so- cial and economic development. Construction of library collections: the acquisition and secretion of the material, Ljubljana, 2000, 163–172 N. JAMAR, J. JAMAR: MATERIALS AND TECHNOLOGY: HISTORICAL OVERVIEW 156 Materiali in tehnologije / Materials and technology 42 (2008) 4, 151–156 F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS: MECHANISMS, MICROSTRUCTURES AND HIGH TEMPERATURE OXIDATION BEHAVIOUR NITRIRANJE KOVINSKIH ZLITIN S FLUKSOM Z MAJHNO ENERGIJO IN VELIKO GOSTOTO: MEHANIZMI, MIKROSTRUKTURE IN VISOKOTEMPERATURNO OKSIDACIJSKO VEDENJE Fernando Pedraza Université de La Rochelle. Laboratoire d’Etudes des Matériaux en Milieux Agressifs (LEMMA, EA 3167). Avenue Michel Crépeau, 17042 La Rochelle cedex 01, FRANCE fpedrazauniv-lr.fr Prejem rokopisa – received: 2007-09-17; sprejem za objavo – accepted for publication: 2008-06-07 Nitridation is typically carried out to improve wear and erosion of different metal and alloy substrates. In the case of "stainless" alloys, the nitridation temperature needs to be lowered to avoid the precipitation of CrN that would reduce the overall corrosion resistance. Low energy – high flux nitridation allows to nitride relatively thick layers in short times at low temperatures depending on the substrate crystal structure and chemical composition as shown for pure Ni, a Ni-20Cr model alloy, a conventional AISI 304L stainless steel and an ODS FeAl intermetallic alloy. The mechanisms of nitridation, the phases and microstructures are discussed in this work with the support of X-ray diffraction, atomic force, scanning and transmission electron microscopy techniques. The high temperature oxidation behaviour of the nitrided matrices is thereafter evaluated in air and the results are compared to non nitrided specimens. The oxidation kinetics are determined with thermogravimetry and the mechanisms are discussed in light of the oxide phases and microstructures resulting from the previous nitridation treatment. It will be shown that a reduction of the high temperature oxidation resistance occurs for the shortest oxidation times because of trapping of the protective elements. Key words: nitridation, ion implantation, nitrided layer, austenite alloys, ODS Fe-Al alloys, surface oxidation Nitriranje pove~a obrabno in erozijsko odpornost podlag iz kovin in zlitin. Pri nerjavnih jeklih je treba zni`ati temperaturo nitriranja, da bi se izognili izlo~anju CrN, ki bi zmanj{alo splo{no korozijsko odpornost. Nitriranje s fluksom z majhno energijo in veliko gostoto omogo~a, da se ustvarijo relativno debeli sloji v kratkem ~asu in pri nizki temperaturi, odvisno od mikrostrukture in kemijske sestave podlage, kot je prikazano za ~isti Ni, modelno zlitino Ni-Cr20, konvencionalno jeklo AISI 304 L in za intermetalno zlitino FeAl ODS. V tem delu razpravljamo o mehanizmu nitriranja, fazah in mikrostrukturah na temelju rezultatov difrakcije rentgenskega sevanja, opazovanja atomske sile ter vrsti~ne in presevne elektronske mikroskopije. Ocenili smo visokotemperaturno vedenje nitriranih matic na zraku in ga primerjali z nenitriranimi vzorci. Kinetiko oksidacije smo ugotovili s termogravimetrijo in o rezultatih razpravljamo z upo{tevanjem oksidnih faz in mikrostruktur, ki so nastale pri nitriranju. Ugotovili smo, da se zmanj{a visokotemperaturna oksidacijska odpornost pri najkraj{ih ~asih oksidacije zaradi ujetja varovalnih elementov v pasti. Klju~ne besede: nitriranje, ionska implantacija, nitrirana plast, avstenitne zlitine, Fe-Al ODS zlitina, oksidacija povr{ine 1 INTRODUCTION Nitriding of austenitic stainless steels has been exten- sively studied owing to the significant improvements in surface hardness and tribological behaviour 1 as well as in corrosion resistance 2 so long as precipitation of CrN is avoided 3. All these improvements obtained at mode- rate temperature (T < 450 °C) seem to be associated with the formation of an interstitial solid solution of nitrogen in the steel matrix: face centred cubic (fcc) N or "expanded austenite". Various studies suggest that the  would correspond to a fcc phase with a high density of stacking faults likely induced by the internal stresses in the nitrided layer 4-6. However, the effect of the nitriding process to other alloy systems has been poorly investigated to date. For high temperature applications, Ni-base superalloys are typically employed as they show good corrosion and oxidation resistance and excellent resistance to creep and rupture at high temperatures 7. However, they exhibit poor wear resistance. Therefore, plasma nitriding studies have been carried out for instance on Inconel 718 (con- taining the mass fraction of Cr 20 %) at temperatures between 550 °C and 750 °C leading to precipitation of chromium nitride, CrN, and subsequent increase in Knoop hardness 8 and wear resistance until the nitrided layer is worn away 9. Further studies on plasma assisted nitriding of Inconel 690 (containing the mass fraction of Cr 30 %) have been carried out at temperatures between 300 °C and 400 °C 10 where the different depths of nitrogen diffusion have been related to the grain orientations and the anisotropic dependence of stress on strain 11. The low energy-high flux nitrogen implantation approach has rarely been addressed. This is also known Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 157 UDK 621.785.5:669.14.018.8 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(4)157(2008) as an implantation-diffusion technique at relatively low temperatures to promote nitrogen diffusion while arresting CrN precipitation in stainless steels 5,6,12,13. Williamson et al. 14 studied a collection of 16 fcc metals nitrided under the same conditions (0.7 keV, 2 mA cm–2, 400 °C and 15 min). It was shown that the Ni-rich alloys contained much less nitrogen with correspondingly thinner layers than the Fe-rich alloys. Besides, no nitrogen could be detected in the pure Ni specimens but an isolated diffracted peak corresponding to the Ni3N phase. The second study dealt with the tribological properties of Inconel 600 (containing the mass fraction of Cr 16 %) in comparison with the AISI 316 stainless steel, both nitrided at 400 °C for 1 h under 1.2 keV and 1 mAcm–2 15. Again, a thin layer with a maximum concentration of the mole fraction of N 9 % was found in the Ni-rich substrates compared to a 25 at% in the stainless steel, but still offering an increase in hardness and a reduction in wear rate. Despite the extensive use of Ni base superalloys, their significant weight is a limitation in the aeronautic domain as fuel consumption must be reduced. To this end, various intermetallic alloys based on TiAl and on FeAl represent solid alternatives to replace the heavier Ni superalloys 16. In these materials, the nitridation of TiAl have received most of the attention concerning the treatment itself 17–19, their corrosion properties 20 or their high temperature behaviour 21–24. However, little is known on the nitridation of FeAl intermetallic alloys. To the best of our knowledge, only the oxidation kinetics and the likely mechanisms of a nitrided ODS FeAl alloy were reported by Dang et al. 25. Contrary to most of the studies devoted to wear and erosion, the purpose of this work is to review the mechanisms of nitridation by implantation-diffusion (also called low energy-high flux nitridation) in different model (pure Ni, Ni20Cr), commercial (AISI 304L) and candidate materials (ODS FeAl) and the effect on their high temperature oxidation behaviour. The roles of "physics" (crystal structure, grain orientation) and "chemistry" (alloying elements) will be discussed to elucidate the mechanisms involved upon nitridation. On the basis of the resulting phases and microstructures, the high temperature oxidation behaviour will thereafter be interpreted. 2 EXPERIMENTAL Table 1 gathers the base composition and crystal structure of the materials of study. The samples con- sisted of round coupons of varying diameter and 1 mm thick cut from the bars. The main surfaces were mecha- nically polished to a final roughness of 0.01 µm. They were then ultrasonically degreased in acetone and rinsed in 96 % ethanol. Low energy – high flux nitrogen (N2+, N+) implan- tation was carried out at LMP (Poitiers, France) with a Kaufman type ion source at 1.2 keV and a current density of about 1 mA cm–2 for 1 h, corresponding to an estimated dose of about 2.25 · 1019 cm–2. The temperature of the samples was carefully controlled with a thermo- couple attached on the back of the samples. Prior to the nitridation treatment, Ar+ sputtering (1.2 keV, 0.5 mA cm–2 for 15 min) was carried out on each main coupon face to remove the rigid oxide layer that precludes nitridation 26. The backing pressure in the chamber upon the nitriding process was better than 10–2 Pa. Implantation was carried out on both principal coupon faces for the subsequent oxidation experiments, representing about 85 % of the overall surface. Oxida- tion of the nitrided specimens was conducted in a Setaram TG92 thermobalance of 10–6 g of accuracy at 800 °C for 24 h under synthetic air. Heating and cooling rates were fixed at 50 °C/min. Thermodynamic calculations have been performed using the HSC Chemistry software 27 to assess the thermodynamically stable compounds expected to form within the different matrices. The calculations have been carried out at equilibrium conditions at 10–2 Pa (implan- tation conditions) and at atmospheric pressure (after implantation) disregarding collision cascades and sputtering of the surfaces. Only the gas species N2+ (g) or N2 (g) have been considered to react with the substrates, thus taking into account the splitting of the molecules into 2 nitrogen atoms and the corresponding energy release. The characterisation of the implanted and the oxidised specimens was undertaken using contact mode atomic force microscopy (AFM) with an Autoprobe CPR (Veeco Instruments), by X-ray diffraction in a Bruker AXS D-5005 equipment in the –2 configuration and grazing incidence (GIXRD) using Cu K1 ( = 0.15406 nm) radiation as well as by scanning electron micro- scopy (SEM) coupled to energy-dispersive spectrometry (EDS) in a JEOL JSM-4510 LV. Cross sections of the implanted specimens were also prepared for transmission electron microscopy (TEM) studies in a JEOL-JEM 2010 operating at 200 kV. For such purpose, careful mechanical polishing in SiC# 4000 emery paper was performed down to a thickness of about 50 µm. Then, Ar bombardment at 3 keV was carried out in a GATAN PIPS (precision ion polishing system) model 691 at F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 158 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 Table 1: Substrates major composition (/%) and the initial crystal structure Tabela 1: Osnovni sestavni elementi v odstotkih in za~etna kristalna struktura podlag ODS-zlitine, utrjene z disperzijo oksidov substrate Fe Cr Ni Al Y2O3 matrix Ni - - 100 - - fcc Ni20Cr - 20 80 - - fcc AISI 304L 70 20 10 - - fcc ODS* FeAl 60 - - 38 2 orderedB2 * ODS = oxide dispersion strengthened different angles. Vickers microhardness measurements were also performed at increasing loads to get acquainted of the effects of the implantation. 3 RESULTS AND DISCUSSION 3.1 Nitridation by implantation-diffusion After nitridation, all the substrates undergo increased surface microhardness compared to the untreated specimens as depicted in Figure 1. In comparison with the untreated specimens, the hardness increase is of about (8, 20, 250 and 280) % for pure Ni, Ni20Cr, AISI 304L and ODS FeAl, respectively. From these results, it can be considered that nitridation does not effectively occur in pure Ni. This can be due to two interconnected mechanisms. The first one is due to the incorporation of N as an interstitial solid solution and/or to the formation of hard metal nitrides, i. e. "structural deformation", i. e. the appearance of harder crystalline phases. The second one is related to an increased plastic deformation typically occurring upon implantation, i. e. "microstruc- tural deformation", i. e. surface roughness. Regarding the crystallographic phases, the XRD patterns after implantation clearly reveal various features and striking differences among the different substrates as shown in Figure 2. In the case of pure Ni Figure 2(a) the patterns of the untreated and the nitrided specimens are rather similar. Calculations of the lattice parameters of both untreated and nitrided substrates leads to the F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 159 40 60 80 100 120 γ 40 0 γ 22 2 γ 31 1γ 220 γ 200 γ 111 (a) NID untreated In te n s it y , a .u . 