KOVINE ZLITINE TEHNOLOGIJE METALS ALLOYS TECHNOLOGIES <0 s ko cn o Glavni urednik / Editor: F. Vodopivec, IMT Ljubljana, Slovenija Gostujoči urednik / Guest Editor: M. Jenko, IMT Ljubljana, Slovenija Izdajatelji / Publishers: Inštitut za kovinske materiale in tehnologije Ljubljana, ACRONI Jesenice, Institut Jožef Štefan, IMPOL Slovenska Bistrica, Kemijski inštitut Ljubljana, Koncem Slovenske železarne, Metal Ravne, Talum Kidričevo KOVINE LETNIK STEV. LJUBLJANA NOV.-DEC. ZLITINE 30 6 1996 TEHNOLOGIJE VOLUME NO. SLOVENIJA Navodilo avtorjem Prosimo avtorje, da pri pripravi rokopisa za objavo članka dosled-noupoštevajo naslednja navodila: - Članek mora biti izvirno delo, ki ni bilo v dani obliki še nikjer objavljeno. Deli članka so lahko že bili podani kot referat. - Avtor naj odda članek oz. besedilo napisano na računalnik z urejevalniki besedil: - VVORDSTAR, verzija 4, 5, 6, 7 za DOS - WORD za DOS ali WINDOWS. Če avtor besedila ne more dostaviti v prej naštetih oblikah, naj pošlje besedilo urejeno v ASCII formatu. Prosimo avtorje, da pošljejo disketo z oznako datoteke in računalniškim izpisom te datoteke na papirju. Formule so lahko v datoteki samo naznačene, na izpisu pa ročno izpisane. Celoten rokopis članka obsega: - naslov članka (v slovenskem in angleškem jeziku), - podatke o avtorju, - povzetek (v slovenskem in angleškem jeziku), - ključne besede (v slovenskem in angleškem jeziku), - besedilo članka, - preglednice, tabele, - slike (risbe ali fotografije), - podpise k slikam (v slovenskem in angleškem jeziku), - pregled literature. Članek naj bi bil čim krajši in naj ne bi presegal 5-7 tiskanih strani, pregledni članek 12 strani, prispevek s posvetovanj pa 3-5 tiskanih strani. Obvezna je raba merskih enot, ki jih določa zakon o merskih enotah in merilih, tj. enot mednarodnega sistema SI. Enačbe se označujejo ob desni strani besedila s tekočo številko v okroglih oklepajih. Preglednice (tabele) je treba napisati na posebnih listih in ne med besedilom. V preglednicah naj se - kjer je le mogoče - ne uporabljajo izpisana imena veličin, ampak ustrezni simboli. Slike (risbe ali fotografije) morajo biti priložene posebej in ne vstavljene (ali nalepljene) med besedilom. Risbe naj bodo izdelane praviloma povečane v merilu 2:1. Za vse slike po fotografskih posnetkih je potrebno priložiti izvirne fotografije, ki so ostre, kontrastne in primerno velike. Vsi podpisi k slikam (v slovenskem in angleškem jeziku) naj bodo zbrani na posebnem listu in ne med besedilom. V pregledu literature naj bo vsak vir oštevilčen s tekočo številko v oglatih oklepajih (ki jih uporabljamo tudi med besedilom, kadar se želimo sklicevati na določeni literarni vir). Vsak vir mora biti opremljen s podatki, ki omogočajo bralcu, da ga lahko poišče: knjige: - avtor, naslov knjige, ime založbe in kraj ter leto izdaje (po potrebi tudi določene strani): H. Ibach and H. Luth, Solid State Physics, Springer, Berlin 1991, p. 245 članki: - avtor, naslov članka, ime revije in kraj izhajanja, letnik, leto, številka ter strani: H. J. Grabke, Kovine zlitine tehnologije, 27, 1993 ,1-2, 9 Avtorji naj rokopisu članka priložijo povzetek v omejenem obsegu do 10 vrstic v slovenskem in angleškem jeziku. Rokopisu morajo biti dodani tudi podatki o avtorju: - ime in priimek, akademski naslov in poklic, ime delovne organizacije v kateri dela, naslov stanovanja, telefonska številka, E-mail in številka fax-a. Uredništvo KZT - odloča o sprejemu članka za objavo, - poskrbi za strokovne ocene in morebitne predloge za krajšanje ali izpopolnitev, - poskrbi za jezikovne korekture. Instructions to Authors Authors are kindly requested to prepare the manuscripts accord-ing to the follovving instructions: - The paper must be original, unpublished and properly prepared for printing. - Manuscripts should be typed vvith double spacing and vvide margins on numbered pages and should be submitted on flop-py disk in form of: - VVORDSTAR, version 4, 5, 6, 7 for DOS, - VVORD for DOS or VVINDOVVS, - ASCII text vvithout formulae, in vvhich čase formulae should be clearly vvritten by hand in the printed copy. Preparation of Manuscript: - the paper title (in English and Slovenian Language)* - author(s) name(s) and affiliation(s) - the text of the Abstract (in English and Slovenian Language)* - key vvords (in English and Slovenian Language)* - the text of the paper (in English and Slovenian Language)* -tables (in English Language) - figures (dravvings or photographs) - captions to figures (in English and Slovenian Language)* - captions to tables (in English) - acknovvledgement - references * The Editorial Board will provide for the translation in Slovenian Language for foreign authors. The length of published papers should not exceed 5-7 journal pages, of revievv papers 12 journal pages and of contributed papers 3-5 journal pages. The international system units (SI) should be used. Equations should be numbered sequentially on the right-hand side in round brackets. Tables should be typed on separate sheets at the end of manuscript. They should have a descriptive caption explaining dis-played data. Figures (dravvings or photographs) should be numbered and their captions listed together at the end of the manuscript. The dravvings for the line figures should be tvvice the size than in the print. Figures have to be original, sharp and well contrasted, enclosed separately to the text. References must be typed in a separate reference section at the end of the manuscript, vvith items refereed too in the text by numerals in square brackets. References must be presented as follovvs: - books: author(s), title, the publisher, location, year, page num-bers H. Ibach and H. Luth, Solid State Physics, Springer. Berlin 1991, p. 245 - articles: author(s), a journal name, volume, a year, issue num-ber, page H. J. Grabke, Kovine zlitine thenologije, 27, 1993 , 1-2, 9 The abstract (both in English and in Slovenian Language) should not exceed 200 vvords. The title page should contain each author(s) full names, affiliation vvith full address, E-mail number, telephone and fax number if available. The Editor - will decide if the paper is accepted for publication, - will take care of the refereeing process, - language corrections. The manuscripts of papers accepted for publication are not re-turned. Rokopisi člankov ostanejo v arhivu uredništva Kovine zlitine tehnologije. KOVINE ZLITINE TEHNOLOGIJE METALS ALLOYS TECHNOLOGIES KOVINE ZLITINE TEHNOLOGIJE Izdajatelj (Published for): Inštitut za kovinske materiale in tehnologije tjubljana Soizdajatelji (Associated Publishers): SŽ ŽJ ACRONI Jesenice, IMPOL Slovenska Bistrica, Institut Jožef Štefan, Kemijski inštitut Ljubljana, Koncem Slovenske Železarne, Metal Ravne, Talum Kidričevo Izdajanje KOVINE ZLITINE TEHNOLOGIJE sofinancira: Ministrstvo za znanost in tehnologijo Republike Slovenije (Journal METALS ALLOYS TECHNOLOGIES is financially supported by Ministrstvo za znanost in tehnologijo, Republika Slovenija) Glavni in odgovorni urednik (Editor-in-chief): prof. Franc Vodopivec, Inštitut za kovinske materiale in tehnologije Ljubljana, 1000 Ljubljana, Lepi pot 11, Slovenija Urednik (Editor): mag. Aleš Lagoja Tehnični urednik (Technical Editor): Jana Jamar Lektorji (Linguistic Advisers): dr. Jože Gasperič in Jana Jamar (slovenski jezik), prof. dr. Andrej Paulin (angleški jezik) Uredniški odbor (Editorial Board): doc. dr. Monika Jenko, prof. Jakob Lamut, prof. Vasilij Prešeren, prof. Jože Vižintin, prof. Stane Pejovnik, dipl. ing. Sudradjat Dai, Jana Jamar Mednarodni pridruženi člani uredniškega odbora (International Advisory Board): prof. Hans Jurgen Grabke, Max-Planck-lnstitut fur Eisenforschung, Dusseldorf, Deutschland prof. Thomas Bell, Faculty of Engineering School of Metallurgy and Materials, The University of Birmingham, Birmingham, UK prof. Jožef Zrnik, Technicka Univerzita, Hutnicka fakulteta, Košice, Slovakia prof. Ilija Mamuzič, Sveučilište u Zagrebu, Hrvatska prof. V. Lupine, Istituto per la Technologia dei Materiali Metallici non Tradizionali, Milano, Italia prof. Gunther Petzov, Max-Planck-lnstitut fur Metallforschung, Stuttgart, Deutschland prof. Hans-Eckart Oechsner, Universitat Darmstadt, Deutschland Izdajateljski svet (Editorial Advisory Board): prof. Marin Gabrovšek, prof. Blaženko Koroušič, prof. Ladislav Kosec, prof. Alojz Križman, prof. Tatjana Malavašič, dr. Tomaž Kosmač, prof. Leopold Vehovar, prof. Anton Smolej, dr. Boris Ule, doc. dr. Tomaž Kolenko, dr. Jelena Vojvodič-Gvardjančič Članki objavljeni v periodični publikaciji KOVINE ZLITINE TEHNOLOGIJE so indeksirani v mednarodnih sekundarnih virih: (Articles published in journal are indexed in international secondary periodicals and databases): - METALS ABSTRACTS - ENGINEERED MATERIALS ABSTRACTS - BUSINESS ALERT ABSTRACTS (STEELS, NONFERROUS, POLYMERS, CERAMICS, COMPOSITES) - CHEMICAL ABSTRACTS - ALUMINIUM INDUSTRY ABSTRACTS - REFERATIVNYJ ŽURNAL: METALLURGIJA Naslov uredništva (Editorial Address): KOVINE ZLITINE TEHNOLOGIJE IMT tjubljana Lepi pot 11 1000 Ljubljana, Slovenija Telefon:+386 61 125 11 61 Telefax: +386 61 213 780 Žiro račun: 50101-603-50316 IMT pri Agenciji Ljubljana Na INTERNET-u je revija KOVINE ZLITINE TEHNOLOGIJE dosegljiva na naslovu: http : // www. ctk. si /kovine/ (INTERNET LINK: http://www.ctk.si/kovine/) Elektronska pošta (E-mail): cobissimtlj @ ctklj.ctk.si Oblikovanje ovitka: Ignac Kofol Tisk (Print): Tiskarna PLANPRINT, Ljubljana Po mnenju Ministrastva za znanost in tehnologijo Republike Slovenije št. 23-335-92 z dne 09.06.1992 šteje KOVINE ZUTINE TEHNOLOGIJE med proizvode, za katere se plačuje 5-odstotoi dg^ek od prometa proizvodom Laudation in honour of Professor Dr. EVanc Vodopivec on the occasion of his 65th birthday Professor Dr. Franc Vodopivec, scientific councillor, former director of Institute of Metals and Tech-nology and member of the State Council of Republic Slovenia is celebrating his 65th birthday. This birthday is the occasion to look at the background and the development of this well known scientist and at the influence which his research work has in the field of elaboration, transformation and use of metals and alloys in Slovenia and abroad. F. Vodopivec was born in Rakitnik, a small village in the former Italy on 8th October 1931. After ftnishing with distinction the secondary school education, he studied Metallurgy at the University of Ljubljana. In 1956 he passed the final examinations and second degree thesis as the first of his class. During the university study he was for three years assistant-student for lectures of Mechanics and Kinematics. In 1956 he joined Metallurgical Institute, present Institute of Metals and Technology in Ljubljana directed by the founder Professor Ciril Rekar. After the military service 1958/59 he re-ceived through the International Agency of Atomic Energy in Vienna a scholarship from the French Government. Working in the Institute de Recherche de la Siderurgie, in St.Germain en Laye, France, from 1960 to 1962 he prepared his Dr.-thesis and graduated in 1962 at the University of Pariš, France with the thesis: Study of the behaviour of arsenic and phosphorous by selective oxidation of iron alloys vvith low contents of both elements. He returned in 1962 to the Metallurgical Institute and worked as founder and head of the Laboratory for Metalography to 1972, head of Technology Department to 1978, assistant director to 1990 and director from 1990 to April 1996 when he retired. In 1992, Professor Vodopivec was elected in the Council State of Republic Slovenia by the community of researchers and engineers. He is the editor-in-chief of Slovenian scientific journal Metals Alloys Technologies since 1994. Professor Vodopivec is full of development špirit and creative ideas. He has been doing research work on the behaviour of metals in oxidative atmosphere, microstructure characterization of metals by optical and electron microscopy, electron probe analysis, mechanical testing; behaviour of material in use at medium and high temperature, hot and cold working of metals, recovery, recrystallization and grain growth. His present research interest includes: ductile permanent magnet alloys, non oriented electrical steel sheets, grain growth induced by selective surface segregation, topology of microstructure and behaviour of metals in use. Professor Vodopivec has published over 150 papers in intemational journals and conferences and 240 papers in Slovenian journals and conferences on topics of science, technology and use of metals and alloys. Professor Vodopivec has been supervisor to several Ph.D. and Master Degree students at the Univer-sities of Ljubljana, Maribor, Belgrade and Zagreb. He is also very active in the intemational academic field. He was a chairman of intemational scientific conferences and project evaluator in EU COST actions. He is the president of Slovenian Society of Materials, member of executive council of Slovenian Vacuum Society, member of Slovenian Electron and Microelectronics Society, Slovenian Society of Chemistry, Historical Society of Ljubljana, chairman of the R&D group of the Slovenian Association of Engineers, chairman of annual Conferences on Materials and Technologies from 1990 to present, and member of Vacuum Metallurgy scientific division of IUVSTA - International Union for Vacuum Science, Technique and Applications. He wrote in Slovenian newspaper several tens of articles of industrial and research policy. In 1978 he was awarded by the Boris Kidrič Foundation Award and in 1984 the Boris Kidrič State Award for Science. His many projects were supported by 21 industrial societies and associations in Slovenia and the former Yugoslavia from Metallurgy, over mechanical industry to power stations as well as the Slovenian and the Yugoslav governments. He was involved also in the projects of international cooperations EU RD actions and USA- Slovenia projects. He prepared forensic analysis of several industrial failures vvhich qualified Slovenian societies to win arbitration for retributions of damages from foreign companies suppliers of industrial equipment. His colleagues hope very much that he vvill instead of the retirement, take part in discussions, lectures and publications. Most of ali we would like to wish him and his family many years to come in good health. Monika Jenko Vsebina - Contents ZNANSTVENI PRISPEVKI - SCIENTIFIC PAPERS Kovinski materiali - Metallic Materials Surface and Grain Boundary Segregation of Antimony and Tin - Effects on Steel Properties Segregacija antimona in kositra na površini in po mejah zrn - Vpliv na lastnosti jekel H. J. Grabke........................................................................ 