2 2 2 2 Θ Θ Θ Θ /degrees /degrees /degrees /degrees 40 60 80 100 120 311 220 γγ N γ N γ γ 222 γ 400 200 γ N γ γ N γ 111 untreated NID (b) In te n s it y , , a . u . (c) γγ γγγγ γγ γ γ γ In te n s it y , a . u . In te n s it y a . u . Figure 2: X-ray diffraction patterns of the different substrates untreated and nitrided by implantation diffusion –NID- (a) pure Ni, (b) Ni20Cr, (c) AISI 304L and (d) ODS FeAl Slika 2: Diagrami difrakcije rentgenskega sevanja za razli~ne podlage, nenitrirane in nitrirane z implantacijsko difuzijo (NID): (a) ~isti Ni, (b) NiCr20, (c) AISI 304 in (d) ODS FeAl H a rd n e s s G P a E E s s t t i i m m a a t t e e d d d d e e p p t t h h , /h , /h µ µ m m (a) H a rd n e s s G P a (b) Figure 1: Evolution of Vickers microhardness with estimated depth of the (a) untreated specimens and (b) nitrided by implantation-diffusion –NID- Slika 1: Evolucija mikrotrdote po Vickersu z ocenjeno globino; (a) nenitriran vzorec (b) nitriran z ionsko implantacijo in difuzijo – NID same results (ao 0.351 nm) hence indicating no expan- sion of the matrix volume. The only remarkable changes involves an attenuation of the <111> directions after nitridation compared to the untreated Ni. Williamson et al. 14 also claimed the absence of N peaks in pure Ni at a lower energy and a higher flux than in our studies. However, they observed a hexagonal Ni3N phase and detected a small shift to lower angles, thus implying retention of a very small amount of nitrogen. Contrary to pure Ni, the nitrided Ni20Cr and AISI 304L substrates Figure 2(b) exhibit a fcc N phase 28 at lower diffraction angles and the original  phase peaks have shifted towards higher diffraction angles 29. For the sake of comparison between both implanted Cr-con- taining substrates a rough estimation of the retained nitrogen has been carried out using the Vegard’s law for substitutional solid solution as follows: aN = a + · CN, where aN and a are the lattice parameters for the N-containing and N-free  phases, respectively, and is the Vegard’s law constant (0.00072 for Fe alloys, also assumed for Ni alloys in this study 14). The concentration of nitrogen is the mole fraction in x(N)%. The results are gathered in Table 2. Table 2: Lattice parameters of the N-containing N and N-free austenite phases, the relative expansion induced, and their corresponding average atomic nitrogen contents, x(N)%, as a function of the diffraction plane (hkl) in Ni20Cr and AISI 304 L Tabela 2: Mre`ni parametri avstenitnih faz -faz z du{ikom in brez njega, relativna inducirana raz{iritev in ustrezna povpre~na atomska vsebnost du{ika x(N)% za razli~ne difrakcijske ravnine (hkl) v Ni20Cr in AISI 304L hkl 111 200 220 311 Ni20Cr aN/nm 0.3580 0.3637 0.3589 0.3612 a/nm 0.3538 0.3540 0.3545 0.3548 expan- sion/% 1.2 2.8 1.2 1.8 x(N)/% 6 13.5 6 9 AISI 304L aN/nm 0.3666 0.3716 0.3666 0.3683 a/nm 0.3572 0.3583 0.3583 0.3583 expan- sion/% 2.6 3.7 2.3 2.7 x(N)/% 13 18.5 11.5 14 Table 2 shows that the retained amount of nitrogen is highly anisotropic. In Ni20Cr the N content is signi- ficantly lower than in the AISI 304L steel regardless of the crystallographic plane. In both substrates however, the highest amount of nitrogen seems to concentrate in the (200) planes and the lowest in the (220). The different partitioning of nitrogen in the various planes also brings about different expansion of the lattice, which in turn may induce strains and stresses. Menthe et al. 30 suggested that a tetragonal distortion of the fcc phase had occurred whereas Fewell et al. 31 proposed a triclinic distortion. Marchev et al. 32,33 considered instead the formation of a martensitic phase. However, any of these would imply the presence of extra peaks never observed on the diffraction patterns. A new structural model nitrogen expanded austenite has been recently proposed by Blawert et al. 4 assuming the effects of deformations and twin faulting commonly observed in fcc metals or alloys. The  expanded austenite would correspond to a fcc phase with a high density of stacking faults likely induced by the internal stresses existing in the nitrided layer 5,6. Indeed, it has been shown that the presence of stacking and twin faults in a perfect fcc lattice produces angular displacements of peaks in XRD patterns 34. The three nitrogen solid solutions observed by Leroy et al. 10 after plasma nitriding of the Ni base alloy Inconel 690 (Ni-30Cr-10Fe, w/%) has not been observed in this work using low energy-high flux implantation. In the ODS FeAl intermetallic, the major contri- bution arises from the (110) and (220) reflections before and after nitridation. At grazing incidence, the hexagonal AlN appears as inferred by three XRD peaks (2 = 33.2°, 36.1° and 38°) and a large and high (110) peak corresponding to the substrate matrix 25. In this alloy, the chemical affinity of N to Al is much greater than that to Fe (e. g., Hf° = –318.0 and –10.5 kJ mol–1 for AlN and Fe4N, respectively) 35 and thus iron nitride formation was not expected to occur. The surface state after nitridation is also quite different among the substrates as shown by plane view SEM in Figure 3. In pure Ni some grains are darker and the orientation of the dislocation slipping bands composing each grain is underpinned; while other grains are lighter in colour and of smoother appearance. In addition, a significant number of protrusions appear throughout the entire surface, especially at grain boundaries. AFM investigations confirm that the roughness can vary between 17.5 nm and 27.5 nm and the aligned bands can be ascribed to the slipping bands due to the presence of stress, as also reported in fcc AISI 316 L stainless steel 36. In Ni20Cr the surface is rather uniform and smooth with no protrusions but with relatively coarse pores. The average roughness is of about 5 – 8 nm but more significant height differences among grains compared to nitrided Ni. The AISI 304L surface is the most heterogeneous of all three fcc nitrided substrates. Some grains are very smooth and deeper and contain large pores thus reminding of the Ni20Cr grains, whereas other grains resemble more the nitrided Ni by underlining the slipping bands, hence being rougher. A common feature observed on the three fcc alloys is the occurrence of twinning within the grains, but again the morphology of twins differs from one matrix to the other. On the contrary, the ordered B2 cubic structure ODS FeAl, the surface seems very uniformly implanted with no twins but some protrusions at the external surface and remaining porosity. This latter feature can be mainly explained by the manufacturing process of this material, which is powder metallurgy. The elongated shape of the protrusions would be related to "softer" areas of the base material, where the strengthening effect F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 160 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 of Y2O3 particles is less important, as revealed by AFM studies Figure 4(a). This microstructure is accom- panied by the highest roughness values, which can attain up to 50 nm. According to the work of Pranevicius et al. 37, the surface roughness can derive from the competition between surface kinetics and bulk diffusion. Nucleation of roughness would first occur by relocation of adatoms, formation of surface vacancies and removal of atoms, which in turn lead to the appearance of clusters of atoms in other regions of the surface. The development of surface roughness subsequently occurs by further reloca- tion and sputtering of atoms displaced by the ion beam. Thereafter, diffusion of nitrogen seems to occur mainly along grain and sub-grain boundaries creating com- pressive stresses 38. Within the metallic substrate, atomic nitrogen can then recombine as molecular nitrogen, raising locally the pressure and inducing plastic defor- mation. Therefore, the amount of deformation would depend on the yield stress of the host material. As a result, a blistered surface appears 39,40. Due to the recession of the metal surface upon implantation, the blisters are peeled off and the pores are then clearly visible in pure Ni and in Ni20Cr Figure 4(b). Since the solubility of nitrogen in nickel is very low the observed porosity is rather shallow. The larger number of pores and blisters are however found at the grain and twin boundaries rather than within the grains as also inferred in a previous study 41. This seems to support the idea that diffusion of nitrogen might be more prone to occur along these short circuit paths, which also become readily F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 161 Figure 4: AFM images of (a) nitrided ODS FeAl showing ridges pinned by Y2O3 particles (b) nitrided Ni20Cr showing the resulting porosity (views of (10 × 10) µm areas) Slika 4: AFM-sliki (a) nitrirani ODS FeAl, ki prikazuje grebene, zasidrane z delci Y2O3 in (b) nitriranega NiCr 20, ki prikazuje nastalo poroznost (ploskvi (10 × 10) µm Figure 3: SEM surface morphology after low energy-high flux nitridation of (a) pure Ni, (b) Ni20Cr, (c) AISI 304L and (d) ODS FeAl Slika 3: SEM-morfologija povr{ine po nitriranju s fluksom z majhno energijo in veliko gostoto pri (a) ~istem niklju, (b) NiCr20, (c) AISI 304 L in (d) ODS FeAl saturated in nitrogen inducing significant plastic deformation. Indeed, EDS microanalyses indicate that no nitrogen has been retained in pure Ni either within the grains or at the grain boundaries where more protrusions are observed. Conversely, in the Cr-bearing alloys the distribution of nitrogen is uneven and confirms the XRD results. For instance, whereas about the mole fraction of N 10 % is present at the surface of Ni20Cr regardless of the location, in AISI 304L stainless steel some of the grains only incorporate about 12 % N and some others contain up to 17 % N, which is close to the chromium content in the substrate. Because of the anisotropic incorporation of N, different compressive stresses are generated. This leads to distortions, plastic deformation and even lattice rotations in an anisotropic fashion 42. As a result of the anisotropic deformation, heterogeneous diffusion will occur modifying the nitrogen ingress rate 36. On the contrary, in the FeAl intermetallic alloy the average composition is Fe-25Al-20N (X/%). This suggests that the N content being introduced could be limited by the Al amount at the surface of the substrate and therefore is only dependent on Al diffusion 43. The SEM cross section morphologies clearly reveal that the only well defined nitrided layers appear on the AISI 304L and the ODS FeAl substrates after a chemical etch (Figure 5). However, the EDS composition profiles (Figure 6) indicate that N has effectively been incor- porated in the Ni20Cr matrix. The maximum N content is found for the ODS FeAl alloy but the depth is the lowest because of N inward diffusion is arrested by the formation of AlN. On the contrary, the shape of the N content is similar in Ni20Cr and AISI 304L. As higher N contents are present in the steel, the nitrided layer is about 1 µm thicker in the steel than in the Ni20Cr alloy. At the substrate/nitrided layer interface, a steep N drop occurs in the steel in comparison with the Ni20Cr alloy. Some explanations can be found from thermodynamic calculations and TEM analyses. Nitrogen has a very low solubility 44 and permeability 45. Upon nitrogen implan- tation chromium shows a strong tendency to form either the fcc CrN (H° = –40 kJ mol–1) or the hcp Cr2N (H° = –38 kJ mol–1) phases, which have not been observed experimentally in Ni20Cr. However, the hexagonal Cr2N phase seems to precipitate at the nitrided layer / AISI 304L interface as shown by cross section TEM and selected area diffraction patterns (SADPs) (Figure 7, Table 3). Fe2N nitride could be also present at the nitrided layer/steel interface but its heat of formation (–18 kJ mol–1) suggests that Cr2N should be the major nitride. This means that the formation of metal nitrides at the nitrided layer/substrate interface would arrest further N inward diffusion and could explain the steep drop of the N content shown in Figure 6. This may indicate that Cr allows to significantly increase the N solubility in Ni. Because nickel rejects nitrogen, the nickel-rich substrate (Ni20Cr) incorporates less nitrogen. On the other hand, from a thermodynamic point of view the free enthalpy (G) is more negative F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 162 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 0 1 2 3 4 5 6 0 5 10 15 20 25 30 D x (N )/ % istance fromsurface, /µmds N in Ni20Cr N in AISI 304L N in ODS FeAl Figure 6: N profile from EDS microanalyses of the cross sections of the nitrided materials. (NB: EDS of ODS FeAl from TEM cross sections) Slika 6: N-profil iz EDS-mikroanalize na prerezu nitriranih mate- rialov (Opomba: EDS ODS FeAl iz TEM prereza) Figure 5: SEM cross section of the nitrided (a) AISI 304L stainless steel and (b) ODS FeAl showing protrusions and the nanograined structure of the substrate Slika 5: SEM-prerez nitriranega (a) nerjavnega jekla AISI 304 in (b) ODS FeAl s protruzijami in nanozrnata struktura podlage (thus, more spontaneous reaction) upon the formation of chromium nitrides than that of iron nitrides (Figure 8). However, the iron effect cannot be neglected if the chemical potential of the species is also taken into account; i. e. when one mole of nitrogen encounters the substrate surface 70 % of the atoms are composed of iron F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 163 Table 3: Data from the selected area diffraction patterns (SADPs) shown in Figures 7 (b) and (c) and the corresponding compounds identified by TEM Tabela 3: Podatkih iz difrakcijskih slik izbranih ploskev (SADPs), ki jih prikazuje slika 7 (b) in (c), in ustrezna spojina, identificirana s TEM experi- mental d-spacing N (experimental) Cr2N (JCPDS 79-2159) Fe2N (JCPDS 73-2102) d-spacing hkl d-spacing hkl d-spacing hkl 2.40c 2.37 (110) 2.39 (110) 2.25b 2.21 (002) 2.21 (002) 2.10c 2.10 (111) 2.09 (111) 2.10 (111) 1.86c 1.86 (200) 1.86 (201) 1.87 (201) 1.52c 1.55 (210) 1.48 (211) 1.46c 1.46 (211) 1.47 (003) 1.35 b 1.37 (300) 1.38 (300) 1.16 b 1.16 (302) 1.17 (302) 0.92c not assigned not assigned not assigned 0.89c 0.89 (400) b data from Figure 7 (b) and c from Figure 7 (c) Figure 7: (a) TEM cross section of the nitrided AISI 304L stainless steel. SADPs of the (b) innermost zone corresponding to a single grain oriented 010Cr2N; and (c) outermost zone representative of various grains Slika 7: (a) TEM-prerez nitriranega nerjavnega jekla AISI 304 L. SDAP (b) notranje cone, ki ustreza enemu zrnu z orientacijo 010Cr2N, in (c) zunanja cona, ki ima razli~na zrna 340 360 380 400 420 440 460 480 500 520 0.00 0.05 0.10 0.15 0.20 Temperature, /°CT , /°CT (a) CrNin Ni20Cr CrNin AISI 304L Cr 2 Nin Ni20Cr Cr 2 Nin AISI 304L 340 360 380 400 420 440 460 480 500 520 0.0 2.0x10 -7 4.0x10 -7 6.0x10 -7 8.0x10 -7 (b) Fe 4 Nin AISI 304L Fe 2 Nin AISI 304L T 2+ x x (n it ri d e )/ (N ) 2+ x x (n it ri d e )/ (N ) emperature Figure 8: Evolution of mole of metal nitride produced per mole of N2+(g) as a function of temperature at 10–2 Pa according to the HSC thermochemical calculations 27 (a) chromium nitrides formation in Ni20Cr and AISI 304L and (b) iron nitrides in AISI 304L Slika 8: Evolucija molarnosti kovinskega nitrida na mol N2 (g) v odvisnosti od temperature pri 10–2 Pa na podlagi termokemi~nih izra~unov 27 (a) nastanka kromovih nitridov v NiCr20 in AISI 304 L in (b) nitridi `eleza v AISI 304 L and only 20 % of chromium. As a result, iron can also enhance incorporation of nitrogen at least to some extent. Indeed, Rivière et al. 5 found that nitrogen was always detected in a nitride type state and that it was preferentially bound to chromium, without specific nitride formation, which agrees well with the trapping-detrapping mechanism proposed by Möller et al. 46. Similarly, a small amount of iron atoms showed the same nitride type bonding but only at the outermost surface. Therefore, iron interaction together with a lower nickel content (which rejects nitrogen) results in higher nitrogen supersaturation in the superficial layers of AISI 304L than in Ni20Cr. Thereafter, because of the difference in chemical potentials between the external layer and the bulk, diffusion will be enhanced. As a result, the Fe-based alloy, which incorporates more nitrogen, will exhibit a higher degree of deformation. This induces significant swelling of the grains, thus developing rougher surfaces than Ni20Cr. For the ODS FeAl intermetallic alloy, the nitrided layer has a nanostructured morphology and at the nitrided layer / substrate interface an iron band segre- gates (Figure 9). Diffraction patterns of the different areas point out the different features observed in these samples such as the nanometre scale of the nitrided layer characterised by the typical rings corresponding to FeAl as well as some spots at shorter distances belonging to AlN. As summarised in Table 4, some of the distances may also correspond to -Fe. Sanghera and Sullivan 35 found that nitrogen implanted at low energy and low flux into pure aluminium did not render stoichiometric AlN because the radiation damage induced many vacancies, interstitials and defects. From our EDS analyses, only the outermost layers would contain enough nitrogen to produce the hexagonal AlN phases massively and therefore, once the average values of nitrogen decrease, a mixture of FeAl containing dispersed particles of AlN occurs closer to the nitrided layer/substrate interface. From the TEM results a combined mechanism of nitrogen diffusing inwardly and aluminium outwardly during the nitridation treatment would occur. This countercurrent diffusion would be promoted by the creation of short-circuit diffusion paths, i.e. the grain boundaries of the nanostructured layer. Indeed, diffusion of indium (isoelectronic with aluminium) has been found to be faster than that of iron by a factor of about two in Fe66Al34 and Fe50Al50 47, which helps in corroborating the suggested mechanism. Table 4: Experimental d-spacings obtained with 0.15 µm-diaphragm SADPs at the nitrided layer/substrate interface in the as-nitrided intermetallic alloy and their correspondence to the planes of the identified compounds Tabela 4: Eksperimentalne d-razdalje, izmerjene pri SDPS z 0,15 µm veliko zaslonko, na medpovr{ini nitridna plast/podlaga v nitrirani intermetalni spojini in njihova lega glede na ploskev indentificiranih spojin Experimental d-spacing, nm FeAl JCPDS 33-20 -Fe JCPDS 89-4186 AlN JCPDS 25-1133 0.252 – – 002 0.207 110 110 – 0.160 111* – 110 0.143 200 200 – 0.119 211 211 202 * superstructure peak 3.2 High Temperature oxidation behaviour Because of their specific uses, the oxidation tests were conducted at different temperatures and the results will be therefore presented independently. 3.2.1 Oxidation of Ni and Ni20Cr: 700 °C and 800 °C Figure 10 shows the mass gain curves against time for both untreated and nitrided specimens. It can be observed that in nitrided Ni no significant difference is observed at both temperatures. On the contrary, in Ni20Cr nitridation increases significantly the overall mass gain. Assuming parabolic behaviour, the oxidation constants have been calculated by the (M/S)2 vs. time F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 164 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 Figure 9: TEM cross section showing (a) the nanostructured morphology of the nitrided layer and (b) the nitrided layer/substrate interface. The band of -Fe segregated at this interface is indicated between arrows Slika 9: TEM-prerez, ki prikazuje (a) nanostrukturirano morfologijo nitridne plasti, in (b) medpovr{ina nitrirana plast/podlaga. Plast segregiranega -Fe na tej medpovr{ini je prikazana med pu{~icami method. In Ni, the parabolic rate constant (kp) values of 4 · 10–12 g2 cm–4 s–1 and 2.5 · 10–11 g2 cm–4 s–1 are found for 700 °C and 800 °C, respectively. However, in Ni20Cr the kp values increase about one order of magnitude from (1.0 · 10–15 to 8.3 · 10–15) g2 cm–4 s–1 at 700 °C and from (2.3 · 10–14 to 2.3 · 10–13) g2 cm–4 s–1 at 800 °C after the whole oxidation test. The XRD patterns have revealed the formation of NiO oxides in both untreated and nitrided Ni samples, together with some weak peaks of the substrate, indicating a relatively thick oxide layer at both temperatures. The oxide species developed on Ni20Cr are the same for both the untreated and nitrided specimens at either temperature and these include NiO, NiCr2O4 and Cr2O3. At the highest temperatures, more contribution of Cr2O3 oxide is found to occur. However, the substrate/oxide intensity ratios are always higher at any temperature than in the nickel substrates. This means that a thinner oxide layer is obtained in the Ni20Cr samples after 24 h of isothermal oxidation. Regarding the expanded austenite (N) phase (Figure 11) oxidation at 700 °C for 24 h brings about shifting of the N and  peaks towards the original  phase (2 = 44.28°) giving rise to the observed doublet. This clearly implies redi- stribution of nitrogen in the matrix but no nitride phase can be derived from the XRD results. The SEM morphologies are also completely diffe- rent. Whereas the untreated specimens develop even and homogeneous oxide scales, the oxide layers spall off or oxide plates develop in nitrided Ni [Figure 12(a) and (b)]. In Ni20Cr oxidation occurs preferentially depen- ding on the grain orientation and grain boundary. At the lowest temperatures, the Ni20Cr samples are distinc- tively covered of oxides [Figure 12 (c) and (d)], which are more developed at 800 °C [Figure 12(e) and (f)], F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 165 40 45 50 55 γγ N γ γ γ N γ N untreated nitrided nitrided + 700°C nitrided + 800°C In te n s it y , a .u . 2Θ/degrees Figure 11: Selected range of the obtained on the untreated, as-nitrided and nitrided and oxidised at 700 °C and 800 °C Ni20Cr substrates. N.B: Only the matrix peaks are indicated. See text for further information concerning oxide species. Slika 11: Izbrana podro~ja na nenitriranem, nitriranem in nitriranem ter oksidiranem Ni20Cr podlagah. (Opomba: Prikazani so le vrhovi matice. V tekstu je pojasnilo o vrstah oksidov). 0 5 10 15 20 25 0.0 0.5 1.0 1.5 O O ∆M S/ /( m g c m ) – 2 ∆M S/ /( m g c m ) – 2 x x i i d d a a t t i i o o n n t t i i m m e e , /htoxid , /htoxid 700°C 800°C (a)untreated 700°C 800°C nitrided 700°C 800°C 0 5 10 15 20 25 0.00 0.02 0.04 0.06 0.08 (b) 700°C 700°C 800°C 800°Cuntreated 700°C 800°C nitrided 700°C 800°C Figure 10: Isothermal oxidation at 700 °C and 800 °C for 24 h in synthetic air (a) untreated and nitrided Ni and (b) untreated and nitrided Ni20Cr Slika 10: Izotermna 24-urna oksidacija pri 700 °C in 800 °C v sinteti~nem zraku; (a) nenitriran in nitriran Ni in (b) nenitriran in nitriran NiCr20 Figure 12: SEM surface morphologies developed at high temperature on (a) and (b) nitrided Ni at 700 °C and 800 °C; (c) and (d) nitrided Ni20Cr at 700 °C and (e) and (f) Ni20Cr at 800 °C Slika 12: SEM-morfologija povr{in, ki so nastale pri visoki tempe- raturi na (a) in (b) nitriranem Ni pri 700 °C in 800 °C, (c) in (d) nitriranem Ni20Cr pri 700 °C in (e) ter (f) pri 800 °C thus suggesting that the N implantation effect is lost at the highest temperature, as confirmed on the cross sections by SEM and EDS microanalyses. Indeed, the N content drops from about the mole fraction 10 % at the surface of the as-nitrided specimens to 3.5 % and 0 % after 24 h of oxidation at 700 °C and 800 °C. At 800 °C, some tiny metal nitrides precipitate (about 3 % N). 3.2.2 Oxidation of AISI 304L: (400, 450, 500 and 550) °C Figure 13 shows the mass gain curves as a function of time for both the untreated Figure 13 (a) and the nitrided Figure 13 (b) specimens. Oxidation is more significant in the nitrided samples than in the untreated steel upon the first oxidation times at any temperature as a result of both a chemical and physical effect 48. The first one is related to the amount of implanted nitrogen, whereas the second refers to the defects induced upon implantation. The XRD patterns of the untreated steel show mainly the substrate peaks, i. e. austenite and ferrite  phases are observed, indicating the low thickness of the scale. The small participation of the ferrite  phase has been previously reported to occur as a result of both plastic deformation induced upon grinding 49 and after high temperature exposure due to chromium outward diffusion, which partially destabilise the austenitic phase until oxide formation is accomplished 50. Only in GIXRD at 15° a weak hematite (-Fe2O3) signal appears at 550 °C. In the nitrided specimens, the N phase is present up to 500 °C Figure 14 but it evolves towards a more stable state, which implies rejection of nitrogen in solid solution in the nitrided layer. Mändl et al. 51 after annealing of a nitrided austenitic stainless steel at 425 °C found that the lattice expansion was considerably reduced, yielding a new N2 phase and additional CrN peaks under 8° of incidence. The XRD results of Figure 14 indicate that the decomposition of the N phase occurs by formation of CrN and two other FeNi phases  (bcc) and (fcc) probably containing a small amount of Cr. The precipitation of the cubic CrN phase is detected from 500° C since at 400°C the mobility of chromium in the AISI 304L stainless steel is low 52. The presence of the  phase can be explained as for the untreated steel (see above) as well as from the nitrogen partial dissolution from the (Fe,Cr)2N leading to a ’-(N) martensite 53. This fact, together with the substantial decrease of superficial nitrogen observed by EDS, indicates that upon oxidation, nitrogen may mainly diffuse inwardly towards the bulk. Such diffusion coupled to the outward diffusion of chromium from the bulk alloy gives rise to the more thermodynamically and kinetically stable CrN nitride. Öztürk and Williamson 54 also found the formation of CrN upon the post-annealing of the AISI 304 stainless steel at 400 °C. However, decomposition of such phase was not observed even after 36 h but a dramatic reduction in N content due to inward and outward diffusion. The oxide scales developed in the untreated steel evolve mainly through bulk alloy outward diffusion and not via the grain boundaries Figure 15 (a) and (b). On the contrary, oxide development is more pronounced on the surface of the nitrided steel even at the lowest oxidation temperatures as a result of the deformation induced through ion implantation Figure 15 (c). Again, as the oxidation temperature increases, the oxide cove- F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 166 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 30 40 50 60 70 80 90 α γ γ N α γ N ,γ γ N ,γ γ γ γ γ N γ α,a,b a ,ba a a b a a = Fe x Cr 2-x O 3 b = CrN 400°C 450°C 500°C 550°C In te n s it y , a .u . 2Θ/degrees Figure 14: GIXRD patterns at 15° of the nitrided AISI 304L stainless steel after oxidation in air for 24 at 400, 450, 500 and 550 °C Slika 14: GIXRD-odsevi pri 15° za nitrirano nerjavno jeklo AISI 304L po 24-urni oksidaciji na zraku pri (400, 450, 500 in 550) °C O O ∆M S/ /( m g c m ) – 2 ∆M S/ /( m g c m ) – 2 x x i i d d a a t t i i o o n n t t i i m m e e , /htoxid , /htoxid Figure 13: Isothermal oxidation of the AISI 304L stainless steel at 400, 450, 500 and 550 °C for 24 h in synthetic air (a) untreated and (b) nitrided Slika 13: Izotermna oksidacija nerjavnega jekla AISI 304 L 24 h pri (400, 450, 500 in 550) °C v sinteti~nem zraku; (a) nenitrirano, (b) nitrirano rage increases depending on the roughness of each grain [Figures 15 (d), (e) and (f)]. Contrary to the untreated steel, diffusion seems to occur through both the bulk alloy and the grain boundaries. The EDS microanalyses show the only presence of oxygen, chromium and iron on the scales (Figure 16). It can be observed that oxide formation is clearly promoted with increasing temperature whereas the superficial nitrogen content decreases. The ratios Fe/Cr after oxidation of the untreated steel at any temperature are relatively the same in comparison with the unoxidised steel. According to the cross section analyses, the oxidising temperature seems to provide enough energy to induce chromium and nitrogen diffusion so that tiny precipitation of CrN might occur as for the Ni20Cr substrates. This in turn leads to the N disappearance and the formation of the  phase. Öztürk and Williamson 54 observed the vanishing of the magnetic state of the N phase as the post-annealing treatment of the fcc AISI 304 steel at 400 °C progressed with time, in agreement with the above results. 3.2.3 Oxidation of ODS FeAl: 800 °C After the 24 h exposure at 800° C, the weight gains of the nitrided specimens was four-fold that of the un-nitrided, with kp values of about 4.7 · 10–8 mg2 cm–4 s–1 for the latest stage 25, i.e. even 10 times faster than those of nitrided -iron 55. F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 167 400 450 500 550 10 20 30 40 50 60 70 Ountreated Onitrided Cr untreated Cr nitrided Fe untreated Fe nitrided C o m p o s it io n , O x /% x (N )/% xidation temperature, /°CToxid. 4 6 8 10 Nnitrided N itro g e n c o n te n t, Figure 16: EDS surface composition of the oxidised surfaces of both untreated (blue) and nitrided (red) as a function of the oxidation temperature Slika 16: EDS-sestava oksidirane povr{ine nenitrirane (modro) in nitrirane povr{ine v odvisnosti od temperature oksidacije Figure 15: SEM surface morphologies developed the AISI 304L stainless steel (a) and (b) untreated and oxidised at 500 °C and 550 °C, respectively; and of nitrided and oxidised at (c) 400 °C, (d) 450 °C, (e) 500 °C and (f) 550 °C Slika 15: Morfologija povr{ine, nastale na nerjavnem jeklu AISI 304L: (a) izhodna in (b) oksidirana pri 500 °C in 550 °C; nitrirana in oksidirana pri (c) 400 °C, (d) 450 °C, (e) 500 °C in (f) 550 °C 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 0 10 20 30 40 50 60 70 (c) O Nx5 Al Fe C o m p o s it io n , D x /% istance fromsurface, /µmds Figure 17: (a) Bright field TEM image of the stratified oxide scale, (b) SADP at the inner scale/substrate interface and (c) EDS micro- analyses across all the layers Slika 17: (a) TEM-slika v svetlem polju za plastasti oksidni sloj, (b) SADP na medploskvi notranja plast {kaje/podlaga in (c) EDS-analiza po prerezu vseh plasti After cooling, the oxide scales are shown to extensively spall off, depicting at least two subscales, a very convoluted top layer and an inner layer with white needles. Although the surface EDS microanalysis and XRD only indicated the presence of iron oxide (hematite), TEM inspection reveals a more complex oxide scale [Figure 17 (a)]. As shown in Figure 17 (b), the selected area diffraction patterns (SADPs) at the top and bottom inner scale suggests the existence of the FeAl2O4 phase, with a contribution of hexagonal AlN at the oxide/substrate interface, whose reflections are summarised in Table 5. The EDS (3 nm spot) microanalyses across the entire scale Figure 17 (c) confirm that the outermost oxide layer is only composed of iron and oxygen and is about 0.25 µm thick. According to the XRD patterns 25 this phase has been identified as -Fe2O3. Underneath, a 0.2 µm thick layer exists, which is mostly enriched in aluminium, which may correspond to -Al2O3. The inner oxide layer is the largest (about 0.5 µm thick) with more of Al than Fe, hence suggesting the presence of the FeAl2O4 phase found by SADP. At the spinel substrate interface nitrogen is found to concentrate, accompanied with a drop in the oxygen content. Table 5: Experimental d-spacings obtained from SADPs at the inner oxide layer / substrate interface after oxidation at 800 °C of the nitrided ODS FeAl and their correspondence to the planes of the identified compounds. Tabela 5: Esperimentalne d-razdalje iz SADPs na medpovr{ini notranja oksidna plast/podlaga po oksidaciji nitrirane ODS FeAl pri 800 °C in njihova lega glede na ravnine identificirane spojine d-spacing, nm FeAl2O4 (JCPDS 34-192) AlN (JCPDS 25-1133) 0.465 111 – 0.270 – 100 0.246 311 002 0.204 400 – 0.186 331 102 0.139 – 103 Such complex structure allows to shed some light on the oxidation mechanisms after nitridation of ODS FeAl. Although outward diffusion of indium (isoelectronic with aluminium) is two times faster than that of iron in Fe66Al34 and Fe50Al50 47, there is not enough aluminium available at the top surface to form the oxide since this is trapped as AlN throughout the compact layer. At the diffusion layer, -Fe was found to segregate at the fragmented FeAl matrix together with some AlN. Such iron is readily available for outward diffusion through the important number of short circuit paths that represent the grain boundaries and defaults created upon nitrida- tion. However, once the outer iron scale is developed, the oxygen partial pressure decreases and only alumi- nium oxide is able to form owing to its higher thermodynamic stability. At reduced pressure only the -Al2O3 phase should develop but its reaction with either Fe2O3 56 or FeO 57, a FeAl2O4 spinel oxide forms. At reduced oxygen partial pressures dissolution of AlN also takes place 58 and indeed, no nitrogen is found in any of the oxide layers except at the spinel / substrate interface. This implies that after dissolution of the nitride, nitrogen seems to diffuse further inwardly towards the substrate whereas the resulting aluminium tends to be transported outwardly, stabilising the spinel oxide phase. In previous works 59 it was already claimed that the spinel layer would only be partially effective in hindering outward aluminium diffusion when the grains coarsened with increasing temperature. 