483 Applications of Surface Analytical Techniques in Corrosion Research (Mainly High Temperature Corrosion) Uporaba površinskih analiznih tehnik v raziskavah korozije H. Viefhaus......................................................................... 497 Aluminium and Magnesium Based Metal Matrix Composites Kompoziti na osnovi Al in Mg K. U. Kainer........................................................................ 509 Microstructural Considerations Limiting the Mechanical Properties of HSLA Steel Mikrostrukturne omejitve mehanskih lastnosti HSLA jekel L'. Parilak ......................................................................... 517 Predicting of Reactions During Carburization and Decarburization of Steels in Controlled Atmospheres Napovedovanje reakcij, ki potekajo med naogljičenjem in razogljičenjem jekla v kontroliranih atmosferah B. Koroušič, M. Stupnišek............................................................. 52.1 Fusion of Low Carbon Steel Scrap in the Middle Carbon Steel Melt Taljenje niskougljičnog čeličnog otpatka u talini srednje ugljičnog čelika V. Grozdanič........................................................................ 527 Equilibrium Grain Boudary Segregation of Antimony in Iron Base Alloys Ravnotežna segregacija antimona po mejah zrn v zlitinah železa in antimona R. Mast, H. Viefhaus, M. Lucas, H. J. Grabke ............................................ 531 Sn Influence on the Recrystallization of Non-Oriented Electrical Sheet Vpliv Sn na rekristalizacijo neorientirane elektro pločevine M. Godec, M. Jenko, R. Mast, F. Vodopivec, H. J. Grabke, H. Viefhaus ....................... 539 Corrosion Resistance of NdDyFeB Basic Alloys Korozijska obstojnost osnovnih zlitin NdDyFeB S. Kobe Beseničar, L. Vehovar, B. Saje .................................................. 545 Some Aspects of Impurity Grain Boundary Segregation in Low Alloy Cr-Mo-V Steel Segregacije nečistoč v nizko legiranih Cr-Mo-V jeklih J. Janovec, V Magula, P. Sevc ......................................................... 551 Mechanical Properties of High Temperature Vacuum Brazed HSS on Structural Carbon Steel vvith Simultaneous Heat Treatment Mehanske lastnosti visokotemperaturno vakuumsko spajkanih in istočasno toplotno obdelanih spojev V. Leskovšek, D. Kmetic, B. Šuštaršič................................................... 557 Discontinuous Al-SiC Composites Formed by a Low Cost Chemically Activated Infiltration Technique Pridobivanje in kemijska infiltracija poroznih SiC vzorcev Al-Si talino V M. Kevorkijan .................................................................... 565 Letno kazalo - Index................................................................ 573 Surface and Grain Boundary Segregation of Antimony and Tin - Effects on Steel Properties Segregacija antimona in kositra na površini in po mejah zrn - vpliv na lastnosti jekel H. J. Grabke1, Max-Planck-lnstitut, Dusseldorf, Germany Dedicated to Prof. Dr. F. Vodopivec on the occasion of his 65th birthday. Prof. dr. Francu Vodopivcu za njegov 65. rojstni dan. Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The tramp elements Sb and Sn have a strong tendency to surface segregation on iron. By LEED and AES surface structures and concentrations of Sb and Sn segregated on single crystal vvere determined. The surface segregation is strongly dependent on orientation. therefore recrystallization of steel sheet is affected since the surface energies of different grains are reduced to different extent - this effect may be used to obtain advantageous textures of electrical steel sheet and deep dravving steels. Surface segregation of Sb and Sn retards surface reaction kinetics as vvas shovvn for the gas carburization of čase hardening steels. Surface segregation of tin in creep cavities of turbine steels vvas shovvn to accelerate the creep fracture. The grain boundary segregation of both elements in iron is minor, and furthermore Sb and Sn are displaced from grain boundaries by carbon so that most steels are not endangered by grain boundary embrittlement due to Sb and Sn, but some low alloy turbine steels are susceptibie to temper and long-term embrittlement. Key vvords: surface and grain boundary segregation Fe-Sb alloys, Fe-Sn alloys, Fe-Sb-C alloys, Fe-Sn-C alloys, intergranular fracture embrittlement Elementa v sledeh Sb in Sn močno segregirata na površini železa. Površinska struktura in koncentracija Sb in Sn v segregirani plasti sta bili določeni z metodami LEED in AES. Površinska segregacija je odvisna od kristalografske orientacije, rekristalizacija jeklenih pločevin je aktivirana, ker imajo posamezna kristalna zrna različno znižano površinsko energijo - pojav se lahko uporabi za pridobivanje prednostnih tekstur elektro pločevin in pločevin za globoki vlek. Površinska segregacija Sb in Sn zavira kinetiko površinske reakcije kar je prikazano pri procesu naogljičevanja jekel. Površinska segregacija kositra v vdolbinah pri lezenju jekel za turbine povzroča pospešenje lezenja do preloma. Segregacija obeh elementov po mejah kristalnih zrn je v železu minimalna zato ker Sb in Sn na mejah zrn izpodrine ogljik. Tako večina jekel ni ogroženih zaradi krhkosti kristalnih mej. ki bi jih povzročala Sb in Sn, le nekatera nizka ogljična turbinska jekla so občutljiva na popuščno krhkost. Ključne besede: površinska segregacija, segregacija po mejah zrn, Fe-Sb zlitine. Fe-Sn zlitine, Fe-Sb-C, Fe-Sn-C, interkristalna krhkost 1 Introduction 1.1 The role of tramp elements in steels The effects of the so-called tramp elements in steels, Ni, Cu, P, S, Pb, As, Sb, Sn etc. are generally deleterious, the greatest problems they cause are 'hot shortness' and 'temper embrittlement' of steels. The hot shortness, a lack of hot workability can have different reasons, one possible reason is the copper enrichment due to surface scaling1-2. Beneath the scale the more noble elements Cu, As, Sb, Sn are enriched and form a liquid phase which causes surface cracking by grain boundary penetration. Sb and Sn greatly reduce the solubility of Cu in austenite and hence lead to precipitation of a molten phase and its grain boundary penetration, under conditions of much less enrichment and down to lower temperatures. The enrichment of tramp elements below the oxide scale upon reheating or hot rolling of steels and could be detected by electron microprobe (EPMA). This enrichment also can have strong effects on the scale adherence and mor- 1 Prof.Dr.Sc. Hans Jiirgcn GRABKE Max-Planck.Inslilut fiir Eisenfurschung GmhH Postfach 140444. 40074 Dussedldorf. Germany phology as has been studied extensively by F. Vodopivec et al. in vvork started at the IRSID311: by the presence of the more noble elements Cu, Ni, Sb, Ag, S the scale adherence is enhanced vvhereas the elements Si, Al, P and B which are oxidized and form silicate, aluminate, phos-phate or borate layers cause formation of voids and cavities at the scale/metal interface. The other way of enrichment which leads to deleterious effects of tramp elements is equilibrium segregation, so the 'temper embrittlement' is caused by segregation of P, Sn or Sb in the temperature range 400 - 700°C to the steel grain boundaries, e.g. during sIow cooling after tempering, but also during application of steels in this temperature range. It vvas suspected since long that temper embrittlement is caused by grain boundary segregation, but this suspect could be confirmed only after the arrival and spreading of interfacial analysis by Auger-electron spectroscopy (AES) in the eighties. But the tramp elements do not have only deleterious effects, e.g. it is knovvn that Cu can enhance the resistance against atmospheric corrosion. Even positive effects of Sb and Sn vvere detected and studied at the IMT Ljubljana and the MPI fiir Eisenforschung Diisseldorf12"20, these tramp elements can improve the texture and magnetic proper- ties of nonoriented silicon steel sheets, caused by surface segregation and its effect on surface energies as dis-cussed in the following chapter. 1.2 Fundamentals of surface and grain boundary segregation In this revievv the equilibrium segregation of Sb and Sn will be described and only the effects will be dis-cussed which are caused by equilibrium surface and grain boundary segregation. Most elements which are dissolved in iron tend to enrich at elevated temperatures at surfaces, grain boundaries and interfaces21"25, and dis- tribution equilibria are established at sufficiently high temperature. A (dissolved) A (segregated) (D There are different driving forces for such equilib-rium segregation: 1. free bonds at the surface or interface can be satu-rated by interaction vvith the atoms A 2. the iron surface may be covered with a layer of atoms A which has a lower surface energy than the initial iron surface 3. the release of atoms A from the bulk solution leads to release of elastic energy, especially in the čase of in- Gibbs In a Langmuir - Mc lean fl - 1*Ka O. 4_L RT* R ln a [n a Figure 1: Schematic diagrams on the Gibbs isotherm (a-c) and the Langmuir-McLean isotherm (d-f) a) surface energy y vs activity a of the adsorbed or segregated element A, b) y vs In a and c) the latter plot for two orientations vvith different surface energies and different adsorption or segregation behaviour - upon increasing activity a and coverage 0 the surface firstly instable becomes stable, a reason for tertiary recrystallization or facetting, d) degree of coverage 0 vs activity of the absorbed or segregated element A, e) plot for the evaluation of studies at constant activity or concentration of the element A, f) isosteres for determination of the thermodynamic data at constant coverage Slika 1: Shematski diagrami Gibbsove izoterme (a-c) in Langmuir-McLeanove izoterme (d-f) a) površinska energija y v odvisnosti od aktivnosti a adsorbiranega ali segregiranega elementa A, b) y v odvisnosti od ln a in c) zadnji grafikon za dve orientaciji z različnimi površinskimi energijami in različno adsorbcijo oziroma segregacijo - po zvišanju aktivnosti a in pokritja ©je sprva nestabilna površina postala stabilna, vzrok za terciarno rekristalizacijo ali facetiranje, d) stopnja pokritja 0 v odvisnosti od adsorbiranega ali segregiranega elementa A, e) grafični prikaz za ovrednotenje študij pri konstantni aktivnosti ali koncentraciji elementa A, f) izostere za določitev termodinamičnih podatkov pri konstantnem številu atomov A terstitial atoms or substitutional atoms larger than the iron atoms. The latter effect is certainly true for Sb and Sn since both elements have large atoms, causing a strain in the iron lattice and increase of lattice parameter26. In fact, ali equilibrium segregation processes should lead to a de-crease in surface energy (or interfacial energy) according to Gibbs' law dy d ln a, = -RTT4 (2) where y is the surface energy, aA the thermodynamic ac-tivity of the segregating species A and Ta the surface concentration (mol/cm2), R gas constant, T temperature (K) (Figure la-c). The effect of adsorption or segregation on surface energy can be measured by the so-called zero creep method27 but onIy at very high temperatures. One example of the result for a measurement on Fe-Sn foils vvith different Sn concentrations at 1420°C28'29 is given in Figure 2a. Combining such a study with measuring 'grain boundary grooving', i.e. the dihedral angle of the thermally etched grain boundary grooves at the surface gave the ratio of the grain boundary to surface energies and thus the dependence of grain boundary en-ergy was derived as a function of the bulk tin content (Figure 2b). From these 'Gibbs isotherms' also the iso-therms for surface resp. grain boundary segregation could be derived27"29. Hovvever, these techniques were time consuming, difficult and tedious and since the arri-val and spreading of AES they are no more used. In manv cases, segregation can be described by a simple equation, the Langmuir-McLean isotherm (Figure ld-f), describing segregation to a limited number of sites which leads to a maximum coverage TAsat when ali sites are occupied, and with a free energy AGa which is independent of coverage. Then the degree of coverage o = FA/rA is given by 0A/(1-0A) = Xa exp (-AGa/RT) (3) (4) Since AGa = AHa - T ASa (5) this leads to the form of the Langmuir-McLean equation ln- ©A 1-0. AHa AS, RT R • + ln x. (6) vvhich is used to derive the enthalpy and entropy of segregation from measurements of ©a at a constant bulk concentration xa of the segregating species in dependence on temperature. Such measurements have been conducted, e.g. for the surface segregation of C, Si, N, P and S on iron and also for the grain boundary segregation of P, Sb and Sn (see Figure 6 and 8). The surface analyses were conducted by AES, observing the concentrations in situ on single or polycrystalline surfaces in dependence on temperature30"40. E.D. Hondros, M.P Seah 1970 0.2 04 tin (wt.%) mt1 icrl tin (wt.