4 SUMMARY AND CONCLUSIONS Similar low energy, high flux nitridation processing conditions on different fcc metallic substrates lead to very different results depending on the chemical composition of the matrix. It has been shown that pure nickel does not develop an expanded austenite phase due to rejection of nitrogen. The tiny retained amount of nitrogen creates blisters and pores as nitrogen tries to be triggered off the substrate. The major surface roughness is then developed by sputtering. On the contrary, with the addition of chromium an expanded austenite phase develops but nitrogen uptake is still limited by nickel rejection. In turn, iron atoms can thermodynamically favour nitrogen uptake at least at the outermost surface. The higher the nitrogen intake, the higher the degree of deformation including grain swelling, which leads to rougher and harder surfaces. On the contrary, in the presence of Al (ODS FeAl alloy) brings about the formation of an outer AlN compact layer and an inner diffusion layer in which AlN, -Fe segregation and fragmentation of the FeAl grains occur. Deformation of the material also seems to be induced upon implantation. The high temperature oxidation behaviour seems to depend thereafter of the microstructure and chemistry of the implanted specimens. Whereas in pure Ni nitridation does not basically change the oxidation kinetics, on Ni20Cr the kinetics are increased by one order of magnitude. This is mainly due to trapping of chromium by the implanted nitrogen, hence impeding the formation of the protective Cr2O3 scale. For the longest exposures enough chromium flux from the matrix seems to be ensured. Deformation induces oxide scale spallation as shown in nitrided Ni. In the AISI 304L stainless steel oxidation of the nitrided specimens brings about progressive disappearance of the N phase accompanied with the appearance of an  phase and precipitation of fcc CrN nitride. This phase transformation phenomenon, in turn, may supply chromium to the oxide scale, since the nitrided samples have shown to be enriched in this metal in comparison with the untreated steel. Our results suggest that oxidation seems to proceed by oxygen inward diffusion through the more nitrogen rich planes composing the grains. In the nitrided ODS FeAl aluminium is trapped as AlN, therefore allowing the formation of a non protective outer Fe2O3 scale. Once the oxygen partial pressure is reduced dissolution of AlN occurs. Thereafter, nitrogen is further transported F. PEDRAZA: LOW ENERGY-HIGH FLUX NITRIDATION OF METAL ALLOYS ... 168 Materiali in tehnologije / Materials and technology 42 (2008) 4, 157–169 inwardly, whereas aluminium diffuses outwardly. As a result, the FeAl2O4 spinel inner layer is developed under the Fe2O3 top layer. Therefore, the nitridation treatment changes the oxidation mechanisms. Overall, the role of implanted nitrogen is to retard the establishment of an external alumina scale, but does not seem to impede it. Longer oxidation tests should be carried out to confirm this possibility. Acknowledgments J.P. Rivière and G. Abrasonis from the LMP labo- ratory at Poitiers (France) are gratefully acknowledged for the nitridation experiments. 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NUMERICAL AND EXPERIMENTAL ANALYSES OF THE DELAMINATION OF CROSS-PLY LAMINATES NUMERI^NA IN EKSPERIMENTALNA ANALIZA DELAMINACIJE V KRI@NIH PLO[^ATIH LAMINATIH Robert Zem~ík, Vladislav La{ University of West Bohemia in Pilsen, Department of Mechanics, Univerzitní 22, 306 14, Plze, Czech Republic zemcikkme.zcu.cz Prejem rokopisa – received: 2007-09-20; sprejem za objavo – accepted for publication: 2008-01-16 This article focuses on a numerical and experimental investigation of the delamination of cross-ply FRP laminates. The tested samples were cut from plates 0/90/90/0 made of unidirectional carbon-epoxy prepregs in a vacuum autoclave. The experimental tests were performed on precracked double-cantilever beam samples according to ASTM standards with Mode-I loading. The load-displacement relations were recorded during the test and the crack length was measured optically with a digital camera. The corresponding numerical simulations were performed in the finite-element code MSC.Marc. The goal was to assess the critical value of the energy-release rate G, which was chosen as the interlaminar fracture toughness. The simulation used the equality between the energy-release rate and the J-integral for the elastic case. Keywords: composite, cross-ply, beam, delamination, experiment, simulation V delu obravnavamo numeri~no in eksperimentalno raziskavo delaminacije v kri`nih plo{~atih FRP-laminatih. Preizku{anci so bili odrezani iz plo{~ 0/90/90/0, ki so bile izdelane v avtoklavu iz surovcev ogljikovo vlakno-epoksi. Preizkusi so bili izvr{eni na vnaprej razpokanih dvojno vpetih nosilcih skladno z ASTM-standardi z obremnitvijo Mode 1. Odnos obremenitev – pomik je bil ugotovljen med preskusom, dol`ina razpoke pa se je merila opti~no z digitalno kamero. Za numeri~no simulacijo je bila uporabljena metoda kon~nih elementov s kodo MSC.Marc. Cilj je bil dolo~iti kriti~no vrednost hitrosti sprostitve energije G, ki je bila opredeljena kot interlaminarna `ilavost loma. V simulaciji je uporabljena enakost med hitrostjo sprostitve energije in J-integralom za primer elasti~nosti. Klju~ne besede: kompoziti, kri`na plo{~a, nosilec, delaminacija, eksperiment, simulacija 1 INTRODUCTION One of the key damage mechanisms in laminates is the origin and propagation of a failure or a crack between individual layers; this is known as delamination and it must be considered in the design of a laminated structure. Delamination can be caused by imperfections during the manufacturing process or due to static and dynamic loads. The existence of delamination in a composite material degrades its stiffness and in certain cases it can degrade the stability to a critical level. The dangerous factor is the propagation of the delamination, which is influenced by the geometrical parameters, the material characteristics and the loading type. Both types of propagation can be present: slow and stable as well as fast and unstable. In this investigation we have looked at a numerical simulation of the delamination of a laminated specimen made of unidirectional fiber-reinforced composite layers. This follows on from the work started in 1. The aim is to design a numerical model for the simulation of crack propagation. An important parameter used in such analyses is the critical value of the energy-release rate (Gc). The majority of studies investigating the delami- nation on laminated structures use either the critical value of energy-release rate Gc 2–4 or the stress-intensity factor Kc 5 as the interlaminar fracture toughness. The early studies were focused mainly on the experimental procedures and analytical solutions. These were later followed by studies dealing with numerical simulations based on a finite-element analysis and introducing new, special element types for the modeling of the delami- nation 6–8. It is known that the critical value of Gc as calculated according to the ASTM standard 9 does not behave as a constant. The critical value changes during the pre- scribed test by as much as tens of percent. Therefore, it is questionable as to which value should be chosen, for instance, for the numerical simulation. In this work the value of Gc is assessed by experi- mental measurements on a Mode-I delamination specimen. The analysis was performed on a cross-ply laminate manufactured from unidirectional plies reinforced with carbon fibers in an epoxy resin. It is known 10 that in the case of linear stress-strain behavior there is equality between G and the value of the J-integral (J). Therefore, it was possible to substitute the evaluation of the decisive parameter Gc with the calculation of Jc in the numerical simulations. The numerical calculations are carried out using the finite- element method in MSC.Marc. The difference between the experimentally and numerically obtained dependencies of force vs. displace- Materiali in tehnologije / Materials and technology 42 (2008) 4, 171–174 171 UDK 678.6/.7:620.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(3)171(2008) ment for the critical value Gc is presented. The basic material parameters (elasticity constants) used in the simulations were identified in previous studies 11,12, where the damage and failure of specimens made of the same material are investigated. 2 EXPERIMENT Firstly, an experimental assessment of the inter- laminar toughness for Mode-I delamination was carried out according to the standard ASTM D 5528-01 9. The specimens used were rectangular strips denoted as DCB (Double Cantilever Beam) having dimensions l × b = (160 × 19.7) mm (see Figure 1). Each specimen was cut out using a water jet from a 4.6-mm-thick laminated cross-ply composite plate with a lay-up 0/90/90/0, whereas the middle part was 0.3 mm thick (see Figure 2). The plate was manufactured from epoxy prepregs reinforced with continuous Toray T600SC carbon fibers using autoclave technology. A non-adhesive aluminum foil, which served as the initial crack, was inserted into the midplane of the plate (i.e., between the 90 degree plies) during manufacture (see Figures 1 and 4). The thickness of the foil was 11 µm. The experiment was performed on the Zwick/Roell Z050 testing machine. The opening force was applied using two piano hinges bonded on the lower and upper surfaces of the specimen (see Figures 1 and 3). During the testing process the dependence of the opening force F vs. the transverse (or load point) displacement  was recorded (see Figure 3). The crack propagation with time (crack length a) was inspected optically using a digital camera (Canon EOS 400D, Sigma 105/2.8 EX DG MACRO) every 10 s (resolution approximately 20 px/mm). The optical measurement also served as a verification of the transverse displacement values, which might differ, in general, from the grip movement, as recorded by the testing machine, due to its imperfect rigidity. The delamination process had a slow and stable character up to the final rupture. The speed of the grip movement was 10 mm/min. The major problems identified during the experi- ments were the inhomogeneity of the initial crack zone (the end of the foil) and the so-called fiber bridging between the upper and lower parts of the specimen (see Figure 5). The bridging of fibers occurs in composite specimens within the unidirectional layers, as in this case. 3 ESTIMATION OF THE ENERGY-RELEASE RATE The energy-release rate G is defined mathematically, in general, as G b U a = − 1 d d (1) where dU is the differential increase in the strain energy, da is the differential increase in the delami- nation (or crack) length, and b is the specimen width. Concerning the investigated Mode I for the DCB specimen, the corresponding energy-release rate GI can be expressed using the Euler-Bernoulli theory of beams as 3,13 R. ZEM^ÍK, V. LA[: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE DELAMINATION ... 172 Materiali in tehnologije / Materials and technology 42 (2008) 4, 171–174 Figure 1: DCB specimen with attached piano hinges and inserted foil Slika 1: DCB-vzorec s pritrjenim {arnirom in z vstavljeno folijo Figure 4: Crack surfaces of the fractured DCB specimen Slika 4: Povr{ina razpok pretrganega DCB-vzorca Figure 3: DCB specimen experimental setup Slika 3: DCB-vzorec z eksperimentalno postavitvijo Figure 2: Material lay-up (showing the part without the foil) Slika 2: Zlog materiala (prikazan je del brez folije) G F a bEII = 2 2 (2) where EI is the bending stiffness, i.e., the sum of the Young’s moduli in the axial direction multiplied by the inertia moments for all the plies (see Figure 1). There are several possible ways of calculating the strain-energy release (or eventually the interlaminar fracture toughness) according to the ASTM standard 9. One way is the so-called Modified Beam Theory (MBT) method. The energy-release rate in this case is calculated as G F baI = 3 2 δ (3) Since the beam is not perfectly built-in (a certain amount of rotation can occur at the crack front, i.e., where a clamped condition is assumed), a correction can be applied, which assumes that the delamination is larger by the amount . Hence, the crack length is (a+) and the corrected energy-release rate is then given by G F b aI = + 3 2 δ ( )∆ (4) where the value of ∆ is calculated according to the procedure given in the standard 9 and the values obtained for both types of specimens are shown in Table 1. Equation (2) can be corrected in a similar way. 4 NUMERICAL SIMULATION The numerical simulation was performed using the finite-element method in the MSC.Marc system. In this case the analysis was solved as a plane-strain problem. The geometry was modeled using 4-node rectangular elements. The 0/90/90/0 lay-up was meshed with 8/2/2/8 layers of elements in the thickness direction and 600 elements along the length. A different model was prepared for each crack length a. The equality between the values of the energy-release rate and the J-integral in the case of the elastic analysis was used in the simulations. MSC.Marc calculates the J-integral using the DeLorenzi method 14. The material is assumed to be homogeneous, linearly elastic, and orthotropic, and to have the following elasticity constants: longitudinal Young’s modulus EL = 110000 MPa, transverse modulus ET = 7700 MPa, Poisson’s ratio νLT = 0.28 and shear modulus GLT = 4500 MPa, which were identified previously 11,12. The critical values of the energy-release rate and the parameter ∆ were calculated using (4) in order to best characterise the averaged experimental dependency in Figure 6. The hypothetical delamination curves for the interval 0.1 < GIc < 1.5 kJ/m2 are shown in the graph together with the optimum curve corresponding to the values displayed in Table 1. The reconstructed load-displacement curve from the FEA is compared with the experiment. Table 1: Measured and calculated parameters of the DCB specimen Initial crack length /mm a0 62 Crack-length correction /mm ∆ 5.12 Critical energy-release rate /(kJ/m2) GIc 1.373 R. ZEM^ÍK, V. LA[: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE DELAMINATION ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 171–174 173 Figure 6: Load-displacement curves for the DCB specimen Slika 6: Krivulje obremenitev – pomik za DCB-vzorec Figure 5: Sequence of photographs showing the fiber-bridging progress Slika 5: Sekvenca posnetkov, ki prikazuje napredovanje vlaknate premostitve 5 CONCLUSIONS An experimental investigation of low-speed Mode-I delamination was carried out on carbon-fiber-reinforced epoxy specimens in order to estimate the critical energy-release rate value GIc. The double-cantilever beam specimens with a cross-ply lay-up were manufactured with inserted aluminum foil serving as an initial crack. The critical energy-release-rate value was determined using the Euler-Bernoulli theory and the ASTM standard from the averaged experimental data. A numerical model was created in the FEA code MSC.Marc for the simulation of the experiment. The model took advantage of the equality between the energy-release rate and the J-integral in the case of an elastic material. The load-displacement dependency was reconstructed and compared with the experimental results. It is clear that it is possible to consider the critical energy-release rate value as the interlaminar fracture toughness in the case of the delamination of a composite with transversely oriented fibers. Acknowledgement This work was supported by the research project GA AV IAA200760611 and research project MSM 4977751303. 6 REFERENCES 1 V. La{, R. Zem~ík, P. M{ánek: Numerical simulation of composite delamination with the support of experiment. Acta Mechanica Slovaca, 1 (2006), 303–308 2 M. Hojo et al.: Mode I delamination fatigue properties of inter- layer-toughened CF/epoxy laminates. Composite Science and Technology, 66 (2006), 665–675 3 T. Kusaka et al.: Rate dependence of Mode I fracture behaviour in carbon-fibre/epoxy composite laminates. Composites Science and Technology, 58 (1998), 591–602 4 A. B. Pereira, A. B. Morais: Mode I interlaminar fracture of carbon/epoxy multidirectional laminates. Composites Science and Technology, 64 (2004) 2261–2270 5 W. T. Chow, S. N. Atluri: Stress intensity factors at the fracture parameters of delaminations crack growth in composite laminates. Composites Part B, 28B (1997), 375–384 6 R. Borst, J. J. C. Remmers: Computational modelling of delami- nation. Composites Science and Technolgy, 66 (2006), 713–722 7 C. H. Roche, M. L. Accorsi: A new finite element for global modelling of delaminations in laminated beams. Finite Elements in Analysis and Design, 31 (1998), 165–177 8 F. Shen, K. H. Lee, T. E. Tay: Modeling delamination growth in laminated composites. Composites Science and Technology, 61 (2001) 1239–1251 9 ASTM D5528-01, Standard Test Method for Mode I Interlaminar Fracture Toughness of Unidirectional Fiber-Reinforced Polymer Matrix Composites, Annual Book of ASTM Standard, 2002, 249–260 10 D. P. Miannay: Fracture mechanics, Springer, 1998 11 R. Zem~ík, V. La{: Identification of composite material properties using progressive failure analysis. Proc. of Computational mechanics 2005, Ne~tiny, 2005, 695–700 12 R. Zem~ík, V. La{: Numerical simulation of damage in fiber-rein- forced composites and comparison with experiment. Proc. of 22nd Danubia-Adria symposium on experimental methods in solid mechanics, Monticelli Terme, Parma, 2005, 174–175 13 M.-S. Sohn, X-Z Hu: Comparative study of dynamic and static delamination behaviour of carbon fibre/epoxy composite laminates. Composites, 26 (1995), 849–858 14 MSC.Marc Volume A: Theory and user information, Version 2005. MSC.Software Corporation, 2004 R. ZEM^ÍK, V. LA[: NUMERICAL AND EXPERIMENTAL ANALYSES OF THE DELAMINATION ... 174 Materiali in tehnologije / Materials and technology 42 (2008) 4, 171–174 K. STRANSKY ET AL.: APPLICATION OF THE THEORY OF PHYSICAL SIMILARITY ... APPLICATION OF THE THEORY OF PHYSICAL SIMILARITY FOR THE FILTRATION OF METALLIC MELTS UPORABA TEORIJE FIZIKALNE PODOBNOSTI ZA OPIS FILTRIRANJA KOVINSKE TALINE 1Karel Stránský, 2Jií Ba`an, 2Jana Dobrovská, 3Martin Balcar, 3Pavel Fila, 3Ludvík Martínek 1VUT Brno, FSI, Czech Republic 2V[B-Technical University of Ostrava, Czech Republic 3@AS, a.s. @ár nad Sázavou, Czech Republic stranskyfme.vutbr.cz Prejem rokopisa – received: 2007-10-08; sprejem za objavo – accepted for publication: 2008-05-07 The Bernoulli equation is the basis for the primary description of the flow of a real metallic melt through the pouring system for the filling of the casting mould with an inserted ceramic filter. In principle, a modified, dimensionless form of the Bernoulli equation can be used for the determination of the loss coefficient as a general function of the dimensionless criteria – the Reynolds, Froude and Euler numbers. It was verified by modelling the flow of the modelling liquid (in this case water) through ceramic filters. In the same interval of Reynolds numbers the loss coefficient was greater for foam filters than for filters with direct holes (strainers); however, the outlet coefficient µ of the foam filters was, in identical conditions, significantly lower than that of filters with direct holes. Key-words: similarity criteria, modelling, steel flow, ceramic filters, steel filtration Bernoullijeva ena~ba je podlaga za opis pretoka taline skozi livni sistem z vstavljenim kerami~nim filtrom. Na~eloma je mogo~e uporabiti modificirano brezdimenzijsko obliko Bernoullijeve ena~be za dolo~itev koeficienta izgube kot funkcije brezdimenzijskega kriterija – Reynoldsovega, Froudovega ali Eulerjevega {tevila. S pretokom modelne teko~ine (vode) je bilo preverjeno, da je pri enakem intervalu Reynoldsovih {tevil koeficient izgube ve~ji pri penastem filtru kot pri filtru z direktnimi luknjicami, vendar pa je izhodni koeficient µ za penaste filtre pri enakih pogojih pomembno ni`ji kot pri filtrih z neposrednimi luknjicami. Klju~ne besede: merila podobnosti, modeliranje, tok jeklene taline, kerami~ni filtri, filtriranje jeklene taline 1 INTRODUCTION The theory of physical similarity makes it possible to investigate the regularities of physical and other phenomena with similar behaviour, and it is possible to conclude from the known behaviour of one phenomenon, the behaviour of a second. This theory is based on simi- larity criteria, i. e., on dimensionless quantities (simila- rity numbers), which can substitute in the investigated phenomenon for dimensional physical quantities. The determination of similarity criteria is the basic task for the application of the theory of physical similarity of an investigated problem, since, on the one hand, they reduce the number of variables describing the given problem, and, on the other hand, they specify the similarity relations between the pertinent phenomenon and its model. The theory of physical similarity is used for the determination of similarity criteria and after- wards, for criteria dependencies, three methods of generalised variables, including the dimensional analysis, the analysis of the mathematical model and finally the analysis of the physical model. The objective of this study is to prove that the tech- nology of filtration of metallic melts can be efficiently investigated by the application of the theory of physical similarities using appropriately determined dimen- sionless similarity numbers – i. e., criteria. This study is based on extensive previous and current experimental research 1–6. 2 PRINCIPAL DIMENSIONLESS CRITERIA OF SIMILARITY FOR THE FILTRATION OF METALLIC MELTS The Bernoulli equation is the basis for the primary description of the flow of a real metallic melt through the pouring system for the filling of the relevant casting mould with an inserted ceramic filter. With respect to the pressure losses caused during the flow through the filter, this equation can have the following form: ρ ρ ξ νρw h g p w d 2 2 + + + = const. (1) where w2/2 (Pa) represents the kinetic energy, hg (Pa) represents the positional or gravitational energy of the capacity of a liquid unity, which is determined by gravity, and p (Pa) represents the potential pressure energy, which is usually dependent on the external Materiali in tehnologije / Materials and technology 42 (2008) 4, 175–178 175 UDK 669.14.018:669.18 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(4)175(2008) forces being exerted. The values in the equation are: w = the real liquid flow-rate (m s–1),  = the density of the flowing liquid (kg m–3), h = the real position, the real height of the flowing liquid (m), g = the acceleration due to gravity (m s–2),  = a dimensionless loss coeffi- cient expressing the filter’s hydraulic resistance (–), ν = the kinematic viscosity of the flowing liquid (m2 s–1) and d = the diameter of the filter capillaries with direct holes (m). The pressure of the flowing real liquid p is expressed in a basic unit (Pa) with the basic physical dimension of (kg s–2 m–1). It follows from the Bernoulli equation that the sum of the kinetic, positional, pressure and loss energy of an ideally flowing liquid remains constant, i.e., const in volume unit at each place during its flow. The original Bernoulli equation (1) is transformed by division with the expression w2/2 (Pa) to a dimension- less form 1 2 2 2 2 2 2 + + + =hg w p w wd wρ ξ ν ρ 2 const. (2) which contains the already standard (internationally) adopted dimensionless similarity criteria, namely: the Froude criterion Fr = w2(hg)–1, the Euler criterion Eu = p–1 w–2 and the Reynolds criterion Re = wd–1. With the use of these similarity criteria (numbers) it is possible to express the Bernoulli equation in a dimen- sionless form: 1 2 2 2 2 + + + = Fr Eu w ξ ρRe 2 const. (3) The Froude criterion Fr = w2(hg)–1 expresses the ratio of the inertia and the gravitational forces in the flowing metallic liquid, including its undulation, vortices and also surface phenomena. It is often also called the Froude criterion 1, Fr1, in contrast to the Froude criterion 2, Fr2, expressed in the form Fr2 = w(hg)–1/2, which characterises the so-called "kinetic head". The Euler criterion Eu = p(w2)–1 expresses the ratio of the pressure and the inertia forces. It characterises the loss of pressure during the flowing of a real metallic liquid and influences the hydraulic resistance of the flowing liquid caused by viscous forces (viscosity). The Reynolds criterion Re = wd–1 belongs to the basic hydrodynamic criteria and it characterises the ratio of the inertia and the viscous forces, i.e., the forces caused by the viscosity of the filtered metallic melt. Its absolute value contains information about the basic character of the flowing of viscous liquids, laminar or turbulent, and also the information about the transition from one type of form of flowing to another. The modified dimensionless form of the Bernoulli equation (3) can, in principle, be used for the determi- nation of the loss coefficient , as a general function of dimensionless criteria ξ ρ = − + +⎛ ⎝⎜ ⎞ ⎠⎟ ⎧ ⎨ ⎩ ⎫ ⎬ ⎭ Re 2 1 2 2 2 2 const. w Fr Eu (4) which indicates that during the flowing of a metallic melt through the ceramic filter the loss coefficient is proportional to the Reynolds number, whereas the constant of proportionality is also dependent on the Froude and Euler numbers and on other parameters of the Bernoulli equation. The expression 2 const. (w2)–1 is also dimensionless, since the original physical dimen- sion of this constant is in pascal (Pa). 3 MODELLING OF A METALLIC MELT FLOW THROUGH CERAMIC FILTERS The practical use of the interdependence between the loss coefficient and the used ceramic filter requires that an equation is determined that contains certain specific numeric values. It is possible to determine such an equation by an appropriately chosen and arranged type of modelling. For these purposes the hydraulic charac- teristics of filters with direct holes (strainers) and foam filters were chosen for the modelling. Water served as the modelling liquid, since its viscosity is quantitatively comparable with the viscosity of liquid steel. The measurement of the hydraulic parameters during the flowing of a liquid through ceramic filters was carried out on the water measuring line in the laboratory of the Department of Airplanes and Engines at the Military Academy in Brno. The results of the measurements for two different types of filters are described in detail in the reports7–9: a) For the same type of filter with direct holes (strainer) manufactured by the company Keramtech @aclé, s. r. o., Czech Republic, a filter number of 0217, a diameter of 70 mm and a basic filter height (thickness) L = 10 mm with 19 holes of diameter D = 6 mm, in which a total of six filter-slenderness ratios were measured. The slender- ness ratios were defined as the ratio of the filter height (thickness) L and the diameter of one filter-capillary aperture D. The changes of slenderness were achieved by changes of the height (thickness) of the same type of filters and by their serial ordering (stacking). b) For the same type of foam filter (schaum filter) with the dimensions (90×90×25) mm (8.0 ppi), manu- factured by the company FOSECO, in which two different hole slendernesses were measured. The first series of modelling, the results of which are briefly presented in this study, was focused on the determination of the loss coefficient  and of the filter outlet coefficient µ. The specific energy loss ez was defined for the purposes of the assessment of selected filters by the relation ez = (w2/2) (m2 s–1), which means that the energy loss is directly proportional to the filter loss coefficient (filter specific resistance) and the square of the flow rate of the modelling liquid – the water in the pipeline of the modelling system. The Reynolds number K. STRANSKY ET AL.: APPLICATION OF THE THEORY OF PHYSICAL SIMILARITY ... 