%) Figure 2: a) Surface energy and b) grain boundary energy of iron-tin alloys at 1420°C plotted as a function of the bulk tin content28,29 in (b) also the grain boundary segregation isotherm is given, which can be derived from the measurements Slika 2: a) površinska energija in b) energija kristalnih mej zlitine železo-kositer pri 1420°C kot funkcija vsebnosti kositra v osnovni zlitini28'29 (b) podana je tudi izoterma segregacije po mejah zrn, ki jo lahko izračunamo iz meritev The grain boundary analyses are also performed by AES, but after annealing the specimens for sufficient time at elevated temperature, then introducing them into the UHV system and fracturing in-situ by impact or ten-sile test35,36,39'40. The analysis of intergranular fracture facets yields the grain boundary concentration, assuming that the content of impurity A has been distributed equally to both sides upon fracture. The sites and structures attained in surface segregation can be elucidated using LEED (=low energy elec-tron diffraction). In most cases the elements A are en-riched on Fe(100) up to half a monolayer, corresponding to a c(2x2) structure, only for oxygen a complete mono-layer and p(lxl) structure is attained. At grain bounda-ries rather high coverages are possible, for P in ferrite coverages nearly up to one monolayer have been observed. The observation of the LEED structures on single crystal surfaces gives a good possibility for calibrating the AES measurements, also at grain boundaries, since the coverage for the saturated LEED structures in known. Further information on segregated species can be ob-tained using photoelectron spectroscopy (XPS), the pho-tolines obtained can indicate the ionization state of ions and the charge transfer between substrate and segregated atom4'"45. Generally, there is a transfer of negative charge (electrons) to the segregated atoms, which means that these (C, N, S, O, P etc.) are present as negatively charged atoms (anions) on the metal surface. This most probab!y is also the čase in the grain boundary segregation, and it is supposed that such charge transfer vveakens the cohesion of grain boundaries46'47 - leading to temper embrittlement of steels. In the čase that two elements are segregating simulta-neously to a surface or a grain boundary, there is gener-ally a competition for the sites available and the relative amount of both species in the surface depends on their free energy of segregation and concentrations in the bulk. Cases of competitive segregation have been studied on the iron surface for carbon and silicon38, and at grain boundaries: carbon-phosphorus36, carbon-sulfur48, nitro-gen-phosphorus37... The simple formalism for competitive segregation without further energetic interaction of the segregating species is given by 0A / (1 -0A-0B) = xA ■ exp (-AGa / RT) (7) 0B / (1-©A-0B) = xB ■ exp (-AGB / RT) (8) which could be applied in the cases mentioned above. In the literature on temper embrittlement there is a lot of fuss about 'cosegregation', the mutually enhanced segregation of two species where attractive energetic interaction is to be assumed. In some cases the enhanced segregation can be explained in a different way - in other cases which are important here (Ni-Sn, Ni-Sb) formation of two- or three-dimensional phases at the grain boundaries may be suspected (see below, chapters 2.2 and 3.4). 1.3 Systems Fe-Sn and Fe-Sb The solid solutions of Sn in a-Fe vvere determined by lattice parameter measurements49,50. Accordingly, the solubility ranges from a maximum at 9,2 at% (17,7 wt%) at 900°C to 3,2 at% (6,56 wt%) at 600°C. The solubility limit in y-Fe has been determined24'26 the y-loop extends to 0,92 at% (1,93 wt%). Own investigations on Fe-0,054 wt% Sn and Fe-0,080 wt% Sn [unpublished], however, showed precipitation of Sn-rich particles on the grain boundaries after long-term annealing at 550°C; accord-ingly, there are uncertainties on the solubility at temperatures <600°C. The solubility of Sb in a-Fe has been determined by several authors, the results are in substantial agree-ment49. The solubility at 900°C is 4,19 at% Sb (8,71 wt%) decreasing at 600°C to 2,58 at% (5,46 wt%). The Y-loop extends till 1,1 at% Sb (2,36 wt%). Experimental and theoretical studies have been con-ducted on the effects of other alloying elements on the antimony solubility, they vvere found to be the largest for M = Ti, Mn and Ni and small for M = Cr, Co. The pres-ence of Ni e.g. reduces the solubility strongly, the phase precipitating is a hexagonal NiAs type: Fe96Sb2Ni2. A cubic CaFei type Fe97Sb2Ti is formed with Ti, which reduces the solubility at 900°C to 1,91 at%52. Strong interaction of Ni and Sb is also observed in surface segregation53. 2 Interfacial segregation of Sn and Sb on and in iron and steels 2.1 Surface segregation of Sn and Sb on iron The surface segregation of tin on Fe-Sn single crys-tals has been studied in the temperature range 450°C to 650°C, mainly on crystals vvith relatively high Sn concentrations so that always saturation coverages vvere observed, no dependence of coverage on temperature, so that the segregation enthalpy was not obtained54 55. Each of the low index orientations exhibits a characteristic behaviour of the segregating Sn, the coverages attained are governed by segregation kinetics (Figure 3a). After heating the specimen for a short time a c(2x2) structure is observed, corresponding to half a monolayer coverage. But the segregation continues which leads to an order-disorder transition and a coverage somevvhat higher than a monolayer (corresponding to 1,4 • 1015 atoms Sn/cm2). The transition is accompanied by a shift of the photoli-nes observed by XPS to values closely corresponding to the values characteristic for pure elemental tin, Figure 3b. Most probably the transition can be explained by formation of a two-dimensional nearly close packed layer of tin on Fe(100) vvith a high surface mobility. This segregation behaviour is different from the segregation in most systems Fe-A (A = C, N, S, P, Sb...) vvhich always leads to a saturation at a surface coverage of 0,5. The driving force for the segregation of tin to higher coverages is probably the strong decrease of surface energy by the presence of a layer of tin. This layer at high coverage has properties similar to a layer of pure molten tin on iron, as indicated by the results of the XPS measurements. The segregation behaviour on Fe-Sn(lll) is similar, there is an inflection point in the kinetics vvhen the p(lxl) structure vvith one monolayer coverage is reached, vvhich corresponds to 7 ■ 1014 atoms Sn/cm2 on Fe(lll). After this surface structure is reached, further Sn segregation occurs, an order-disorder transition is observed and a Sn monolayer is attained. The segregation behaviour is different on Fe-Sn(llO), here no intermedi-ate adsorption structures vvere observed, but only structures vvith high Sn content, firstly a hexagonal structure corresponding to one monolayer of grey tin. Upon fur- Figure 3: a) Kinetics of the tin surface segregation on Fe-4 wt% Sn(100) during heating to 650°C54, at the inflection point indicated the structuraJ phase transition from the ordered monolayer c(2x2)Sn to the disordered multilayer occurs; b) photolines observed during increasing surface concentration demonstrating the shift caused by the transition Slika 3: a) Kinetika površinske segregacije na zlitini Fe-4 ut.% Sn(100) med žarjenjem do 650°C54, prevoj označuje strukturni fazni prehod iz urejene monoplasti c(2x2)Sn v neurejeno večplastnost; b) fotolinije med naraščanjem površinske koncentracije prikazujejo kemijski premik, nastal zaradi prehoda ther segregation a structure is formed which corresponds to a layer of the intermetallic compound FeSn of one unit celi thickness, Figure 4a. The segregation behaviour of Sb on Fe-4 wt% Sb55'56 is similar on the orientations (100) and (111) to the behaviour of tin, on both orientations on ordered adsorp-tion structure is formed, c(2x2) on (100), see Figure 5, and p(lxl) on (111) but upon continued segregation no elevated Sb surface concentration were observed, in con-trast to Sn. On Fe(l 10) the presence of Sb caused facet-ing, the LEED patterns indicated formation of (111) and (111) planeš, Figure 4b. Accordingly, the segregation enthalpy of Sb to Fe(lll) must be very exothermic (negative), due to a strong decrease of the surface energy of Fe(l 11) which compensates the increase of total surface area by the faceting. Figure 4: Phenomena on the Fe(l 10) face caused by segregation of Sn or Sb; a) Supposed structure of the surface compound 'FeSn' formed by epitaxial stabilization on Fe-Sn(l 10) as the final saturation structure ; b) Faceting on Fe-Sb(llO) under formation of (111) faces due to Sb segregation56 Slika 4: Pojav na Fe(110) ploskvi, ki gaje povzročila segregacija Sn ali Sb; a) predpostavljena struktura zlitine na površini 'FeSn', ki je nastala z epitaksialno stabilizacijo na Fe-Sn(llO) kot končna nasičena struktura54; b) facetiranje na površini monokristala Fe-Sb(llO), zaradi segregacije Sb se tvorita (111) in (111) ploskvi Three possibilities are demonstrated in the systems Fe-Sn and Fe-Sb for the behaviour upon segregation, (i) formation of adsorption structures such as c(2x2) or p(lxl), (ii) formation of surface phases such as two-di-mensional grey tin and two-dimensional FeSn, or (iii) formation of facets to attain surface energies. 2.2 Grain boundary segregation of Sn and Sb A fundamental study on grain boundary segregation in Fe-Sn alloys has been conducted after annealing in the temperature range 500-750°C for up to 5000 h39. The results of the grain boundary analyses show a wide scatter Ek (eVl Figure 5: Surface segregation of Sb on Fe-Sb (100)56; a) Auger spectrum after segregation at 640°C, corresponding to surface segregation; b) model for the Fe-Sb (100) c(2x2) structure derived from LEED study of the saturated surface Slika S: Površinska segregacija Sb na monokristalu Fe-Sb orientacije (100)56; a) AES spekter posnet po segregaciji Sb pri 640°C; b) model za Fe-Sb (100) c(2x2) strukturo dobljen z metodo LEED na nasičeni površini (Figure 6a) which may be caused by the strong dependence of tin segregation on grain boundary orientation. Ali data have been obtained for Sn concentrations within the a-solid solution range, no precipitates of intermetal-lic compounds should have formed. The tin concentrations are always below a monolayer, in contrast to the surface segregation behaviour. In spite of the large scat-ter the data were evaluated according to the Langmuir-McLean equation (Figure 6b), yielding the values for segregation enthalpy and entropy AH = -22,5 kJ/mol AS = 26 J/mol K for 550°C results in good agreement with previous results of E. D. Hondros and M. P. Seah28 29. The enthalpy value is relatively low (P: AH = -34,3 kJ/mol35'36), this indicates the rather low tendency for grain boundary segregation of Sn! Furthermore, due to o 0.20 %Sn ^ 0.08 '/.Sn o 0.054% Sn x 0022% Sn •----■--- 0-g- * . i temperature (°CI Figure 6: a) Grain boundary concentrations of tin in Fe-Sn alloys after annealing at elevated temperatures, measured by AES on intergranular fracture faces39; b) evaluation of the measurements in (a) applying the Langmuir-McLean equation (6) Slika 6: a) Koncentracija kositra v segregirani plasti na mejah zrn po žarjenju pri povišanih temperaturah, merjeno z metodo AES na interkristalnih prelomnih ploskvah39; b) ovrednotenje meritev (a) z uporabo Langmuir-McLean enačbe (6) its low segregation enthalpy Sn is kept from the grain boundaries effectively by the presence of carbon such as in plain carbon steels (Figure 7). As described in the in-troduction, an equilibrium of site competition between Sn and C occurs according to C(dissolved) + Sn(segregated) = C(segregated) + Sn(dissolved) In the presence of some ppm dissolved carbon, the tin is effectively removed from the grain boundaries. Hovvever, in low alloy steels the concentration of dissolved C is reduced due to the formation of less soluble carbides vvith Cr and Mn. Tin segregation is possible if not Sn is displaced from the grain boundaries by segre-gated phosphorus. For rotor steels, CrMoV steels it is even dangerous to have too low phosphorus contents, since in application at high temperature if Sn segregation prevails, this easily leads to formation of creep cavities, due to the strong tendency for surface segregation of tin. The surface segregation of Sn decreases the surface en-ergy of pores and cavities, stabilizes such defects and ac-celerates their growth (see chapter 3.2). In must be kept in mind that the data given above are average values and have an integral character, since the grain boundary segregation of tin in iron is strongly de- 0.10 0.05 cn - Q \ \ \ \ \ A ir-\ * / v / / -— C S / / __ \ ' ' / \ 1 > \ ' ! \// / /'V / !// o / 550 °C o * -0 Sn Fe-Fe- 0.08 V.Sn 0.054 V.Sn / Fe- 0.022 V.Sn ' » • 20 40 60 [Cl (ppm) 550 600 650 700 750 annealing temperature (°C) Figure 7: Grain boundary segregation of Sn and C in Fe-Sn-C alloys in dependence on the bulk carbon concentration after equilibration at 550°C, demonstrating the displacing effect of carbon on segregated Sn39 Slika 7: Segregacija Sn in C po mejah zrn v Fe-Sn-C zlitinah v odvisnossti od koncentracije ogljika v osnovnem materialu pri ravnotežju pri 550°C pendent on the misorientation, increasing with the tilt angle of misorientation betvveen the grains57. Sn causes grain boundary hardening, excess hardness extending to many mtcrons on either side of the grain boundary, also increasing with the misorientation. This is a well docu-mented effect but not well understood57-58. For temper embrittled Ni-Cr steels there are strong indications that Sn is present at the grain boundaries cou-pled with Ni in an bidimensional phase corresponding to an intermetallic compound such as Ni3Sm, this has been concluded from Mossbauer spectroscopy and TEM work59"61. The grain boundary segregation of Sb was investi-gated for Fe-Sb and Fe-Sb-C alloys after equilibration at temperatures betvveen 550°C for sufficient time62,63. The analysis of intergranular fracture faces by AES calibrated on the base of the surface segregation studies shovvs relativen low interfacial concentrations, see Figure 8, and a wide scatter of results. The plot of the data according to the Langmuir-McLean equation leads to the values for segregation enthalpy and entropy: AH = -19 kJ/mol AS = 28 J/mol K Thus, the segregation enthalpy is even lower than for Sn, which emphasizes the Iow tendency for grain bound-ary segregation of Sb. However, even small grain bound-ary concentrations of Sb cause marked grain boundary embrittlement and prevailing intergranular fracture. Also the segregant Sb is effectively displaced from grain .o LO X XI LO CD X) LO CD b) ■ 0.012% Sb • 0.049% Sb * 0.094% Sb 0,95 1,00 1,05 1,10 103 / T I K 1,15 -11 1,20 1,25 Figure 8: a) Grain boundary concentrations of Sb in Fe-Sb alloys, plotted vs equiIibration temperature62; b) plot of the data in (a) according to the Langmuir-McLean aquation (6) Slika 8: a) Segregacija Sb po mejah zrn v Fe-Sb zlitini v odvisnosti od ravnotežne temperature62; b) prikaz podatkov v (a), ki ustrezajo Langmuir-McLean enačbi (6) boundaries by carbon, small concentrations of dissolved carbon < 60 wtppm can shift the displacement equilib-rium to low Sb segregation and also lead to a marked reduction of intergranular fracture, see Figure 964. Carbon not only removes Sb from the grain boundaries, but also enhances the grain boundary cohesion and enforces transgranular fracture. The effect of carbon also was demonstrated by notch-impact tests on Fe-Sb-C alloys, see Figure 10. As in the čase of Sn, for unalloyed carbon steels the danger of embrittlement by Sb is minor, there will be always enough dissolved and segregated carbon to avoid Sb grain boundary segregation. Only for alloyed steels, in which the carbon is tied up by carbide forming elements, Cr, Mn, etc., embrittlement is possible during heat treatment or use of steels in an elevated temperature range. Several authors have claimed an effect of nickel, en-hancing the grain boundary segregation of Sb, however, this effect could not be reproduced in recent studies on ai k_ 3 u O 10 carbon 20 30 concentration 40 50 ! wt.ppm 80 rs 70 F o 60 —i 50 40 m aj 30 C a> 20 i . o n 10 LL E 0 Fe - 0.05 % 6 ppm C -100 -50 50 100 temperature (°C) Figure 10: Results of notch impact tests on an Fe-Sb al!oy with different carbon concentrations . The ductile-brittle transition temperature is shifted to lower temperatures by carbon, due to the removal of Sb from the grain boundaries and increase of grain boundary cohesion by segregated carbon Slika 10: Rezultati udarnih preizkusov na zlitini Fe-Sb z različnimi vsebnostmi ogljika6-1. Temperatura prehoda duktilno-krhko je premaknjena k nižjim temperaturam zaradi ogljika, le-ta izpodrine Sb z mej zrn in poviša kohezijo a a. 3 C a L. cn k_ a« C Figure 9: a) Grain boundary segregation of Sb and C in Fe-Sb-C alloys after equilibration at different temperatures, plotted in dependence on the bulk concentration, demonstrating the displacing effect of carbon on segregated Sb62,63; b) intergranular part of fracture in dependence on bulk carbon concentration Slika 9: a) Ravnotežna koncentracija Sb v segregirani plasti na mejah zrn v zlitini Fe-Sb-C pri različnih temperaturah, prikazana v odvisnosti od koncentracije C v osnovnem materialu, prikazuje pojav ko ogljik izrine Sb v segregirani plasti62,63; b) interkristalna ploskev preloma v odvisnosti od koncentracije ogljika v osnovnem materialu Fe-Ni-Sb alloys64. In earlier studies65-66 of Fe-Sb and Fe-Ni-Sb alloys at 560°C an increase of Sb segregation was observed with hte Ni-content and Ni also segregates to the grain boundaries, its segregation being only slightly affected by the presence of Sb. For low alloy Ni-Cr steels the authors65'66 conclude that the Sb-segregation is a complex function of the total alloy composition. When Mn is present in these steels it causes precipitation of an antimonide and greatly reduces Sb-segregation. A de-tailed investigation of a 3,5 Ni-ICr-steel after embrittle-ment at 480°C demonstrates a dependence on the microstructure67. Intergranular embrittlement in a quenched and tempered martensitic microstructure was associated with the segregation of phosphorus, which is possible since the carbon activity is reduced by precipitation of chromium rich carbides at the grain boundaries. the embrittlement in the bainitic microstructure was associated with the segregation of antimony, since the carbon activ-ity is relatively high due to the formation of cementite type carbides. Prolonged embrittlement of the bainite produced a low energy fracture. Increased nickel and an-timony concentrations at the grain boundaries vvere associated with the formation of a fine grain boundary pre-cipitate. The increased carbon activity continued to prevent appreciable P segregation but could not inhibit the 'cosegregation' of Ni and Sb67. 2.3 Segregation of Sb and Sn at internat interfaces Sb can be trapped by TiC precipitates in Fe. A dense dispersion of TiC, produced by ion implantation and annealing at 600-700°C, ties up Sb effectively68. Continued annealing leads to slow release of Sb into the matrix in a diffusion and trapping process. The Sb is present at the interface TiC/ferrite, and not in the TiC, the binding en-thalpy is -35,6 kJ/mol68"70. This interfacial segregation may provide a means for keeping Sb from grain boundaries in ferritic steels to suppress embrittlement. Similar trapping has been observed at TaC and Cu precipitates in Fe at 600°C71. Trapping or segregation of Sn at MnS particles has been observed in Fe-3% Si doped vvith tin. The Sn was clearly enriched compared to the grain boundaries, this segregation retards the growth rate of the MnS particles so that in Sn doped alloy they are much smaller than in Sn-free Fe-3% Si72. The size of the precipitates affects the primary and secondary recrystallization, thus influ-encing the magnetic properties of Si steels, see chapter 3.1. In the eutectoid transformation of austenite to čast iron, minor additions of Sb (0,08 wt%) or Sn (0,12 wt%) 0 10 20 30 40 50 60 carbon concentration (wt.ppm) » 700°C .--•-- 650°C —600°C were found to inhibit the y —> a + graphite and the Fe.iC —> a + graphite reaction paths, but did not significantly affect the metastable y —> a + FejC reaction73. Scanning Auger microprobe analysis indicated that Sn and Sb ad-sorb at the graphite/metal interface. The segregated layer acts as a barrier for the access of carbon to the graphite nodules. With the graphite disabled as a sink for carbon, the metal transforms as a nongraphite steel. 3 Effects of interfacial segregation of Sn and Sb on steel properties 3.1 Effects of surface segregation on the texture of electrical sheet A (100) [001] texture of Fe-Si can be achieved vvith the aid of adsorption or segregation of different species: O, S, Sb, Sn etc. The (100) [001] texture cannot compete lossvvise vvith the (110) [001] texture if unidirectional magnetization is important. In applications vvhere the magnetization must occur in ali directions in the plane of the sheet such as in motors or generators the (100) [001] texture is favourable since the plane of the sheet does not contain the hard (111) direction of magnetization, but even in transformers lovver losses can be obtained by using some (100) [001] texture. After the primary recrystal-lization, the grovvth of grains is governed by the surface energy, preferential grovvth of grains vvith a low surface energy occurs in the secondary recrystallization. In ab-sence of oxygen or other adsorbing or segregating species yno is the lovvest surface energy and (110) [001] grains grovv. When sufficient oxygen or sulfur is present yno and (100) grains become stable in the surface74"76, see also Figure 1. Presence of oxygen and sulfur is not vvell possible in the production process of non-oriented electrical sheet. The annealing for secondary recrystallization is done in dry hydrogen at about 900°C. Presence of sulfur vvould cause precipitation of MnS particles in the steels vvhich may hinder the reorientation of the magnetic domains. Thus, other elements such as Sn and Sb vvere success-fully used as alloying additions to improve the texture and magnetic properties of non-oriented steel sheet77-78. The alloying additions may not be too high to obtain the vvanted (100) [001] texture, for too high activities and surface coverages the surface energies of nearly ali ori-entations are decreased so strongly that no preferential grovvth of (100) is attained. Sb has proved to have an-other advantageous effect, it suppresses widely the inter-nal oxidation of the alloying elements Si, Al and Mn vvhich is possible during the decarburization treatment and causes increasing permeability deterioration vvith in-creasing subscale depth79. Also in the production of high induction and high permeability grain oriented Fe-Si, the presence of Sb and Sn can have positive effects, yielding a more precise (110) [001] secondary recrystallization texture than in conventional Fe-Si. In earlier vvork it vvas assumed that Sb and Sn are effective on the primary re-crystalIization, retarding primary grain grovvth in coop-eration vvith BN and S, less S being necessary than vvith-out Sb and Sn. But in recent studies it vvas found that grain boundary segregation of Sb and Sn is negligibly lovv in the silicon steel sheet after the usual thermal treatment. Obviously, the effect of Sb and Sn is caused by the surface segregation during recrystallization annealing. The surface segregation decreases the surface energy of grains vvith (100) orientation in the plane of the steel sheet and the grains vvith lovv surface energy grovv on ac-count of grains vvith other space orientation in the sheet plane. The role of the surface segregation has been con-firmed by extended studies on silicon steel doped vvith Sb and Sn12"20. Only a controlled surface segregation promotes the vvanted selective grain grovvth. For too high Sb and Sn concentrations the surface energy of ali orien-tations are strongly decreased and no preferential grovvth of (100) is obtained. For steels vvith a high Sb content rather the unvvanted grovvth of (111) is to be expected, since the surface concentration on that plane is highest56. Furthermore, it has been stated that Sb and Sn retard the decarburization79, vvhich is also a very important process in the production of electrical steel sheet - so this again vvould be a negative effect of Sb and Sn surface segregation. 3.2 Effect of surface segregation in creep of steels Due to their strong tendency to surface segregation Sn and Sb can very negatively affect the creep behaviour of heat resistant CrMo- and CrMoV- steels used for turbine rotors and blades80. The failure of such steels occurs by formation of creep cavities at the grain boundaries and in the steel matrix and the coalescence of the cavities to cracks. The nucleation of the cavities mostly starts at inclusions, such as sulfides (MnS) and oxides81. But the nucleation is favoured and accelerated by the presence of Sn or Sb vvhich vvill immediately segregate to the free metal surface of a pore forming at an inclusion or at a grain bondary. The segregation decreases the surface en-ergy, the pores are stabilized and can grovv to cavities un-der further surface segregation. This effect of Sn has been observed for a CrMoV- steel39-40 measuring creep curves for melts doped vvith different Sn-concentrations. The higher the Sn content of the steel the earlier they failed by rupture in the creep test, Figure 11. 3.3 Effects on the carburization of čase hardening steels Surface segregation on Sn and Sb can effectively retard the carburization of čase hardening steels82. The carburization is generally conducted at about 930°C in CO-H2-H2O atmospheres. In the beginning its rate is controlled mainly by the surface reaction sequence CO(g) = C (dissolved) + O (adsorbed) O (adsorbed) + H2(g) = H,0(g) Figure 11: Creep curves for 1% CrMoNiV- steel at 300 MPa and 500°C39, effect of different Sn-contents - vvith increasing Sn-content the rupture tirne is markedly decreased Slika 11: Krivulje lezenja za jeklo 1% CrMoNiV pri 300 MPa in 550°C39, vpliv različnih vsebnosti Sn - z naraščajočo vsebnostjo Snje prelomni čas opazno znižan and later on a coupled surface reaction and diffusion control determines the rate of carburization. The surface reaction rate can be described by r - P([C]e„ - [C],) where [C]eq and [C]s are the equilibrium and the actual surface concentration of carbon and (3 is the carbon transfer coefficient, which contains dependencies on partial pressures and temperature83. Extended thermo-gravimetric studies of carburization on steels doped vvith Sn, Sb, Cu, P or Pb demonstrated a strong effect of Sb on the coefficient (3 (see Figure 12a), vvhereas the effect of the other elements is much less. This retarda-tion of carbon transfer is caused by the blocking of surface sites for reaction, the adsorption and dissociation of CO, by segregated Sb. The surface segregation of Sb and Sn on the čase hardening steels vvas demonstrated by AES studies, after exposure in the carburization atmosphere at 900°C, see Figure 12b. Segregation in the UHV chamber leads to displacement of Sb and Sn by sulfur, hovvever, in the carburization atmosphere the sul-fur vvould be removed by the reaction S (absorbed) + H2(g) = H2S(g). The presence of too high levels of Sb in čase hardening steels vvould lead to too lovv carbon con-tents after the usual carburization period and insufficient hardening of the vvorkpieces, see Figure 12c. Thus, a specification for Sb-content < 25 wt ppm vvas recom-mended for čase hardening steels, vvhereas concentration of the other tramp elements may be in the usual range84. 3.4 Temper embrittlement Reversible temper embrittlement occurs upon slowly cooling of steels through the temperature range 550 to 350°C after annealing (tempering) at higher temperatures or during application of steels in this range. Temper em- mass% Sb orSnICu mass%/10) eV distance from the surface (mm) Figure 12: Effects of Sb, Sn and Cu on the gas carburization of a case-hardening steel at 930°C82,84, a) carbon transfer coefficient p in dependence on bulk concentrations of Sb, Sn or Cu, b) Auger spectrum of the Sb-doped steel after heating in hydrogen to 930°C, c) carbon concentration profiles after gas carburization of samples in an industrial furnace Slika 12: Vpliv Sb, Sn in Cu na plinsko naogljičenje jekla za cementacijo pri 930°C82-84, a) (3 prenosni koeficient ogljika v odvisnosti od koncentracije Sb, Sn ali Cu v osnovnem materialu, b) AES spekter jekla legiranega z Sb po žarjenju v vodiku pri 930°C, c) koncentracijski profil ogljika po plinskem naogljičevanju vzorcev v industrijski peči brittlement is caused by grain boundary segregation of P, Sn, Sb and As61"72 but severe embrittlement is observed only if the alloying elements Ni, Cr and Mn are present, such as in low alloy turbine steels. In earlier years this fact was explained by 'cosegregation', e.g. of Cr and P, however especially for this čase it could be clearly shown that Cr alone has no enhancing effect on P-segre-gation35,36. In fact, Cr and Mn decrease the carbon solu-bility in steels, and the effect of carbon on P-segregation, i.e. removal of P from the grain boundaries by displace-ment by C, is reduced in the presence of Cr and Mn, thereby allovving more P-segregation. 