176 Materiali in tehnologije / Materials and technology 42 (2008) 4, 175–178 of the investigated filters was defined by the standard relation ReF = w DF hF ν (5) where wF is the flow rate of the liquid in the filter and DhF is the hydraulic diameter of the filter, which is equal to the diameter of one of its apertures (capillaries). The filter outlet coefficient µ was defined as the ratio of the real volume flow of the modelling liquid (water) with respect to the friction losses and the possible contraction of the water jet to the ideal volume flow through the filter, determined from the equations for the flowing of an ideal liquid7. 4 RESULTS OF THE MODELLING OF THE FLOW THROUGH CERAMIC FILTERS In Figure 1 the dependence of the loss coefficient P of both types of ceramic filters, as determined by modelling, on the Reynolds number ReF and the slenderness of filter apertures L/D is shown. The loss coefficient is considered as the local resistance of the filter inserted into the pipeline (index P), and as the resistance corresponding to the state of the modelling liquid (i.e., water at a temperature of 20 °C) in the filter (index F). The two curves in the upper part of the diagram are valid for foam filters, while the group of curves in the bottom part of diagram relates to filters with direct holes (strainers). With increasing Reynolds number and decreasing filter slenderness the coefficient P, expressing the filter’s resistance and the energy loss of the flowing liquid, decreases. However, for the foam filters the decrease of the coefficient P with the decrease of the filter’s slenderness L/D is distinctly larger than for filters with direct holes. The modelling of the filter-outlet coefficient µ is shown in Figure 2. The two curves in the bottom part of the diagram are related to foam filters, while the group of curves in the upper part of the diagram are related to filters with direct holes. The outlet coefficient µ also expresses the contraction of the modelling liquid flowing out of the filter. It is obvious that the foam filters have, in comparison with the filters with direct holes, significantly lower values of the outlet coefficient µ, although the slendernesses of the apertures (capillaries) were otherwise identical. The values of the outlet coefficients µ for the foam filters vary in the interval 0.23 to 0.35, while the values of the outlet coefficients µ for the filters with direct holes are within the interval 0.65 to 0.80, and the Reynolds numbers are very close ReFi = 2 · 103 to 3.5 · 104 (Figure 2). It is worth noting that the Reynolds number ReFi in the diagram shown in Figure 2 is linked to the Reynolds number ReP in the diagram shown in Figure 1 by equation (6) in the following manner: Rep = ReFi p1+ ξ (6) During the modelling measurements the theory of physical similarities was strictly respected and the obtained results were expressed with dimensionless similarity numbers. 5 CONCLUSIONS The results of the modelling measurements shown in Figures 1 and 2 indicate that it is characteristic for both different types of ceramic filters (i.e., filters with direct holes and foam filters) that lower the values of the coefficient P, expressing the filter losses or the resistance to the flowing of the liquid through the filter, correspond to higher values of the coefficient µ = , expressing the contraction of the liquid flow flowing out from the filter aperture (capillary) and vice versa. The higher hydraulic resistance during the flow of the steel melt through the foam filters favourably affects the micro-mechanism of filtration. The hydraulic resistance influences the filtration of metallic liquid through ceramic filters with direct holes (capillaries) to a consi- derably lesser degree. Table 1 summarises the values for K. STRANSKY ET AL.: APPLICATION OF THE THEORY OF PHYSICAL SIMILARITY ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 175–178 177 Figure 2: Dependence of the filter-outlet coefficient on the Reynolds number and the ceramic filters’ slenderness Slika 2: Odvisnost izhodnega koeficienta od Reynildsovega {tevila in od vitkosti filtra Figure 1: Dependence of the coefficient of the local resistance of the ceramic filters on the Reynolds number and the filter’s slenderness Slika 1: Odvisnost koeficienta lokalnega upora keremi~nega filtra of Reynoldsovega {tevila in od vitkosti filtra the kinematic viscosity of the water at the temperature of the modelling of flow through the filter and carbon steel containing the mass fraction of C 0.25 % at the casting temperature. This work was realised within the frame of the projects GA ^R reg. no. 106/06/0393 and EUREKA E!3192 ENSTEEL. 6 REFERENCES 1 M. Píhoda, J. Ba`an, J. Dobrovská, P. Jelínek, Z. Jon{ta, M. Vro`ina: Nové poznatky z výzkumu plynulého odlévání oceli New findings from research of continuous casting of steel, V[B – Technical University of Ostrava, Faculta of Metallurgy and Materials Engineering, Ostrava (2001), 175 pp. ISBN 80-248-0037-3 2 J. Happ, M. G. Frohberg: Untersuchungen zur Filtration von Eisen- schmelzen, Giessereiforschung 23 (1971) 1, 1–9 3 G., F. A. Acosta et al.: Analysis of Liquid Flow through Ceramic Porous Media Used for Molten Metal Filtration, Metallurgical and Materials Transactions B 26B (1995), 159–171 4 J. Rou~ka et al.: Filtrace slévárenských slitin Filtration of casting alloys, Czech Foundry Society – Technological committee, Brno (2000), ISBN 80-02-01389-1 5 P. Joná{: O turbulence On turbulation, In`enýrská mechanika 5 (1998) 2, 89-106 6 J. Ba`an, K. Stránský: Flowing of the melt through ceramic filters, Materiali in Tehnologije / Mater. Tehnol. 41 (2007) 2, 99–102 7 J. Maxa, D. Rozehnal, P. Onderli~ka: Stanovení hydraulických pomr cedítkového a pnového filtru modelováním Determination of hydraulic characteristic of strainer and foam filter by modelling, Military Academy in Brno, (1996), 50 p 8 J. Maxa, D. Rozehnal, P. Onderli~ka: Matematické zpracování výsledk mení cedítkových a pnových keramických filtr získaných modelováním Mathematical processing of results of measurement of strainer and foam ceramic filters obtained by modelling, Military Academy in Brno, (1996), 52 p 9 J. Maxa, D. Rozehnal: Stanovení hydraulických pomr cedítkových filtr modelováním Determination of hydraulic characteristic of strainer filters by modelling, Military Academy in Brno, (1997), 39 p 10 K. Ra`njevi}: Termodynamické tabulky Thermo-dynamic tables. Vydavatestvo technickej a ekonomickej literatúry ALFA, Bratislava (1984), 313 p 11 T. Myslivec: Fyzikáln chemické základy oceláství Physical- chemical principles of steelmaking. SNTL/ALFA, Praha (1971), 448 p K. STRANSKY ET AL.: APPLICATION OF THE THEORY OF PHYSICAL SIMILARITY ... 178 Materiali in tehnologije / Materials and technology 42 (2008) 4, 175–178 Table 1: Comparison of the kinematic viscosity of liquids flowing through the filters – water and steel Tabela 1: Primerjava kinemati~ne viskoznosti teko~in, ki se pretakajo skozi filter – voda in jeklena talina Kinematic viscosity / m2 s–1 Model − water 10 Reality − carbon steel 11 Temperature /°C Viscosity /m2 s–1 Temperature /°C Viscosity /m2 s–1 20 1.55 ⋅ 10–7 1600 6.23 ⋅ 10–7 S. ^AMPELJ ET AL.: PRIPRAVA NANOKOMPOZITA ZA BIOMEDICINSKE APLIKACIJE PRIPRAVA NANOKOMPOZITA ZA BIOMEDICINSKE APLIKACIJE PREPARATION OF NANO-COMPOSITES FOR BIOMEDICAL APPLICATIONS Stanislav ^ampelj1, Darko Makovec1, Luka [krlep2, Miha Drofenik1,3 1Odsek za sintezo materialov, Institut “Jo`ef Stefan”, Jamova 39, SI-1000 Ljubljana, Slovenija 2Zavod za gradbeni{tvo Slovenije, Dimi~eva 12, 1000 Ljubljana, Slovenija 3Fakulteta za kemijo in kemijsko tehnologijo, Univerza v Mariboru, Smetanova ul. 17, SI-2000 Maribor, Slovenija stanislav.campeljijs.si Prejem rokopisa – received: 2007-11-23; sprejem za objavo – accepted for publication: 2008-04-24 Kompozitni nanodelci, ki vsebujejo superparamagnetno maghemitno jedro, prevle~eno s tanko plastjo amorfnega silicijevega oksida so zelo obetaven material za uporabo v biomedicini. Magnetno jedro omogo~a manipulacijo z delci z zunanjim magnetnim poljem, medtem ko pla{~ amorfnega silicijevega oksida omogo~a vezavo razli~nih molekul na njihovo povr{ino. Vezava razli~nih organskih molekul, na primer zdravilnih u~inkovin, zahteva pripravo nanodelcev, ki imajo na povr{ini sloj funkcionalizacijskih molekul z razli~nimi funkcionalnimi skupinami. Funkcionalizacijo nanodelcev smo dosegli s kovalentno vezavo razli~nih silanskih molekul: (3-aminopropil)trietoksisilan (APS) in viniltrietoksisilan (VTS), na njihovo povr{ino. Reakcija je potekla v me{anici etanola, v katerem je bila predhodno raztopljena izbrana silanska molekula, in stabilne vodne suspenzije kompozitnih nanodelcev. Vezavo razli~nih silanskih molekul na povr{ino nanodelcev smo spremljali z elektrokineti~nimi meritvami in s konduktometri~no meritvijo koncentracije molekul na njihovi povr{ini. Izkazalo se je, da lahko ve`emo na delce molekule APS v povr{inski koncentraciji, ki se sklada s koncentracijo silanolnih skupin na povr{ini amorfnega silicijevega oksida. Klju~ne besede: nanodelci, nanokompoziti, silani, zeta-potencial, funkcionalizacija Composite nano-particles of superparamagnetic maghemite particles coated with a thin layer of silica are very promising material for biomedical applications. The magnetic core of the composite nano-particles allows manipulation of particles with external magnetic field while the silica shell allows additional bonding of molecules to the surface. Different organic molecules, such as medical drugs, require nano-particles with a layer of functionalization molecules with different functional groups. The functionalization of nano-particles was achieved with covalent bonding of different silanol molecules: (3-aminopropyl) triethoksysilane (APS) and vinyltriethoksysilane (VTS) to their surface. Reaction took place in a mixture of ethanol with previously dissolved silane and stable aqueous suspension of composite nano-particles. The bonding of different silanol molecules was monitored with electro-kinetic measurements and with conductometric measurements of molecules on the surface. The concentration of APS molecules which can be bondend to the surface of the composite nano-particles is in accordance with the concentration of silanol groups on the surface of silica. Key words: nanoparticles, nanocomposites, silanes, zeta potential, fuctionalisation 1 UVOD V zadnjih letih se veliko pozornosti namenja uporabi magnetnih nanodelcev v medicini. Magnetizem delcev nam omogo~a, da lahko z njmi na daljavo manipuliramo z zunanjim magnetnim poljem, spremljamo njihov polo`aj ali jih segrevamo. Uporabljamo jih tako v diagnosti~ne namene, kot je na primer za pove~evanje kontrasta pri slikanju z NMR-tehniko, kot tudi v terapevtske namene, kot sta na primer magnetna hipertermija in ciljni vnos zdravilnih u~inkovin. Primeren magnetni material za uporabo v medicini je maghemit (-Fe2O3), ki velja za nestrupen material1,2. Pogoj za uporabo nanodelcev v medicini je poleg majhne velikosti in zadovoljivih magnetnih lastnosti tudi njihova nestrupenost in specifi~ne povr{inske lastnosti. ^e `elimo, da je magnetni nanokompozit primeren za uporabo v medicini, je treba na njegovo povr{ino vezati razli~ne biolo{ke u~inkovine. Vezavo u~inkovin na povr{ino nanokompozita dose`emo s funkcionali- zacijskim slojem molekul, ki so vezane na njegovo povr{ino. Ta sloj molekul zagotavlja funkcionalne skupine za kemijsko vezavo razli~nih u~inkovin, hkrati pa prepre~uje aglomeracijo nanodelcev med uporabo. Narava funkcionalizacijskih molekul dolo~a tudi povr{inske lastnosti delcev (povr{inski naboj, polarnost) in s tem njihovo zdru`ljivost z biolo{kimi sistemi3. Za razli~ne biolo{ke uporabe je treba na povr{ino delcev vezati razli~ne biolo{ke molekule, kot so na primer proteini, antigeni ali deli DNA-molekule, kar zahteva plast molekul z razli~nimi funkcionalnimi skupinami. Pomembno je, da so molekule mo~no vezane na povr{ino delcev4,5,6. Ker je povr{ina `elezovega oksida relativno inertna, ne omogo~a mo~ne vezave molekul. Mo~no kovalentno vezavo razli~nih molekul na povr{ini delcev omogo~imo, ~e oksidne delce pre- vle~emo z amorfnim silicijevim oksidom. Silicijev oksid ima namre~ na povr{ini mo~no vezane silanolne OH-skupine. Na povr{inske OH-skupine lahko nadalje kovalentno ve`emo razli~ne funkcionalizacijske mole- Materiali in tehnologije / Materials and technology 42 (2008) 4, 179–182 179 UDK 547:620.1:61 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 42(4)179(2008) kule. Plast silicijevega oksida na oksidnih delcih mora biti neprekinjena in homogena. Debelina plasti mora biti dovolj debela, da za{~iti magnetno jedro nanokompozita, hkrati pa dovolj tanka, da bistveno ne poslab{a njegovih magnetnih lastnosti. Silanolne skupine na povr{ini amorfnega silicijevega oksida niso primerne za vezavo vseh biolo{ko aktivnih molekul. Tako so za vezavo razli~nih molekul potrebne razli~ne skupine, kot sta na primer aminoskupina (NH2) ali karboksilna skupina (COOH). V tem primeru lahko ve`emo na povr{inske silanolne skupine silane, ki nosijo `eleno funkcionalno skupino. Silani so molekule, ki temeljijo na spojini silan, SiH4. Navadno imajo na Si-atom vezane tri etoksidne skupine in eno alkilno verigo, ki se kon~a s funkcionalno skupino. Silani v vodi hidrolizirajo in reagirajo s povr{inskimi silanolnimi skupinami ter tako tvorijo mo~no Si-O-Si-vez s povr{ino silicijevega oksida7,8,9 {e za APS. Pri tem delu smo sistemati~no raziskovali vezavo razli~no funkcionaliziranih silanskih molekul (APS, VTS) na povr{ino kompozitnih nanodelcev iz maghemit- nega jedra in tanke prevleke silicijevega oksida. 2 EKSPERIMENTALNO DELO 2.1 Priprava nanokompozita Sintezni postopek za pripravo kompozitnih nano- delcev je podrobno opisan drugje10. Nanodelce maghemita smo sintetizirali s koprecipitacijo ionov Fe2+ in Fe3+ v vodni raztopini s koncentrirano raztopino amo- nijaka. Sintetizirane nanodelce smo sprali z amonija- kalno raztopino (pH > 10,5). Na delce smo v naslednji stopnji adsorbirali citronsko kislino. Ta deluje kot surfaktant in prepre~i njihovo aglomeracijo. Prekrivanje delcev s silicijevim oksidom je potekalo s hidrolizo tetra- etoksisilana (TEOS). V etanolu raztopljen TEOS smo dodali v stabilno suspenzijo maghemitnih nanodelcev. V suspenziji nastane hidroliza TEOS-a in nukleacije silicijevega oksida na povr{ni maghemitnih nanodelcev. Pomembno je, da v suspenziji ni aglomeratov nano- delcev, saj bi v tem primeru prekrili s silicijevim oksidom aglomerate, in ne posameznih nanodelcev. Tako pripravljene kompozitne nanodelce smo sprali in dispergirali v vodi. Analiza kompozitnih nanodelcev je pokazala, da so iz magnetnega maghemitnega jedra s premerom (13,7 ± 2,9) nm in iz plasti amorfnega silicijevega dioksida z debelino okoli 2 nm. Izra~unana masna plo{~ina nano- kompozita je 94,4 m2/g. Iz magnetnih meritev je razvidno, da ka`ejo kompozitni nanodelci superpara- magnetizem ob relativno visoki nasi~eni magnetizaciji 35 · 10-4 T/g. 2.2 Vezava silanskih molekul na povr{ino kompozitnih nanodelcev Na povr{ino kompozitnih nanodelcev smo vezali naslednje silanske molekule: (3-aminopropil) trietoksi- silan, (CH3CH2O)3Si(CH2)3NH2 (APS) in viniltrietoksi- silan, (CH3CH2O)3SiCHCH2 (VTS) (slika 1). Vezava molekul je potekla v vodem mediju pri temperaturi 50 °C. Stabilni vodni suspenziji (20 mL) z masnim dele`em 0,5 % kompozitnih nanodelcev smo dodali etanol (25 mL), v katerem smo predhodno raztopili silan (0,02 mol). Izra~unan dodatek silana na povr{ino nanokom- pozita je bil v vseh primerih 2 mmol/m2. Za spremljanje prekrivanja povr{ine kompozitnih delcev in dolo~itev koli~ine silana, ki je potrebna za popolno prekritje povr{ine nanokompozita pri prej omenjenih pogojih, smo uporabili APS. Pri tem smo dodatek APS spreminjali od 1 µmol/m2 do 2 mmol/m2. Po 5 h smo delce magnetno lo~ili in jih temeljito sprali z destilirano vodo. Uspe{nost prekritja nanokom- pozita s silani smo spremljali z elektrokineti~nimi meritvami, ki so bile opravljene z zetametrom (Brook- haven Instruments Corp., ZetaPALS). Povr{insko koncentracijo aminoskupin na povr{ini kompozitnih nanodelcev smo dolo~ali s kondukto- metri~no titracijo. Suspenzijo spranih kompozitnih nanodelcev, prevle~enih z APS v vodi, smo titrirali z razred~eno raztopino klorovodikove kisline (HCl). Prevodnost suspenzije je odvisna od koli~ine raztop- ljenih ionov v nosilni teko~ini (vodi). Na za~etku imamo v raztopini le -NH3+ vezane na povr{ini delcev in ione OH– v vodi. Lastni pH suspenzije je 9. Med titracijo pote~e reakcija med klorovodikovo kislino in aminoskupino. Pri tem nastaja nedisociirana molekula vode, ki ne prispeva k prevodnosti raztopine, zato se le-ta bistveno ne spremeni. Ko v raztopini ni ve~ povr{inskih aminoskupin, ki bi bile na voljo za reakcijo s klorovodikovo kislino, se koli~ina ionov z vsakim dodatkom kisline zelo pove~a. Posledica je povi{ana prevodnost raztopine. Ekvivalentna to~ka se dolo~i iz ostre spremembe v naklonu premice odvisnosti prevodnosti od koncentracije dodane HCl. Prevodnost suspenzije smo merili s konduktometrom (Radiometer analytical IONcheck30). 3 REZULTATI IN DISKUSIJA Slika 2 prikazuje elektrokineti~ne meritve za kom- pozitne nanodelce, funkcionalizirane z razli~nimi silani. Sami kompozitni nanodelci ka`ejo v nevtralnem mo~no negativen zeta- potencial, ki je posledica negativno S. ^AMPELJ ET AL.: PRIPRAVA NANOKOMPOZITA ZA BIOMEDICINSKE APLIKACIJE 180 Materiali in tehnologije / Materials and technology 42 (2008) 4, 179–182 APS VTS Slika 1: Shematski prikaz molekul APS in VTS Figure 1: Schematic presentation of APS and VTS molecules nabitih silanolnih OH-skupin na povr{ini. Tako imajo izoelektri~no to~ko (IEP) pri kisli pH-vrednosti okoli 2,5. Iz elektrokineti~nih meritev je razvidno, da je v vseh primerih funkcionalizacije pri{lo do sprememb na povr{ini delcev, kar gre pripisati uspe{ni vezavi silanov na njihovo povr{ino. Po pri~akovanju je sprememba najbolj o~itna, ~e na povr{ino delcev ve`emo APS. Opazen je premik IEP od pH = 2,5 na pH = 9,5. APS ima na koncu alkilne verige aminoskupino, ki je po naravi bazi~na. To pomeni, da je v obmo~ju visokih pH-vrednosti nedisociirana in ima negativen naboj. Ko se pH-vrednost medija premakne na podro~je nizkih pH-vrednosti, se aminoskupina protonira in posledi~no dobimo pozitivno nabito skupino na koncu alkilne verige -NH3+. V primeru vezave VTS na povr{ino nanodelcev se IEP ne spremeni, pa~ pa se pove~a negativni naboj. Vinilna skupina ima zaradi dvojne vezi pove~ano elektronsko gostoto in posledi~no vi{ji negativni naboj. V kislem obmo~ju pH-vrednosti pote~e reakcija med dvojno vezjo in vodo, pri ~emer nastane alkohol. Hidroksilna skupina na koncu alkilne verige je po svojih kemijskih lastnostih podobna silanolnim OH-skupinam, zato se IEP ne premakne. Slika 3 prikazuje elektrokineti~ne meritve za razli~ne dodatke APS, uporabljene v procesu vezave na povr{ino kompozitnih nanodelcev. Mno`ina dodanega APS k nanodelcem je izra`ena v mikromolih APS na kvadratni meter povr{ine nanodelcev. Z vi{anjem dodatka APS se vi{a koncentracija APS, vezana na povr{ini nanodelcev, zato se IEP suspenzije postopno pove~uje. Kon~no pH-vrednost IEP dose`emo pri dodatku APS enakem ali ve~jem od 20 µmol/m2. Sklepamo, da je v tem primeru povr{ina nanodelcev nasi~ena z vezanimi molekulami. Rezultati elektrokemi~nih meritev se skladajo s kon- duktometri~nimi meritvami koncentracije aminoskupin na povr{ini nanodelcev (Tabela 1). Z ve~anjem koli~ine dodanega APS se ve~a tudi izmerjena koncentracija APS na povr{ini nanodelcev. Pri dodatku 20 µmol APS na kvadratni meter povr{ine nanodelcev, se je na povr{ino vezal APS v povr{inski koncentraciji (8,1 ± 0,8) µmol /m2 oziroma 4,8 ± 0,5 molekul APS na kvadratni nanometer povr{ine nanodelcev. Vrednost se dobro ujema s {tevilom silicijevih atomov oziroma silanolnih OH-skupin na povr{ini amorfnega silicijevega oksida, ki je 4,55 molekul OH na kvadratni nanometer silicijevega oksida11. Na vsako od povr{inskih silanolnih skupin se namre~ lahko ve`e po ena molekula APS. Pri vi{jih dodatkih APS je izmerjeno {tevilo aminoskupin na povr{ini ve~je od teoreti~no mo`nega. Vzrok za to je verjetno nastanek polimernih molekul med hidrolizo APS zaradi relativno visoke koncentracije APS v suspenziji. Pri tem nastanejo polimerne molekule ni`jih molekulskih mas, ki se med seboj povezujejo in pri tem nastajajo polimerne molekule vi{jih molekulskih mas. Teh z izpiranjem delcev po funkcionalizaciji verjetno nismo popolnoma izlo~ili. Znano je, da ta proces nastane, ko koncentracija silana prese`e vrednost 3 mmol/L. V primeru dodatka 500 µmol APS na kvadratni meter nanodelcev je bila koncentracija silana v vodni fazi suspenzije 86 mmol/L. Tabela 1: Rezultati konduktometri~nih meritev povr{inske kon- centracije aminoskupin na povr{ini nanodelcev v odvisnosti od dodatka APS uporabljenega v procesu funkcionalizacije. Table 1: Results of conductometric titration measurements for the nano-composite functionalized with different amounts of APS Dodatek APS na povr{ini delcev µmol/m2 1 3 5 20 500 Izmerjena povr{inska koncentracija amino- skupin, camino/(µmol/m2) 3,2 ± 1,5 3,7 ± 1 2,4 ± 1 8,1 ± 0,8 11,0 ± 1 Koncentracija molekul APS na povr{ini delcev, c/nm2 1,9 ± 0,9 2,2 ± 0,6 1,4 ± 0,6 4,8 ± 0,5 6,6 ± 0,6 S. ^AMPELJ ET AL.: PRIPRAVA NANOKOMPOZITA ZA BIOMEDICINSKE APLIKACIJE Materiali in tehnologije / Materials and technology 42 (2008) 4, 179–182 181 Slika 3: Graf elektrokineti~nih meritev za nanokompozit, funkcionali- ziran z razli~nimi dodatki APS Figure 3: Graph of electro kinetics measurements for nano-composite functionalized with different amounts of APS Slika 2: Graf elektrokineti~nih meritev za nanokompozitne delce (polna linija), nanokompozitne delce, funkcionalizirane z APS (pik~asta linija) in nanokompozinte delce, funkcionalizirane z VTS (~rtkana linija) Figure 2: Graph of electro kinetics measurements for nano-composite particles (full line), nano-composite particles functionalized with APS (dotted line) and nano-composite particles functionalized with VTS (dashed line) 4 SKLEP Povr{ino magnetnih nanodelcev smo funkcionalizi- rali z vezavo razli~nih silanskih molekul ((3-amino- propil)trietoksisilan (APS) in viniltrietoksisilan (VTS)) na njihovo povr{ino preko vmesne tanke plasti amorf- nega silicijevaga oksida. Slednji omogo~i s svojimi povr{inskimi silanolnimi OH-skupinami mo~no kova- lentno Si-O-Si-vez silana s povr{ino delcev. Povr{inska koncentracija silanskih molekul. vezanih na povr{ini nanodelcev, se sklada s koncentracijo silanolnih skupin na povr{ini amorfnega silicijevega oksida. 5 LITERATURA 1 Q. A. Pankhurst, J. Connolly, S. K. Jones, J. Dobson, J. Phys. D: Appl. Phys. 36 (2003), R167–R181 2 R. Hiergeist, W. Andrä, N. Buske, R. Hergt, I. Hilger, U. Richter, W. Kaiser, J. Magn. Magn. Mater., 201 (1999), 420–422 3 C. C. Berry, A. S. G. Curtis, J. Phys. D, 36 (2003), R198–R206 4 A. K. Gupta, M. Gupta, Biomaterials, 26 (2005), 3995–4021 5 Y. Ichiyanagi, S. Moritake, S. Taira, M. Setou, J. Magn. Magn. Mater., 310 (2007), 2877–2879 6 M. E. Park, J. H. Chang, Mat. Sci. Eng. C, 27 (2007), 1232–1235 7 Z. Ma, Y. Guan, H. Liu, J. Magn. Magn. Mater. 301 (2006), 469–477 8 X. Liu, J. Xing, Y. Guan, G. Shan, H. Liu, Colloids Surf A, 238 (2004), 127–131 9 K. Woo, J. Hong, J. P. Ahn, J. Magn. Magn. Mater., 293 (2005), 177–181 10 S. ^ampelj, D. Makovec, M. Bele, M. Drofenik, J. Jamnik, Mater. Tehnol., 41 (2007), 103–107 11 R. K. Iler, The chemistry of silica, John Wiley & Sons, New York 1979, 637 S. ^AMPELJ ET AL.: PRIPRAVA NANOKOMPOZITA ZA BIOMEDICINSKE APLIKACIJE 182 Materiali in tehnologije / Materials and technology 42 (2008) 4, 179–182 M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING NUMERICAL MODELLING RAZVOJ KOKILE ZA ORODNA JEKLA Z UPORABO NUMERI^NEGA MODELIRANJA Martin Balcar, Rudolf @elezný, Libor Sochor, Pavel Fila, Ludvík Martínek @AS, a.s., Strojirenska 6, CZ 59171 Zdar nad Sazavou, Czech Republic martin.balcarzdas.cz Prejem rokopisa – received: 2006-10-11; sprejem za objavo – accepted for publication: 2008-03-25 A long lifetime and tools with optimum quality affect the primary production field and require improvements to the technological conditions of tool-steel ingot production. In the design of a newly developed chill-mould shape the specific casting and crystallization conditions of tool steel, large forging ingots are exploited. The casting and solidification processes were modeled numerically by applying the MAGMA software and the ingot shape was optimized with respect to the real solidification conditions, suppressing the ingot’s internal discontinuities and obtaining an acceptable level of structural and chemical homogeneousness. Key words: tool steel, mould, ingot casting, numerical modeling Dolga trajnostna doba in optimalna kakovost vplivata na primarno proizvodnjo jekla in zahtevata izbolj{anja v tehnologiji izdelave jeklenih ingotov. Pri razvoju nove kokile so bili upo{tevani pogoji ulivanja in kristalizacije velikih ingotov iz orodnega jekla za izkovke. Procesa ulivanja in strjevanja sta bila modelirana z MAGMA-programom. Oblika ingota pa je bila optimizirana z upo{tevanjem realnih razmer pri strjevanju, prepre~en je bil nastanek notranjih diskontinuitet in dose`en sprejemljiv nivo strukturne in kemi~ne homogenosti. Klju~ne besede: orodno jeklo, kokila, ulivanje ingotov, numeri~no modeliranje 1 INTRODUCTION In the frame of the traditional production of equip- ment and tools for heavy engineering and metallurgy, large forgings of tool steel are also manufactured at @AS. The tool-steel forgings with good forming properties and tougher requirements in terms of the product’s internal quality can only be forged from ingots with a high internal quality. In earlier papers we have discussed the possible causes for the occurrence of inherent defects in heavy ingots of W. Nr. 1.2343 (X37CrMoV5-1) and W. Nr. 1.2344 (X40CrMoV5-1) tool steels according to EN ISO 4957 and in roll forgings of 8CrMoV or 8Cr3MoSiV steels.1,2. The optimization of the ingot-solidification process in the casting mould, the change of the mould shape and avoiding the formation of the defects in the ingot were achieved using the MAGMA software for the simulation of the casting and solidification process of the 8K8.4 forged ingot with a mass of 7.6 t. The project’s success was achieved with the design and the verification of a new shape for the 8K9.2 chill mould for the casting of tool-steel ingots with a mass of 8.9 t. 2 DEFECTS IN TOOL STEEL INGOTS In ZDAS heavy semi-products for tools are produced from ingots of the 8K series, with weights from 1000 kg to 11700 kg, forged on the CKV 630, CKV 1250 and CKV 1800 presses from the traditional steel grades EN ISO 4957 W. Nr. 1.2842 (90MnCrV8), W. Nr. 1.2842 (X210Cr12), W. Nr. 1.2343 (X37CrMoV51), W. Nr. 1.2344 (X40CrMoV51), W. Nr. 1.2714 (55NiCrMoV7) and special steels for rolls of 8CrMoV, 8CrMoSiV, 8Cr3MoSiV. The increased difficulty of making heavy forgings of tool steels is connected with the specific properties of high-carbon steels, alloyed with chromium, molybdenum and vanadium. Only by ensuring a sufficient forging-reduction ratio of the as-solidified steel, which may have a large number of defects in the final solidification areas in the axial part of the ingot, can a high internal quality of the forging be achieved. To achieve a high reduction ratio it is necessary to ensure a high forming rate and considerable deformation levels per compression. Tool-steel forming with the CKV 1800 press is limited by the ingot reduction that can be achieved with the pressing force. For this reason, the deformation rate depends strongly on the temperature. A poor forging process may lead, in the case of the bar forgings of tool steels with a large diameter (>200 mm), to a low degree of deformation in the axial part of the forging. With respect to the requirements for the tool’s service life and the resistance to considerable dynamic stressing, it is necessary to attain a high internal homogeneity without critical defects. The internal quality of the forgings is checked for crack and cavity occurrence by using ultrasonic examination. Materiali in tehnologije / Materials and technology 42 (2008) 4, 183–188 183 UDK 669.14.018.252:669.18:519.68 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 42(4)183(2008) The analysis of the tool-steel ingot and the forging carried out in connection with the extensive and repeated occurrence of axial defects in forgings has led to interesting findings for the 8K8.4 ingot’s solidification, crystallization and forming. In spite of the exploitation of the limited possibilities of material deformation on the CKV 1800 press, the metallographic examinations of the forgings have not shown an insufficient forging reduction in the ingot axis, while ingot examinations have pointed out the causes of the defect occurrence at the forging axis. Figure 1 shows an image of an 8K8.4 ingot of 8CrMoSiV steel in cross-section at 1/2 ingot body height after a liquid penetrant test, with the considerable occur- rence of regular cavities and shrinkage porosity at the ingot axis. The defects are confirmed to the ingot’s vertical section, Figure 2, with imaging of the distribu- tion and the shape of the defects throughout the vertical section of the ingot’s upper part below the ingot’s head. 3 NUMERICAL SIMULATION OF 8K8.4 INGOT CASTING AND SOLIDIFICATION The conditions for simulations of the existing shape and the design of the optimized shape of the steel chill mould were selected on the basis of practical results and experience of forging the 8K8.4 ingot of 8CrMoSiV tool steel. From the MAGMA database, the GS80CrMo steel was selected because it is very similar to the steel grade investigated in terms of its chemical composition, which is shown for both steels in Table 1. The solidification data were verified on a model of the current polygonal ingot type 8K8.4 of weight 7600 kg. The results of the numerical simulation were obtained as a set of graphical outputs from the MAGMA database and examined in compliance with the practical findings of the ingot’s internal quality assessment. On the basis of the analysis, the following parameters for the simulation of the steel ingot’s casting and solidification (marking according to the MAGMA database) were determined and selected for further work: SOLTIME s LIQTOSOL s NIYAMA 1 PRINCIPAL STRES MPa M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... 