'Cosegregation' was also suspected for Ni and Sb, and Ni and Sn, but most probably the strong effect of these combinations on embrittlement are due to interfacial formation of inter-metallic compounds of these elements. Steels without Ni do not show embrittlement by Sn or Sb85"96. Temper embrittlement is a particular problem for low alloy steels, e.g. Ni-Cr-Mo-V rotor steels and Cr-Mo pressure vessels. Temper embrittlement does not occur in plain carbon steels vvith less than 0,5% Mn. At high Mn concentrations, hovvever, P-segregation is possible in plain carbon steels and also a 'cosegregation' of Mn and Sb is supposed to occur. Hovvever, the effect of Mn can easily be explained by the reduction of carbon activity caused by formation of Mn-rich carbides97. Thereby, carbon segregation is reduced vvhich allovvs grain bound-ary segregation of P, Sn and Sb. 3.5 Hydrogen induced cracking The threshold stress intensity for cracking of a Ni-Cr-Mo steel is strongly reduced by grain boundary segregation of Sb, Sn and p98"100. In the presence of hydrogen this threshold stress intensity is lovvered further, but it is the impurity effect vvhich is dominant, the hydrogen merely accentuates the tendency for brittleness already present. If must be emphasized again that for embrittlement of steels by Sb the presence of Ni and Cr is necessary. Sb causes intergranular fracture in the constant strain rate test, it is five times more effective in inducing intergranular fracture at cathodic potentionals than S, the results are consistent vvith H-permeation studies in Fe as affected by Sb and S101-102. 3.6 Possible effects of grain boundary segregation in in- terstitial free steels One may expect that grain boundary segregation of Sn and Sb is possible in interstitial free steels (i.f. steels). Such steels have very lovv concentrations of C and N in order to attain good deep dravving properties, and thus the tramp elements are not kept away from the grain boundaries by segregated carbon (see Figure 10). The effects of Sn and Sb in deep dravving steels vvere not studied as yet, but a behaviour similar to p103'104 may be expected. Brittle behaviour vvas found for steels vvith very lovv C content (< 300 ppm), caused by P at grain boundaries. Some i.f. steels are alloyed vvith Ti to tie up the interstitial elements, but Ti is also effective in scav-enging the phosphorus forming a very stable phosphide. Also TiC as a precipitate is able to trap phosphorus and to keep it from the grain boundaries to a certain extent, such trapping effect has also been reported for Sb at the TiC/ferrite interface68"70. Anyway, similar as P can be deleterious for the properties of certain deep dravving steels vvith lovv C and no TiC or excess Ti, also Sn and Sb may adversely affect the ductility of such steels. Es-pecially if steels vvith high Sn and/or Sb contents are slowly cooled after batch annealing or coiling, they may segregate to grain boundaries and cause embrittlement. On the other hand, also positive effects may occur on the texture, as in the čase of electrical steel sheet. Effects of Sb and Sn on the texture of deep-dravving steels are cur-rently investigated105. 4 Conclusions Generally, tramp elements such as Sn and Sb can have effects on steel properties only if they enrich at in-terfaces, the enrichment by equilibrium segregation leads to coverages in the range of a monolayer depending on bulk concentration and decreases vvith temperature. Util-izing Auger-electron spectroscopy the thermodynamics of segregation to surfaces and grain boundaries can be elucidated. The solubilities of Sn and Sb in the ferritic matrix are relatively high, the solubility is strongly decreased in the presence of some elements such as Ni vvhich form inter-metallic compounds vvith Sn and Sb. The tendency for surface segregation of Sn and Sb is very high, as yet no thermodynamic data have been determined since always saturation vvas observed. The segregation coverages and structures are very different for different crystallographic orientations, therefore the decrease of surface energy vvill be strongly dependent on orientation and marked effects of Sn and Sb on the sta-bility of different crystallographic planeš are to be ex-pected. Sn and Sb segregate to grain boundaries in ferrite, the extent of segregation strongly depends of the misfit of the grains. The tendency for grain boundary segregation of Sn and Sb is relatively lovv, as indicated by the results on equilibrium segregation in binary alloys in the temperature range 500 to 750°C. This can also be seen from the segregation enthalpies: -22,5 kJ/mol Sn and -19 kJ/mol Sb. Sn and Sb also segregate to interfaces, for Sb at inter-faces ferrite/TiC and for Sn at the interface ferrite/MnS. In the annealing of steel sheet the surface segregation of Sn and Sb affects the stability of certain orientations. For intermediate concentrations the (100) orientation ap-pears to be stabilized, for higher contents the (111) orientation becomes stable. These effects are of importance in the production of electrical Fe-Si steel sheet and may also be useful in the production of deep-drawing steels. The strong tendency for surface segregation of Sn (and Sb) plays a role in the creep of heat resistant steels, since formation and growth of creep cavities is enhanced by surface segregation decreasing the surface energy of the cavities. This was demonstrated for Sn doped CrMoV-steels. Surface segregation of Sn and Sb affects the carburi-zation of čase hardening steels. Especially Sb strongly retards the carbon transfer and may cause insufficient carburization. Grain boundary segregation of Sn or Sb causes em-brittlement. Temper embrittlement of low alloy steels only occurs in the presence of Ni and Cr. Obviously, re-duction of carbon activity by Cr and formation of inter-metallics NixSny resp. NixSby is necessary for temper embrittlement. Hydrogen induced cracking can be favoured by Sn and Sb grain boundary segregation. However, as for temper embrittlement the presence of Ni and Cr appears to be a precondition of such effect of Sn and Sb. In interstitial free steels one may expect strong ef-fects of Sn and Sb since these elements are not kept away from the grain boundaries by carbon. Addition of Ti may scavenge Sn and Sb, either by direct interaction or by trapping effect of TiC. 5 References I D. A. Melford: Phil. Trans. Roy. Soc. London. A295, 1980, 89 2M. Torkar, F. Vodopivec: Železarski zbornik, 15, 1981, 61 3 F. K. Peters, H. J. Engell: Archiv Eisenhiittenwes„ 30, 1959, 275 4 F. Vodopivec, A. Kohn, J. Benard: Comptes Rendus A. S., 255, 1962, July, 296-298 5 F. Vodopivec: Metaus-Corrosion-Industries, 452, 1963, April, 159-170 6F. Vodopivec, A. Kohn: Comptes Rendus A. S., 253, 1963, July, 448- 450 7 F. Vodopivec. A. Kohn, J. Philibert, J. Manenc: Mem. Scient. 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Cesta žeiezarjev 8, 4270 Jesenice te), centrala: +386 64 861-441 tel. dkektor: +386 64 861-443 tel. komerciala: +386 64 861 -474 Fax: +386 64 861-379 Teta: 37219 ZEUSN SI issf8 m*?" STEELS JjgjglS*'" fndtemp«"« OUR PRODUCTION PROGRAM INCLUDES: * general structural steels * finegrained and HSLA structural steels * carbon and alloyed steels - for quenchig and tempering - čase herdening * silicon steels for electrical sheets * stainless steels * hot rolled plates, wide and slit strips and bars * cold rolled sheets, wide and slit strips * cold rolled sections * metal door posts * blanks WE ALSO OFFER: * hot and cold rolling * blankig * torch cutting by drawing * straightening * heat treating of plates, strips and sheets Applications of Surface Analytical Techniques in Corrosion Research (Mainly High Temperature Corrosion) Uporaba površinskih analiznih tehnik v raziskavah korozije H.Viefhaus1, Max-Planck-lnstitut, Dusseldorf, Germany Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The application of materials under various ertvironmental conditions strongly depends on the corrosion properties of those materials. This is in particularly true for high temperature materials to be used in povver plants, petrochemical, chemical and automobil industry, where during application high temperatures and aggressive environments may cause great problems. At lovv temperatures the corrosion of metals is often inhibited by a passive layer on the metal surface. To understand the phenomenon of passivity the formation and nature of this surface film, quite often only a few nm's thick, has to be characterized. Corrosion protecting layers on high temperature materials, either grown during application or precovered before application, have a much larger thickness. In order to study the growth mechanisms, nature and properties of corrosion protecting layers thin films have to be characerized spreading over a quite large range of thicknesses, from a few nm's for the passive layers up to several nm's for the protecting layers on high temperature materials. Different methods for thin film analysis using surface analytical methods will be presented and illustrated by examples from different areas of corrosion research. Key vvords: corrosion, high temperature corrosion, surface analytical techniques Uporaba materialov v različnih okoljih zavisi od njihovih korozijskih lastnosti. To je posebno pomembno pri uporabi materialov pri visokih temperaturah in agresivnih medijih npr. v elektrarnah, petrokemijski, kemijski in avtomobilski industriji, kjer lahko pride do hudih industrijskih havarij. Pri nizkih temperaturah se na površini kovin tvori tanka pasivna plast, ki zavira korozijo. Razumevanje pojava nastanka in narave tanke pasivne plasti, ponavadi debele le nekaj nanometrov je mogoče samo s karakterizacijo teh plasti. Protikorozijske zaščitne plasti materialov, ki se uporabljajo pri visokih temperaturah in nastajajo med samo uporabo ali pa so bile predhodno nanesene, so debelejše. Študij mehanizma rasti, narave in lastnosti protikorozijskih prevlek je mogoč z raziskavami protikorozijskih prevlek. Te so različnih debelin, od nekaj nanometrov debelih pasivnih tankih plasti do nekaj mikronov debelih prevlek za zaščito na visokih temperaturah. Za analizo protikorozijskih plasti se uporabljajo različne metode površinske analize, ki so prikazane v članku, kakor tudi primeri z različnih področij korozije. Ključne besede: korozija, visokotemperaturna korozija, metode površinske analitike 1 Introduction A metal is normally described as beeing passive, if for the existing surrounding atmosphere a high corrosion rate would be expected, instead of the very lovv corrosion rate to be observed. This passivity is caused by very thin dense oxide (and/or hydroxide) layers vvhich are formed on the metal by the corrosion process. In order to get a better understanding on the effect of those passive layers they have to be analysed vvith respect to composition and thickness by very surface sensitive methods. High temperature oxidation and corrosion can cause great problems in povver plants, petrochemical and chemical indus-try. A very important precondition for the practical application of a metalic material at high temperature is its oxidation or high temperature corrosion resistance. This precondition may be fulfilled if on the surface of the material protecting oxide layers are formed. These oxide layers can grovv under use of the material by reaction of the material elements vvith the surrounding oxy-gen atmosphere or by a specific preoxidation at suitable temperatures in oxygen containing atmospheres. The protecting effect of those oxide layers relys on their property to act as a diffusion barrier betvveen the metallic ' Dr. Sc. H. VIEFHAUS Max-Planck-Inslilul fiir Eisenforschung GmbH 40074 Dusseldorf. Postfach 140 444 Germany and the corrosive atmosphere surrounding it. Oxyde lay-ers therefore have a key function for the application of materials in high temperature technology and there is a great need for doing research and testing the materials for such applications. 2 Methods In order to studv the grovvth mechanisms, nature and properties of corrosion protecting layers thin films have to be analysed spreading over a large range of thicknesses, from a fevv nm's for the passive layers up to several tenth of m(i's for the protecting Iayers on high temperature materials. To analyse thin films vvith respect to layer composition and thickness different depth profiling methods can be applied depending on the thickness of the layer under study and on various sample preparation methods. Fol-lovving the mainly applied surface analytical methods to be used for depth profiling of homogeneous and inhomogeneous surface layers are listed. Depthprofiling homogeneous layers inhomogeneous layers AES AES XPS S IMS SIMS SNMS GDOES AES is Auger electron spectroscopy, XPS is X-ray excited photoelectron spectroscopy, SIMS is secondary ion mass spectroscopy, SNMS is secondary neutrals mass spectroscopy and GDOES is glow discharge optical emission spectroscopy. To analyse inhomogeneous surface layers laterally re-solving methods like AES and SIMS have to be applied. For this paper a restriction to the electron spectroscopic methods will be carried out. More detailed information on the application of the remaining methods may be found elsevvhere1. To get a depth profile in most cases destruction of the layer to be analysed has to be performed by one of the follovving methods: a) Sputter depth profiling b) Angle lapping c) Crater edge profiling d) bali cratering For sputter depth profilling the layer is decomposed by bombardement with noble gas ions and parallel or successive analysis of the momentary interface by a surface analytical method is carried out. Angle lapping means that by using normal metal pol-ishing equipment a layer covered sample is polished un-der a very flat angle. Using polishing angles down to about 10 a spreading of the surface layer up to a factor of about 100 is possible and this spreaded part of the layer may be analysed after transfer into a surface analytical system by AES point or line analysis for example. For crater edge profiling the crater edge resulting from ion beam etching accompanying a single sputter depth profile is exploited. The crater edge of such a profile exposes the strata of the interface in a manner related to that produced by angle lapping, but the striking differ-ence is that for crater edge profiling the resulting angles between surface and interface are 3 orders of magnitude less. In the čase of bali cratering a rotating, spherical, steel bali coated vvith fine diamond paste, is used to grind a spherical crater into the sample surface and after that the sample is transferred into the surface analytical system to analyse the sputter cleaned crater vvalls. Advantages and disadvantages of the different methods are discussed in some detail in2. 