184 Materiali in tehnologije / Materials and technology 42 (2008) 4, 183–188 Table 1: Chemical composition in mass fractions (w/%) of the 8CrMoSiVsteel and the GS80CrMo steel for the simulation from the MAGMA database Tabela 1: Kemi~na sestava jekla 8CrMoSiV in jekla GS80CrMo iz baze podatkov MAGMA, ki je bila uporabljena pri simulaciji Figure 2: Longitudinal section of the part below the ingot head. Image of the liquid penetrant test and detail of the axial part of the 8CrMoSiV steel ingot Slika 2: Vzdol`ni prerez pod glavo ingota. Posnetek preiskave s teko~im penetrantom in detajl aksialnega dela ingota iz jekla 8CrMoSiV Figure 1: Cross-section at half ingot body height. Image of liquid penetrant test and detail of the axial part of the 8CrMoSiV steel ingot Slika 1: Pre~ni prerez pri polovici vi{ine ingota. Posnetek preiskave s teko~im penetrantom in detajl aksialnega dela ingota iz jekla 8CrMoSiV MISES MPa w(C) segregation % Figure 3 shows a graphical representation of the numerical simulation results of the 8K8.4 ingot of the GS80CrMo (8CrMoSiV) steel grade. The SOLTIME parameter in the figure shows the evolution of the solidus temperature zone and of the ingot solidification with its dependence on time. The melt holding time at the liquidus and solidus temperatures interphase dividing line, LIQTOSOL in Figure 4, indicates that the metal in the ingot body axis remains at the liquidus–solidus interphase dividing line for a longer time than at the part below the ingot head. The criterion for micro-shrinkage occurrence – NIYAMA – is shown in Figure 5. The vertical axial section through the ingot body in the Niyama criterion representation in scale (0–1) shows the risk points for the occurrence of micro-shrinkage. The Niyama criterion, decisive for the occurrence of micro-shrinkage, is defined as G/ T ( )/ / –K s mm1 2 1 2 1⋅ ⋅ where G/(K·mm–1) is the temperature gradient and T/(K·s–1) is the cooling rate. The critical value of the Niyama criterion, decisive for micro-shrinkage occurrence in the castings is 0.775 K1/2·s1/2·mm–1. Accordingly, the area of overvalues for the critical limit of the Niyama criterion reflects the propensity for micro-shrinkage formation 3. In the conditions of the steel ingots, the Niyama criterion determines the appearance of the material cavities and the micro-porosity. It is, however, not possible to determine with more precision the critical Miyama value on the basis of the available results. When representing the Niyama criterion in scale (0–1), a zone with the values G/ T to 0.5 K1/2·s1/2·mm–1 in the axial part of the ingot in Figure 3, some shrinkage porosity and cavities may form in the area at approximately 30 % of the body height. The principal stress parameter in Figure 5 shows the relative local stressing arising from volume changes in the course of a steel ingot’s solidification. In the cross-section at 1/2 the ingot body’s height, the zone of relatively high local stresses in the ingot body middle part is evident. The occurrence of high stressing in the ingot can also be related to the ingot’s ability to provide the fluid phase in different solidification areas. The MISES parameter in Figure 6 shows the relative stress that is used to compare the triaxial stress state with the uniaxial stress state (with the rupture test). The stress values can be compared with the yield and the ultimate strength obtained with the tensile test. The carbon concentration change throughout the ingot cross-section expressed by the concentration zones in Figure 8, does not show the significant non-mixing expected from the ingot’s solidification process. The results of the numerical simulation agree sufficiently well with the assessed internal quality of the experimental ingots and confirmed the conformity of the theoretical calculations and the practical experience. They also point out some possible causes for the occurrence of axial defects in the ingot and the forging of the tool steel. The result of the simulations and verifications of the real ingot’s internal quality was the starting stage for the design and construction of a new mould shape. M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... Materiali in tehnologije / Materials and technology 42 (2008) 3, 183–188 185 Figure 3: Ingot solidification time 8K8.4 (SOLTIME, s) Slika 3: ^as strjevanja ingota 8K8.4 (SOLTIME, s) Figure 4: Time Tlikvidus – T solidus 8K8.4 (LIQTOSOL, s) Slika 4: ^as Tlik. − Tsol. za ingot 8K8.4 (LIQTOSOL, s) Figure 5: Niyama criterion 8K8.4 (NIYAMA 1) Slika 5: Niyama−kriterij za ingot 8K8.4 (Niyama1) 4 DESIGN AND VERIFICATION OF THE INGOT 8K9.2 The design of the new shape of the mould and ingot evolved gradually with the comparison of the results of the calculations for a number of modifications to the mould’s geometry. It was based on the present polygonal ingot – type 8K8.4, with a weight of 7600 kg. The main geometrical changes were for the slenderness and the taper of the ingot. The shape was finally modified to be suitable for a simulation of the solidification of the polygonal ingot 8K9.2 of weight of 8850 kg. The comparison between the basic parameters of the ingot 8K8.4 and the new shape design of the ingot 8K9.2 is shown in Table 2. The results of the numerical simulation of casting and solidification of the ingot 8K9.2 of steel GS80CrMo (8CrMoSiV) evaluated in terms of selected parameters in a similar way as for the original ingot, are shown in Figures 9 to 14. It is clear from Figure 9 that the change to the mould’s geometry has positively changed the shape and course of the curves of the time zones of the solidus temperature. In Figure 10 the positive effect of the M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... 186 Materiali in tehnologije / Materials and technology 42 (2008) 4, 183–188 Figure 6: Relative local stress 1/2 ingot height 8K8.4 (PRINCIPAL STRESS, MPa) Slika 6: Relativna lokalna napetost pri polo- vici vi{ine ingota 8K8.4 (Glavna napetost, MPa) Figure 7: Relative stress 1/2 ingot height 8K8.4 (MISES, MPa) Slika 7: Relativna napetost pri polovici vi{ine ingota 8K8.4 (Mises, MPa) Figure 8: Carbon concentration 1/2 ingot height 8K8.4 (w(C)/%) Slika 8: Koncentracija ogljika v masnih dele- `ih pri polovici vi{ine ingota 8K8.4 (w/C)/%) Figure 9: Ingot solidification time 8K9.2 (SOLTIME, s) Slika 9: ^as strjevanja ingota 8K9.2 (SOL- TIME, s) Figure 10: Time Tliquidus – T solidus 8K9.2 (LIQTOSOL, s) Slika 10: ^as Tlik. − Tsol. za ingot 8K9.2 (LIQTOSOL, s) Figure 11: Niyama criterion 8K9.2 (NIYAMA 1) Slika 11: Niyama−kriterij za ingot 8K9.2 (Niyama1) Table 2: Basic parameters of the ingot 8K8.4 and the design of the ingot 8K9.2 Tabela 2: Osnovni parametri za ingot 8K8.4 in na~rt za ingot 8K9.2 enlargement of the ingot taper on the time of the melt persistence the phase-to-phase interface liquidus and solidus is shown. A comparison with Figures 3 and 4 shows a significant improvement of the feeding of the fluid metal from the head into the body of the ingot. On the basis of the result of the simulation of the casting and the solidification an improvement to the internal quality of the ingot can be expected. As shown in the ingot section in Figure 11, and with a difference compared to the ingot 8K8.4, no axial cavities and shrinkage porosities formed during the solidification of the ingot 8K9.2. In Figure 12 and Figure 13 are some interesting changes compared with the original ingot shape. The relative proportional local stress (principal stress) due to volume changes during the solidification of the ingot is reduced considerably. It is assumed that this change is connected with the improved flowability of the fluid metal to the phase-to-phase interface and/or to the point of steel solidification in the ingot. The interrelation between the relative stress and the proportional local stress parameter significantly increases the homogeneity of the stress distribution over the ingot cross-section. Thus, it is logical to expect a lower level of internal stresses in the ingot and an improved homogeneity of the steel’s microstructure. The effect of a change of the mould on the un-mixing and segregation processes is evident from the images in Figure 14. The chemical heterogeneity of the ingot 8K9.2 in terms of carbon concentration change over the ingot body cross-section is probably related to the lengthening of the solidification time. This time in the axis and in the middle part of the ingot body, as shown in Figure 3, is 9230 s for the ingot 8K8.4 and 11036 s for the ingot 8K9.2, and in Figure 9 the solidification time for the new ingot shape is increased by 806 s, i.e., by 19.6 % compared to the initial ingot shape. Based on the achieved improvements in terms of internal quality for the simulated ingot shape, two moulds of shape 8K9.2 were manufactured and their suitability for achieving better internal quality for ingots and forgings of the steel processed in the secondary metallurgy equipment in @AS was tested. The tests were carried out with the steel EN ISO 4957 W. Nr. 1.2344 (X40CrMoV51) and two ingots of 8K9.2 were cast. One ingot was then processed by forming on the CVK 1800 press, using open-die forging technology, to a bar of 350 mm in diameter with the maximum material yield. Of crucial importance for the evaluation of the effect of the change of the mould shape is the forging’s internal quality, which is tested ultrasonically according to the standard SEP 1921. These tests showed that in terms of internal defects the experimental forging from the ingot 8K9.2 of steel EN ISO 4957 W. Nr. 1.2344 (X40CrMoV51) conformed with the levels A/a, B/b, C/c, D/d, E/e. For tool steels, the levels D/d and E/e are acceptable. M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... Materiali in tehnologije / Materials and technology 42 (2008) 4, 183–188 187 Figure 12: Proportional local stress 1/2 height of the ingot 8K9.2 (PRINCIPAL STRES, MPa) Slika 12: Relativna lokalna napetost pri polovici vi{ine ingota 8K9.2 (Glavna nape- tost, MPa) Figure 13: Relative stress 1/2 height of the ingot 8K9.2 (MISES, MPa) Slika 13: Relativna napetost pri polovici vi{ine ingota 8K9.2 (Mises, MPa) Figure 14: Carbon concentration 1/2 height of the ingot 8K9.2 (w(C)/%) Slika 14: Koncentracija ogljika v masnih dele`ih pri polovici vi{ine ingota 8K9.2 (w(C)/%) Table 3: Ingot utilization – EN ISO 4957 W. Nr. 1.2344 (X40CrMoV5-1) steel forging and internal quality tested ultra- sonically per SEP 1921 Tabela 3: Uporaba ingota – EN ISO 4957 W. Nr. 1.2344 (X20CrMoV5-1) jekleni odkovek in podatki o notranji kakovosti, presku{eni z ultrazvokom po SEP 1921 From the verifications carried out and the comparison of results of the statistical data on the earlier production of forgings of steel EN ISO 4957 W. Nr. 1.2344 (X40CrMoV51) shown in Table 3, it is evident that a considerable increase of the ingot yield was achieved after the introduction of the new mould shape. The bar forging length complied at all levels with the requirements of the ultrasonic test for the ingot volume yield of 82 %. The difference to 100 % could not be used because of the crack extension from the ingot head during the forming process and also for the waste in the ingot foot due to the material flow and the loss during squaring up the face of the forging with machine cutting. Based on the result of the first assessment it is possible to conclude that an improved internal quality of the ingot 8K9.2 in terms of ultrasonic tests of the internal quality of the tool-steel forging was achieved. This conclusion is confirmed with the results achieved with real shop orders for the tool-steel forgings of heavy bars and blocks with a weight up to 7 t, a bar diameter up to 600 mm and the height of the block up to 500 mm. 5 CONCLUSIONS The MAGMA software was used for the design of a new shape for the 8K9.2 mould. In spite of the possible modeling errors, it is possible to conclude, based on the results achieved, that a great improvement in the quality of forgings has been attained as a direct result of the change in the geometry of the forging ingot of type 8K. The development work will be continued with the aim of a further increase in the quality of the ingots and forgings of selected tool-steel brands and of the verification of the results of a numerical simulation in terms of the internal structure and of the chemical homogeneity related to the solidification and the forming of the ingot. Acknowledgements The presented results are part of the TANDEM program of the FT–TA/061 project. The project was implemented based on state resources with the financial support of the Ministry of Industry and Trade of the Czech Republic. 6 REFERENCES 1 Martínek, L., Balcar, M., Novák, J., Sochor, L.: Rozbor vnitních necelistvostí ingotu z nástrojové oceli. 6. Mezinárodné metalurgické sympozium. Rájecké Teplice, 2003 2 Martínek, L., Balcar, M., Novák, J., Sochor, L.: Vnitní struktura ingotu z nástrojové oceli. 20. konference Teorie a praxe výroby a zpracování oceli, Ro`nov. Tanger, 2003 3 Carlson, K.D., Shouzhu Ou, Hardin, R., Beckerman, Ch.: Develop- ment of New Feeding – Distance Rules Using Casting Simulation: Part I. Methodology. http://www.engineering.uiowa.edu/ becker/ documents.dir/ FeedingPart1.pdf M. BALCAR ET AL.: THE DEVELOPMENT OF A CHILL MOULD FOR TOOL STEELS USING ... 188 Materiali in tehnologije / Materials and technology 42 (2008) 4, 183–188