3 Results a) Very thin surface layers (passive layers) The only nondestructive method to depth profile very thin surface films is by angle dependent XPS - or AES -measurements. The first example to be presented con-cerns an oxide layer vvhich vvas formed at room temperature in air on a pure zine sample. The thickness of this layer vvas assumed to be only a fevv nm's and it should be less than the information depth of the applied electron spectroscopic methods AES or XPS. The next figure 1 compares the Zn - LMM Auger spectrum of a clean zine sample (sputter cleaned by Ar+ ion bombardment) vvith the same sample after oxidation at room temperature in air. The comparison makes clear that for the oxidized sample additional features can be recognized if the analyser energy resolution is adequate (0.05%). XPS studies on the same sample lead to the results that also for the Zn 3d- and Zn 2p- photoelectron signals a distinetion betvveen metal and oxide is possible, but the difference is most pronounced for the Zn - LMM Auger signal (in this čase X-ray excited), figure 2. If the information depth for the Zn - LMM Auger signal corresponding to the oxidized state of Zn is less than the thickness of the oxide layer, angle dependent measurements should reveal a more pronounced oxide signal at lovver angles of analysis. This is illustrated by figure 3. The results of an angle dependent measurement on the X-ray excited Zn - LMM Auger signal are depieted in figure 4 and clearly demonstrate, for angles of analysis ranging from 25° to 90°, that the information deepth is less than the thickness of the oxide layer. For a smooth, homogeneous, contamination free thin oxide layer the measured signals of the metallic and ox-ide components may be used to determine the oxide thickness according to equation (1): L = Nni \m exp[-d / \m sinft] 1„* No* 1—exp[-d / sinp] (D vvhere lm, lox are the intensities of the metal and oxide signals, Nm, Nox are the densities of the metal and the oxide, Xm, \ox are the inelastic mean free pathes of the Auger electrons in the metal and in the oxide, (3 is the angle of analysis in relation to the sample surface, see figure 3. 9B0 990 1000 Kinetic Energy / eV 1010 1020 Figure I: Comparison of the Zn - LMM - Auger spectrum for a clean zine sample and after oxidation at room temperature in air for 2 vveeks Slika 1: Primerjava AES spektrov čistega Zn - LMM pred in po oksidaciji dva tedna na zraku pri sobni temperaturi The (2): above equation (1) may be changed to equation sin(3 N, N,„ k, ■X? (2) zine zine oxide zine hydroxide ^LMM 17.7 20 21 6.7 8 7 According to equation (2) a plot of 1 /sin(3 against the term on the right hand side of equation (2) should give a straight line and from the slope of this straight line the thickness may be derived. angle of analysis /\ surface layer bulk information depth I 1018 1019 1020 1021 1022 1023 1024 1025 1026 1027 10J binding energy Figure 2: Zn - 3d-. Zn - 2p - photoelectron and Zn - LMM - Auger -signal for clean and oxidized zine sample (see figure 1) Slika 2: Zn - 3d, Zn - 2p XPS signal in Zn - LMM AES signal za čisti in oksidirani cink (glej sliko 1) For a determination of the layer thickness values for the A,'s are necessary. There are no experimental data for the X - values of zine, therefore the corresponding X's of zine and zine oxide (of zine hydroxide as well, which will be needed lateron) were calculated according to a relation derived by Tanuma et al.3. The calculated values are (in A): Table 1: The calculated X values (in A) Figure 3: Shematic illustration of the higher surface sensitivity at lower angles of analysis Slika 3: Shematični prikaz višje površinske analizne občutljivosti pri nizkih kotih The angle dependence of the peak areas of the Zn -LMM Auger signal and the Zn - 2p photo electron signal are plotted in figure 5 for the metallic and the oxide con-tributions. From this plot and by using additionally the known values for Nm, Nox and X0* figure 6 can be derived. The observed straight line for the dependence of the Zn - LMM Auger signal leads to a thickness of 15A for the oxide. If we have a closer look to the well resolved oxygen O - ls photo electron signal corresponding to the oxide layer, figure 7, we can detect that the assumption of a pure oxide layer was not correct, additionally to the ox-ide signal centered at about 530 eV peak energy a second signal caused by some hydroxide contribution is observed around 532 eV peak energy. Results of angle de-pendent measurements on the oxygen O - ls signal in figure 8 indicate that we have an inner oxide layer and an outer hydroxide layer. - Zn-LMM metal i - oxide Analysenwlnkel - 24,5' 33.5' - - 42.5' 51.5' metal / 60.5- 82.5' _ 90.0' - _3-m 250 255 260 265 270 Binding Energy / eV 275 280 Figure 4: Angle dependence of the Zn - LMM Auger signal for an oxidized zine sample (as for figure 1) Slika 4: Kotna odvisnost Zn - LMM AES signala za oksidiran cink (kot na sliki 1) metal contribution in% 100 60 40 20 oo°o Zn - 2Po, ,oo O O o o o o o o o O Zn - 2pm„, o Zn - Auger O* Zn - Auger mat °o o o o OOo + + 70 50 30 angle of analysis 10 L L '-ZnOH [l-exp(dZn0H/XZn0H cos©)] / / XZn cZn exp(- A,Zn0 cos© ) exp(- I cos© (3) + o O CvJ t— _c CD CO 22 20 18- 16- 14- 12 10 - d„ 15A Zn - LMM 1,0 1,2 —I-'-1— 1,4 1,6 1 / sirili 2,0 Figure 5: Angle dependence of the Zn - LMM Auger signal and Zn -2p photoelectron signal for an oxidized zine sample (as for figure 1), metal and oxide contributions are plotted Slika 5: Kotna odvisnost Zn - LMM AES signala in Zn - 2p XPS signal za oksidiran cink (kot na sliki 1), prikazana sta deleža kovine in oksida Looking more detailed to the angle dependence of the Zn - LMM Auger signal it can be recognized, figure 9, that also for this signal the "oxide" - component ex-hibits some peak shape variations in dependence on the angle of analysis. If we try to fit the Zn - LMM Auger signal for the oxidized sample at first, figure lOb, by a metal contribution according to figure lOa and the re-maining peak area for the oxide contribution by a single peak this is not possible. We need two different peaks as shown in figure lOc, to get an acceptable fit of the whole Zn - LMM spectrum. We can novv try evaluate the thickness of the oxide and the hydroxide separately. Again we assume that the oxide and hydroxide layer are homogeneous layers. The intensity ratios of the sig-nals of the different layers with respect to the signal of the zine substrate are given according to S. Lecuyer et al.4 by: Figure 6: Determination of the oxide thickness according to equation (2), see text Slika 6: Določitev debeline oksidne plasti po enačbi (2), glej tekst MPI,DUSSELDOHF XPS - Spectrum ZNC148.DAT Region i / 1 Level 1 / 1 Rad iat ion Mg Kalpna Max Count nate 3637 CPS Ana lyser 5 eV Step Size 0.05 cV Dwell Time 100 ms No of Channels 401 No of Scans 30 Time for Region 1203 Sec Acquired 14: 32 08-Jan-96 Plotted 59 OS-Jan-96 532 534 Binding Energy / eV 538 Figure 7: XPS - O - ls signal for an oxidized zine sample (as for figure 1) Slika 7: XPS - O - ls signal za oksidiran vzorec cinka (kot na sliki 1) T2 (®) = ^no CZn0 [ l-exp(- , dzn" J] / I cz„ exp(- ^ZnO COS0 dZnO ^ZnO COS0 (4) where Ci is the concentration within layer i and dj is the layer thickness. For ali the angle dependent measurements of the Zn -LMM Auger signal recorded for the oxidized zine sam- O - 1s - Signal Hydroxid 528 532 Binding Energy eV 536 250 35 30 k 25 C 0 20 u n ib t 10 s 5 0 Zn LMM "A 20° - Me A Ok OH \ - 75° 256 258 260 262 264 266 268 Binding Energy / eV 270 272 274 Figure 9: Zn - LMM Auger signal for an oxidized zine sample (as for figure 1) recorded at 20° and 75° angle of analysis Slika 9: Zn - LMM AES signal oksidiranega cinka (kot na sliki 1) posnet pri kotih analize 20° in 75° ple a peak fitting for the total LMM signal was per-formed in a similar way as discussed before. The measured peak area ratios lox/lme and loH/lme for the oxide and the hydroxide layer in dependence on the angle of analy-sis ZN1 = 0 are listed in the table 2 together with the same peak area ratios calculated according to equation (3) and (4) by using a thickness of 13A for the oxide layer and a thickness of 4A for the hydroxide layer. if 200 C o 150 Figure 8: Angle dependent measurements for the XPS - O - ls signal of an oxidized zine sample (as for figure 1) Slika 8: Meritve XPS - O - ls signala za oksidiran vzorec cinka (kot na sliki 1) v odvisnosti od kota 100 11 t 50 S Q k 10 G 8 0 6 u n 4 f 2 s oJ k 10 C 8 0 6 u n 4 t 2 s 0 /\y metal signal metal signal fitted \ a / \ z // / metal + oxide signal metal signal fitted --b '-■"ST" • -1---- ' Zn LMM metal + oxide signal metal,oxide and /a\ hydroxide fitted /A\ Ox \\ 1 OH M e j \\ --J C Me I// \ i // y - 256 258 260 262 264 266 268 Binding Energy / eV Figure 10: a) Zn - LMM Auger signal for a clean zine sample, metal signal fitted; b) Zn - LMM Auger signal for an oxidized zine sample (as for figure 1), metal part fitted; c) Zn - LMM Auger signal for an oxidized zine sample (as for figure 1), metal oxide and hydroxide parts fitted Slika 10: a) Zn - LMM AES signal za čisti cink; b) Zn - LMM AES signal za oksidiran cink (kot na sliki 1) kovinski del se prilega; c) Zn -LMM AES signal za oksidiran cink (kot na sliki 1) kovinski in hidroksidni del se prilega krivulji Table 2: The angle dependent measurements of the Zn - LMM Auger signal for the oxide and hydroxide layer © lox/lme lox/lmc l0H/lme l0H/lme (exp.) (calc.) (exp.) (calc.) 78.1 13 14.7 11.7 10.5 76 9 9.05 5.8 4.1 73 5.6 5.44 2.8 2.2 70.5 3.5 3.9 1.8 1.65 68 3.15 3.1 1.25 1.05 64 2.3 2.25 0.8 0.7 60.5 1.9 1.8 0.6 0.54 58 1.5 1.6 0.5 0.42 45.4 1 1 0.25 0.28 32.9 0.79 0.77 0.18 0.18 21.4 0.65 0.67 0.15 0.17 13 0.66 0.63 0.14 0.14 For the first thickness determination, assuming a single oxide layer, a thickness of 15A resulted, which is in very good agreement with a thickness of 17A for the double layer consisting of a 13A thick oxide and a 4A thick hydroxide layer. b) Thick layers (high temperature oxidation) For future power plants high strength 9% Cr steels are being considered as construction materials for steam piping, headers and superheater tubes up to 600°C. For those materials it vvas found that the destruction of the protective oxide scale occurs during eposure in simulated combustion gas by the presence of vvater vapour. Al-though this detrimental effect of vvater vapour on the oxi-dation resistance of ferritic Cr - steels is knovvn already for a long time no conclusive mechanism has been eluci-dated. According to figure 11, illustrating schematically the variation vvith alloy Cr content of the oxidation rate and oxide scale structure, several different oxides are ex-pected to appear for a Fe 9% Cr alloy and the above mentioned oxidation conditions. The Fe 9% Cr alloys vvere oxidized isothermally in N2 - 1% O2 vvith and vvithout various H2O contents at 650°C. Using the crater bali equipment a crater vvas ground into the different oxidized samples. The crater bali process is illustrated in figure 12. After transferring Alloy chromium content, wt% Figure 11: Schematics of the variation vvith alloy chromium content of the oxidation rate and oxide scale structure (based on isothermal studies at 1000°C in 0.13 atm oxygen) Slika 11: Shematičen prikaz vpliva različnih vsebnosti kroma na stopnjo oksidacije in struktura oksida (izotermna oksidacija pri 1000°C in 0.13 atm kisika) the samples into the scanning Auger system the samples vvere sputter cleaned and SEM images recorded. Tvvo ex-amples vvill be presented to shovv the possibilities of this method. The first sample vvas oxidized in N2 - 1% O2 -4% H20 at 650°C for 3 hours and the second one for 10 hours, figure 13 to 16. For the application of high temperature materials the formation of even, slovv grovving and well adherend ox-ide layers are desired. AI2O3 layers vvould be the thermo-dynamically most stable layers for nearly ali conditions occuring during the use of the material. Quite often hovvever the nucleation and adherence of the AI2O3 layers are not satisfactory. The improve the adherence of those lay-ers oxidation of a model alloy Fe - 6% Al - 0,5% Ti and 50 to 100 ppm C vvere studied. After oxidation of this alloy the samples vvere inves-tigated by surface analytical methods. The oxide layers vvere fine grained, vvell adherent and represented an ex-cellent protection against carburizing atmospheres. This vvas tested by long term investigations. In order to find out the reason for this improvement of the protecting properties Auger depth profiles of the oxide layer on top of the model alloy vvere recorded. A typical example is shovvn in figure 17. The most striking feature of this depth profile is the simultaneous enrich-ment of carbon and titanium at the oxide metal matrix interface. This observation leads to the assumption that by formation of a TiC layer at the interface this improvement of the corrosion protecting properties could be real-ized. By further detailed surface analytical investigations on a single crystal model alloy of the same composition and applying LEED (lovv energy electron diffraction) and AES it could be found out that on several lovv in-dexed crystal surfaces this TiC layer grovvs epitaxially. d =--------- 8R Figure 12: Scheinatic illustration of the crater bali etching process Slika 12: Shematski prikaz kraterja dobljenega s procesom jedkanja s kroglico 400 500 600 700 400 500 600 700 Figure 13: SEM of a part of the crater etched area of the Fe 9% Cr sample oxidized for 3 h (see text) and characteristic Auger point spectra for the different areas Slika 13: SEM posnetek dela jedkalnega kraterja vzorca zlitine Fe 9% Cr, po 3 urah oksidacije (glej tekst) in karakteristični AES spektri, posneti na označenih mestih Metal dusting is known to be a dangerous high temperature corrosion phenomenon in petrochemistry and in reformer and direct reduction plants. In strongly car-burizing atmospheres and temperatures from 400 to 800°C Iow alloyed Fe, Ni and Co base alloys are subject to a catastrophic carburization leading into a desintegra-tion of the material into a dust composed of fine metal particles and carbon. For the reaction mechanism of the metal dusting process it was assumed Ihat instable car-bides form as an intermediate before they decompose to metal and carbon dust. In order to get a more detailed picture of the metal dusting process an iron sample from the initial stages of the metal dusting process (680°C, 78% H2, 15% CO and 0.5% H2O) vvas removed from the carburizing atmos- Figure 14: a) oxygen; b) chromium; c) and iron - images of the same surface area as shovvn in the SEM image in figure 13 Slika 14: a) kisik; b) krom; c) in železo - slike delov površin prikazanih na SEM posnetku slike 13 phere and using the crater bali equipment a crater vvas etched into the sample surface. After transferring the etched sample into the scanning Auger system a clean surface of the crater area vvas produced by Ar ion bom-bardement. Immediately aftervvards Auger spectra and Auger images vvere recorded. The different chemical states of carbon vvithin graphite and in carbide may eas-ily be distinguished by Auger electron spectroscopy be-cause of a characteristic Auger signal peak shape for the individual compounds. This is demonstrated by the Auger spectra in figure 18, vvhich vvere recorded for different areas of the etched crater. Because of the different peak shape and a slight difference in peak energy graphite and carbide may be imaged separately. This is illus-trated by the follovving figures, figure 19a to figure 19d, vvhich shovv additional to a SEM image Auger elemental maps for carbidic carbon, carbon in graphitic form and iron of the crater region. 400 500 600 "00 400 500 600 700 Figure 15: SEM of a part of the crater etched area of the Fe 9% Cr sample oxidized for 10 h (see text) and characteristic Auger point spectra for the different areas Slika 15: SEM posnetek dela jedkalnega kraterja zlitine Fe 9% Cr, po 10 urah oksidacije (glej tekst) in karakteristični AES spektri, posneti na označenih mestih Figure 16: a) oxygen; b) chromium; c) and iron - images of the same surface area as shown in the SEM image in figure 15 Slika 16: a) kisik; b) krom; c) in železo - slike istih delov površin prikazanih na SEM posnetku slike 15 The sum up the results from different kinds of inves-tigations the following development of surface layers during the metal dusting process can be derived (sche-matically): surface surface surface surface metal (Fe) +C diss. Fe3C graphite coke(Fe+C) metal(Fe) +C diss. Fe3C graphite metal(Fe) +C diss. Fe3C metal (Fe) +C diss. Well adherent and corrosion protecting oxide layers are of great importance for high temperature alloys. The adherence of the oxide layers is affected by the morphol-ogy and the chemical composition of the oxide/metal in-terface. In this study scanning Auger microscopy (SAM) is used to investigate the oxide/alloy interface of oxi-dized Fe-Cr-Al alloys (undoped or doped with Ti, Ce and C'i 20)' C (2) 0 • , . A A AAA AA A AA • . h 4» ■ A ' • • "i." . •'fe "' . Al A ■ * -1-,-T^-1-1- , ~ 20 40 60 80 100 120 KO Sputlertime (min) Figure 17: AES depth profile of a polycrystalline Fe-6Al-0,5Ti-0,01C sample oxidized for h at 1000°C and 10"19 bar oxygen partial pressure Slika 17: AES profilni diagram polikristalne zlitine Fe-6Al-0,5Ti-0,01 C po 1/2 urni oksidaciji na temperaturi 1000°C in parcialnemu tlaku kisika 10'19 bar HPI DUSSELD0RF Auger - Spectrum V.G.Scientific CLU100.DAT Region / 1 Level 1 / 1 Point i / 7 92 p 1 graph K 90 r ea ■-'graph u \ n 84 ^ 80 - 240 250 260 270 280 290 Kinetic Energy / eV HPI DUSSELDORF Auger - Spectrum V.G.Scientific CLU100.0AT Region / i Level 1 / 1 Point i / 7 Rad lat Ion Elactron Man Count Rata 1373161 CPS CRfl 9tap Sira 0.70 aV Dwall Tina 92 K 90 C 88 0 86 Bcarb/ D cart) 2015 No of Scana n 84 Tl«a for Raglon 302 Sac Acgulrad 11:31 04-Aug-94 Plottad 07: 42 08-Aug-94 t s02 80 250 260 270 280 Kinetic Energy / eV Radlatlon Elactron M«x Couot Rate 1373161 CPS CRR Stap Sira 0.70 aV D»all T1m of Channala S of Scana TlM for Raglon 302 Sac Acgulrad 11:31 04-Aug-94 Plot tad 07: 42 08-Aug-94 Figure 18: 'graphitie' (upper spectrum) and 'carbidic' carbon - KVV Auger signal of a crater bali etched Fe metal dusting sample indicating the different peak shapes and illustrating the peak (P,) and background (Bi) energy positions for recording the Auger images Slika 18: 'grafitni' (zgornji spekter) in 'karbidni ogljik' - KVV Augerjev signal s kroglico jedkanega Fe prašnatega vzorca z različnimi oblikami vrhov prikazuje vrh (Pj) in ozadje (Bi) in energijske pozicije za posnete AES spektre Figure 19: a) to d) SEM and SAM images of a crater bali etched and sputter cleaned Fe metal dusting sample Slika 19: a) do d) SEM in SAM posnetki s kroglico jedkanega Fe prašnatega vzorca in z Ar+ ioni jedkan Fe prašnat vzorec Figure 20: a) to c) SEM images of surface areas where the oxide layer is partly (or completely) removed Slika 20: a) do c) SEM posnetki površine, kjer je bila oksidna plast delno (ali popolnoma) odstranjena Figure 21: Sulfur image of the same surface area as shown in the SEM image of figure 20 c) Slika 21: Posnetek žvepla na površini, prikazani na sliki 20 c) 200 400 600 800 1000 1200 1400 Elektronen-Energie [eV] Figure 22: Auger point spectra of a) a void area; b) a rugged surface area Slika 22: AES spekter posnet a) v vrzeli; b) na hrapavi površini Y) after partly removing the oxide layer by in situ bend-ing. Thin Fe-Cr-Al ribbons, doped and undoped, were produced by meltspinning. Rectangular specimens were cut from the ribbons and ultrasonically cleaned in ace-tone. The samples were oxidized at 1273 K in a control- 200 400 600 800 1000 1200 1400 Elektronen-Energie [eV] led He - O2 gas mixture at an oxygen partial pressure of 133 mbar. Bending of the specimens was performed in UHV at a residual pressure of 5 x 10~9 Pa to spali off parts of the oxide layer. The stripped oxide/metal interface was in-vestigated by in situ SEM and SAM. The metal surface shows individual voids and rugged parts, indicating imprints of the removed oxide, see SEM figures 20a to 20c. For the undoped alloy poor adher-ence of the oxide layer is observed. Sulphur is strongly enriched at the surface of the voids, figure 21 shows a sulphur image of the sample area as for the SEM image in figure 20c and 22a and 22b Auger point spectra of a void surface and a rugged part of the interface. On the Ti containing alloys the oxide layer is again poorly adherent. Sulphur is also strongly enriched at the surface of voids. On the Y- and Ce- containing alloys the oxide layer is vvell adherent and the sulphur concentration is belovv the detection limit. The poor adherence of the oxide layers on undoped and Ti doped Fe-Cr-Al alloys is correlated to the pres-ence of sulphur at the alloy surface. Sulphur enrichment is explained by sulphur segregation to the free alloy surface and by additional sulphide formation for the Ti containing alloys. The positive effect of Y and Ce on the ad- herence of the oxide layers is explained by sulphide pre-cipitation in the bulk and thus preventing sulphur segregation to the free surface of the alloy. 4 Summary Depending on the thickness of the surface layers to be analysed by surface analytical methods vvith respect to composition and thickness, different methods of depth profiling have to be applied. Valuable information may be derived from the results of the surface analytical in-vestigations, leading to a better understanding of the grovvth mechanisms and the corrosion protecting properties of oxide layers on metal surfaces. 5 References 'Guidelines for methods of testing and research in high temperature corrosion, European Federation of Corrosion Publications 14, The Institute of Materials, London, 1995, 189 2 in Practical Surface Analysis, eds. D. Briggs and M. P. Seah, Wiley, Chichester, 1990 3S. Tanuma, C. J. Povvell and D. R. Penn, Surf. Interface Anal., 20a, 1993, 77 4 S. Lecuyer, A. Quemerais and G. Jezequel, Surf. Interface Anal., 18, 1992, 257 IMT INŠTITUT ZA KOVINSKE MATERIALE IN TEHNOLOGIJE INSTITUTE OF METALS AND TECHNOLOGY 1001 LJUBLJANA, LEPI POT 11, SLOVENIJA, POB 431 Phone.: +386 61/125 11 61, Fax: +38661/213 780 VACUUM HEAT TREATMENT LABORATORY Vacuum Brazing Universally accepted as the most versatile method of joining metals. Vacuum Brazing is a precision metal joining technique suitable for many component configurations in a wide range of materials. ADVANTAGES • Flux free process yields clean, high integrity joints • Reproducible quality • Components of dissimilar geometry or material type may be joined • Uniform heating & cooling rates minimise distortion • Fluxless brazing alloys ensure strong defect free joints • Bright surface that dispense vvith expensive post cleaning operations • Cost effective Over five years of Vacuum Brazing expertise at IMT has created an unrivalled reputation for excellence and quality. Our experience in value engineering vvill often lead to the use of Vacuum Brazing as a cost effective solution to modern technical problems in joining. INDUSTRIES • Aerospace • Mechanical • Electronics • Hydraulics • Pneumatics • Marine • Nuclear • Automotive QUALITY ASSURANCE Quality is fundamental to the IMT philosophy. The choice of process, ali processing operations and process control are continuously monitored by IMT Quality Control Department. The high level of quality resulting from this tightly organised activity is recognised by government authorities, industry and International companies. Aluminium and Magnesium Based Metal Matrix Composites Kompoziti na osnovi Al in Mg K. U. Kainer1, Technische Universitat Clausthal-Zellerfeld, Germany Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 In motor-cars metal matrix composites (MMCs) are employed in braking systems and engine components. Other applications for these materials have been developed in energy and in information applications. The potentional of composite materials is very great because the properties can be tailored according to the application. There are many possible material combinations and processing techniques which can be employed. For structural applications standard iight metal are often strengthened by ceramic fibres or particles. The performance and potential of composites will be discussed using examples of reinforced aluminium and magnesium alloys. Key words: Al and Mg based MMCs, Al and Mg alloys, SiC and Alz03 as reinforcement, properties, applications, processing techniques Kompoziti s kovinsko osnovo, ojačani s keramičnimi delci ali vlakni se danes že uporabljajo kot deli zavornega sistema in motorjev z notranjim izgorevanjem. Razvite so tudi posebne vrste teh materialov, ki so uporabni na področju energetike in informatike. Uporabnost kompozitnih materialov je vsestranska, ker lahko njihove lastnosti prilagajamo potrebam uporabe. Možne so številne kombinacije materialov (kovinska osnova / keramična ojačitev) in postopkov njihove izdelave. Kompoziti, ki se uporabljajo kot konstrukcijski materiali so najpogosteje sestavljeni iz lahke kovinske osnove in keramičnih delcev ali vlaken. Prispevek obravnava predvsem dosežene lastnosti in možnosti uporabe kompozitov z osnovo iz Al ali Mg zlitin, ki so diskontinuirno ojačane z SiC ali AI2O3 delci oziroma kratkimi vlakni. Ključne besede: kompoziti s kovinsko osnovo, Al in Mg zlitine, SiC in AI2O3 kot ojačitvena faza, lastnosti, uporaba, postopki izdelave 1 Introduction The strenuous efforts to develop metal matrix composites vvith light metal matrices in the eighties have paid off vvith successful applications in automobile and transport systems. Worthy of mention are partially reinforced pistons, hybrid reinforced engine blocks for cars or trucks as vvell as particle reinforced brake discs for light lorries, motocycles, cars or rail vehicles. Further fields of applications are military aircraft and space craft. The in-novative materials are interesting possibilities in the development of modern materials because the properties of MMCs can be tailored for a particular application and hence MMCs can fulfill ali requirements of the designer. Such materials become important vvhen the property profile cannot be achieved by the conventional light metal alloys. The specific strength as outstanding advantage of light metal MMCs is hovvever under pressure from com-peting technologies such as povvder meta!lurgy of polmer technology. The advantages of composites are only realised if a resonable cost performance ratio is achiev-able on production of the component. In this respect it is important for economic and ecological reasons to recycle scrap components, production vvaste, etc. The aims in the reinforcement of metal matrix func-tional or structural materials are on the one hand the op-timisation of some critical properties at the same time as maintaining other properties and on the other hand a 1 Dr.-Ing. habil. K. U. KAINER Tcchniscbe Universitat Clauslhal ■Institut lur VVerkštofflcunde und VVerkstofftechnik. Clausthal-Zellerfeld. Germany complete change in the property profile of a class of materials. The reinforcement of light metals opens, for ex-ample, an extension of the application potential vvhere vveight reduetion of components is very desirable at the same time as optimisation of component properties. The development aims of light metal matrix composites thus can be sumarised as follovvs: - inerease in yield strength, ultimate tensile strength and fatigue strength at room temperature vvhilst maintaining minimum values of ductility or toughness, - inerease in hot strength, fatigue strength and creep resistance at elevated temperatures compared to conventional materials, - reduetion in the coefficient of thermal expansion of light metal alloys to values comparable vvith steels, - improvement in the stability of light metals to temperature changes, - improvement in damping behaviour, - improvement in the vvear resistance through addition of hard materials, - improvement in vveight specific properties (strength and E-modulus). Discontinuous particle, fibre or vvhisker reinforced light alloys are most likely to fulfill design eriteria because the components are relatively cheap and production of components in large numbers is possible. Further advantages are the relatively high isotropy of properties compared to the long continuous fibre reinforced light metals and the possibility of further forming by forging and machining. 2 Combination of materials for light alloy compos-ites The obvious candidates for light metal matrices for composite materials are the easily workable, conven-tional a!loys. Particulary when powder metallurgical (P/M) production techniques are employed it is possible to consider special alloys with specific compositions. P/M technology allows the use of alloys with super sa-torated or metastable phases. The alloys are free from segregation problems as often observed after conven-tional solidification. Examples of extensively investigated matrix alloys are1"8: Conventional Casting Allovs: Al allovs: AlSil2CuMgNi AlSi9Mg AlSi7 (A 356) Mg allovs: MgA19Znl (AZ91) MgA12RE2Zrl (MSR, QE 22) Conventional wrought allovs: Al allovs: AlMgSiCu (6061) AlCuSiMn (2014) AlZnMgCul 5 (7075) Mg allovs: MgA13Zn (AZ 31) MgZn6Zr (ZK 60) MgZn6Cu3 (ZC 63) Special allovs: Al allovs: Al-Cu-Mg-Li (8090) Mg allovs: Mg99,5 + RE, Ca, Zr, Ba. Br, Sb or Sn (1-2.4%) A wide variety of reinforcement materials are avail-able vvith a vvider range of properties. The choice depends on the method chosen for production and on the matrtx alloy system. In general the requirements are: - lovv density, - mechanical compatability (a thermal coefficient of expansion vvhich matches the matrix), - chemical compatability, - thermal štabi lity, - high elastic modulus, high compressive and tensile strength, - good workability, - economicy. These demands can be fulfilled virtually only by in-organic reinforcing materials. Often only ceramic particles or fibres or carbon fibres are used to reinforce met-als. The use of metallic fibres results in prohibitive increases in density. Which component is chosen depends on the matrix material and the property profile of the particular application. Information of available particles, short fibres, vvhiskers and continuous fibres for reinforcement of metals is collected in Table 1 and in references910. The preparation, vvorking and means of applications of the various reinforcements depends on the method chosen to produce the composite (see1). A combinated application of tvvo and more reinforcement material is possible (hybrid technique)1-9. 3 Production of light metal composites There are several possible methods of producing semi finished material and components in light metal composites, vvhich depend primarly on the component geometry and the material systems (matrix / reinforcement). The process must be divided into preparation of suitable starting material, production of the semi finished material or component and finishing operations. For eco-nomic reasons near net shape production should be at-tempted to minimise mechanical finishing operations. In general the follovving production techniques are available: • Casting techniques - infiltration of short fibres, particle or hybrid pre-forms by squeeze casting, vacuum infiltration or pressure infiltration1'4'7'8, Table 1: Examples of particles, whisker, continuous and discontinuous fibres used a reinforcements in metal alloys (*CTE = coefficient of thermal expansion, nPAN based fibres, 2lpitch based fibres) reinforcement producer diameter densitv E-modulus tensile strength CTE* (10"6K"') (pm) (gcm5) (GPa) (MPa) axial FP OC-AI2O3 Du Pont 20 3.9 380 > 1400 7.6 Altex alumina fibre Sumitomo 17 3.2 300 2000 8.8 Nicalon SiC-fibre Nippon Carbon 15 2.6 185 2700 3.5 Torayca T-300" Toray 7 1.8 230 3530 -0.26 Torayca M-4011 Toray 5.5 1.8 392 2650 -1.3 Thornel P 752) Amoco 10 2.0 520 2370 -1.4 Saffil RF disk a-Al203 ICI plc. 1-5 3.3 300 2000 4.7 SiC-whiskers Silar DWA Composites Specialities 0.6 3.2 690 6900 4.1 SiC-particles Norton AS, ESK. Kempten various 3.2 ca.400 - 4.7 alumina platelets Elf, ESK, Kempten various 3.9 ca.380 - 3.6 alumina particles H.C. Starck, ESK. Kempten various 4.0 ca.380 - 9.5 'P' 4 9 * Figure 1: Collection of typical microstructures of various light metal composites as a function of reinforcement and production process. a) AI2O3 short fibre reinforced magnesium. b) AbO.i-SiC hybrid reinforced magnesium, c) SiC particle reinforced aluminium (chill čast), d) SiC particle reinforced aluminium (pressure die čast), e) SiC particle reinforced aluminium (čast and e.\truded), f) SiC particle reinforced aluminium (extruded povvder blend). g) SiC particle reinforced magnesium (spray formed), h) SiC particle reinforced magnesium (spray formed and extruded) 11 12 - reaction infiltration of fibre or particle preforms ' , - production of prematerial by stirring particles into metallic melts vvith subsequent sand casting, chill casting or pressure casting2,3. • Powder metallurgy techniques - extrusion or forging of metal povvder - particle mix-tures5'6, - extrusion or forging of spray formed semi finished material1'13'14. • Further processing of semi finished čast material by thixocasting or forming, extrusion15, forging, cold forming or superplastic forming, • Joining or vvelding of semi finished products, • Finishing by machining. 4 Structure and properties of light metal composites The structure of composites is determined by the na-ture and shape of the reinforcing components, their distribution and orientation by the production process. Typi-cal microstructures of various short fibre and particle reinforced light metals are shown in Figure 1. In the čase of short fibre reinforced composites a planar iso-tropic distribution of the short fibres is formed as a result of the production of the fibre preform. The pressure sup-ported sedimentation technque leads to a layer like structure (Figures la & b)10. The direction of infiltration is generally normal to these planeš. The čast particle reinforced light metals show, depending on the vvorking processing, typical particle distributions. Gravity čast material exhibit as a result of the casting conditions particle free regions (Figure lc), vvhereas pressure die čast A 500 i 400 E E Z 300 si CD S 200 -4—» (/) U 100 0 0 0 100 200 300 400 °c temperature -► Figure 2: Comparison of the temperature dependence of the tensile strength of the unreinforced and reinforced piston alloy AlSil2CuMgNi (KS 1275)7 a) KS 1275 vvith 20 vol.% SiC vvhiskers, b) KS 1275 vvith 20 vol.% AI2O3 short fibres, c) KS 1275 unreinforced materials shovv a much better particle distribution (Figure ld). An even distribution is achieved by extrusion of semi finished material (Figure le). An extremely homo-geneous particle distribution is obtained by extrusion of mixed povvders or spray formed materials (Figures 1 f-g)- Properties of short fibre reinforced aluminium An increase in strength with increasing fibre content in short fibre reinforced aluminium is actually observed as the example AlSil2CuMgNi vvith 20 vol.% AI2O3 shovvs in Figure 2. Composites of light metal casting al-loys is not made just to increase only the strength. The effect alone vvould not be justifiable economically. The improvement of the properties at high temperature vvith a doubling of the strength (Figure 2) and the rotating bending fatigue strength at 300°C (Figure 3), opens up possibilities for use as piston material or cylinder liners. A dramatic increase of the thermal shock resistance can be achieved at temperature of 350°C as is shovvn in Figure 4. Properties of particle reinforced aluminium In general addition of particles to light metals, such as magnesium and aluminium increases the elastic modulus, yield strength, ultimate tensile strength, the hardness and the vvear resistance and also decreases the coefficient of the thermal expansion. The degree of improvement of these properties depends on the volume fration of the particles and the chosen means of production. Tables 2 and 3 shovv a collection of properties of various particle reinforced aluminium alloys. The particle volume fraction in stirred in particle reinforced Al al-loys is limited to about 20 vol.%. This limit is imposed by the process. A maximum tensile strength of over 500 MPa and E-moduli of 100 GPa are possible for this particle content. Higher particle contents can be achieved by £ 150 a C S! «1 ra o o. 3 S s —100 ™ 5 - n | b S 50 O) C n o i. 0 0 100 200 300 °C 400 temperature Figure 3: Change in the rotating bending fatigue strength of the unreinforced and reinforced (20 vol.% AI2O3) piston alloy AlSil2CuMgNi (KS 1275) vvith increasing temperature® (GK = chill čast GP = squeeze čast) - 1- 1 1 l Ng=2.5* 10' P„=50% N . N ' N ' . ............ "............GP - KS1275/20%AI20^ GP -KS1275 ---GK -KS1275 "*-• 1 i i i O 2000 4000 6000 1000 3000 5000 temperature cycls-► Figure 4: Temperature shock resistance of the fibre reinforced piston alloy AlSil2CuMgNi as a function of the fibre content for a temperature of 350°C7: a) unreinforced. b) 12 vol.% AI2O3 short fibres, c) 17.5 vol.% AJ2O3 short fibres, d) 20 vol.% AI2O3 short fibres infiltration of particle preforms with higher particle volume fraction. The materials then assume increasingly the characteristics of ceramics. On tensile loading premature failure occurs. The small thermal expansion is an excel-lant characteristic despite the metallic features. There is a limit to the particle content of about 13-15 vol.% also for spray formed materials. The use of special alloys e.g. with lithium additions can nevertheless lead to high specific properties. If povvder metallurgical tech-nique involving extrusion and forging are applied then the particle content can be increased to more than 40 vol.%. As the result of high particle content and the re-sultant fine grain of the matrix very high strength of up to 760 MPa, very high E-Moduli of 125 GPa and lovv coefficients of expansion of 17 x 106K"' can be achieved. Unfortunately the elongation to fracture and the fracture toughness deteriorate. The values lie, hovv-ever, in part above those for casting alloys. Properties of discontinuously reinforced magnesium al- loys In general the strengthening effects in discontiuous reinforced composites is smaller than in continuous fibre reinforced materials but the properties are more iso-tropic. In the follovving the properties of short fibre or particle reinforced magnesium composites are listed. Table 4 shows the 0.2 yield strength, the temperature de-pendence of the ultimate tensile strength and the ductil-ity of different magnesium alloys reinforced vvith 20 vol.% Saffil short fibre. The properties are compared vvith of these of unreinforced alloys. Information about hardness, Young's modulus and coefficient of thermal ex-pansion (CTE) are included. The results shovv that the main advantages of this type of composite material are the high specific strength at elevated temperatures, the increase of Young's modulus and the reduction of the CTE. The improvement of the properties depends on the volume content of the short fibres. In the range of 15 -22 vol.% short fibres the most promising properties vvere measured416. With a higher fibre content problems in the infiltration arises vvhich reduces the strength and ductil-ity of the composites. Table 2: Selected properties of typical čast aluminium composites, prepared b^chill, pressure die casting or reaction infiltration2,3,11. (T6 = solution annealed and aged, T5 = aged; 'after ASTM G-77: čast iron 0.66 mm3; *CTE = coefficient of thermal expansion, a) after ASTM E-399 and B-645; b) after ASTM E-23), n.i. = no information Material Identification Composition Yield stress (MPa) Tensile strength (MPa) Elongation to fracture (%) Young's modulus (GPa) a) Fracture toughness. b) impact strength VVear' volume decrease (mm3) Thermal conductivity 22°C (cal/cm s K) CTE" 50- 100°C (KTK'1) Gravity casting (chill casting) a) (MPa m1'2) A356-T6 AlSi7Mg 200 276 6.0 75.2 17.4 0.18 0.360 21.4 F3S.10S-T6 AlSi9MglOSiC 303 338 1.2 86.9 17.4 n.i. n.i. 20.7 F3S.10S-T6 AlSi9Mg20SiC 338 359 0.4 98.6 15.9 0.02 0.442 17.5 F3K.10S-T6 AlSilOCuMgNilOSiC 359 372 0.3 87.6 n.i. n.i. n.i. 20.2 F3K.20S-T6 AISi 10CuMgNi20SiC 372 372 0.0 101 n.i. n.i. 0.346 17.8 Die casting b) (J) A390 AlSil7Cu5Mg 241 283 3.5 71.0 1.4 0.18 0.360 21.4 F3D.10S-T5 AISi 1 OCuMnNi 1 OSiC 331 372 1.2 93.8 1.4 n.i. 0.296 19.3 F3D.20S-T5 AlSil0CuMnNi20SiC 400 400 0.0 113.8 0.7 0.018 0.344 16.9 F3N.10S-T5 AlSilOCuMnMglOSiC 317 352 0.5 91.0 1.4 n.i. 0.384 21.4 F3N.20S-T5 AISi 10CuMnMg20SiC 338 365 0.3 108.2 0.7 0.018 0.401 16.6 Reaction infiltration Bending strength (MPa) Density (g/cm3) a) (MPa m"2 ) MCX-693™ Al+55-70 % SiC 300 2.98 255 9.0 n.i. 0.430 6.4 M.CX-724™ Al+55-70 % SiC 350 2.94 226 9.4 n.i. 0.394 7.2 MCX-736™ Al+55-70 % SiC 330 2.96 225 9.5 n.i. 0.382 7.3 Table 3: Properties of aluminium wrought alloy composites, manufactures information after5-613-15. (T6 = solution annealed and aged), "after ASTM G-77: čast iron 0,66 mm3; **CTE = coefficient of thermal expansion, n.i. = no information. Material Identification Yield stress (MPa) Composition Tensile strength (MPa) Elongation lo fracture (%) Young's modulus (GPa) a) Fracture toughness. b) impact strength Wear volume decrease (mm3) Thermal conductivity 22°C (cal/cm s K) CTE" 50-100°C (10"6K"') Čast starting material (extruded or forged) 6061-T6 AlMglSiCu 355 375 13 75 30 10 0.408 23.4 6061-T6 + 10% AI2O3 335 385 7 83 24 0.04 0.384 20.9 6061-T6 + 15% AI2O3 340 385 5 88 22 0.02 0.336 19.8 6061-T6 + 20% AI2O3 365 405 3 95 21 0.015 n.i. n.i. Powder metallursicallv prepared starting material (extruded) 6061-T6 AlMglSiCu 276 310 15 69.0 n.i. n.i. n.i. 23.0 6061-T6 + 20% SiC 397 448 4.1 103.4 n.i. n.i. n.i. 15.3 6061-T6 + 30% SiC 407 496 3.0 120.7 n.i. n.i. n.i. 13.8 7090-T6 AlZn8Mg2Col.5Cul 586 627 10.0 73.8 n.i. n.i. n.i. n.i. 7090-T6 + 30% SiC 676 759 1.2 124.1 n.i. n.i. n.i. n.i. 6092-T6 AlMglCulSil7.5SiC 448 510 8.0 103.0 n.i. n.i. n.i. n.i. 6092-T6 AlMglCulSi25SiC 530 565 4.0 117.0 20.3 n.i. n.i. n.i. Spray formed starting material (extruded) 6061-T6 + 15% AI2O3 317 359 5 87.6 n.i. n.i. n.i. n.i. 2618-T6 + 13% SiC 333 450 n.i. 89.0 n.i. n.i. n.i. 19.0 8090-T6 AlLi2.5CuMg 480 550 n.i. 79.5 n.i. n.i. n.i. 22.9 8090-T6 + 12% SiC 486 529 n.i. 100,1 n.i. n.i. n.i. 19.3 Table 4: Properties of short as čast fibre reinforced magnesium composites (CTE = coefficient of thermal expansion, n.d. = not determined, rt = room temperature, 0.2 YS = 0.2 yield strength. UTS = ultimate tensile strength)4 Cp-Mg AS 41 AZ 91 QE 22 matrix comp. matrix comp. matrix comp. matrix comp. 0.2 YS (MPa) (rt) 70 220 125 240 160 230 180 250 UTS (MPa) (rt) 80 240 193 270 220 280 250 300 Elongation (%) (rt) 5.0 2.2 9.0 1.0 4.8 1.8 4.5 1.6 Young's modulus 46 56 49.8 77.7 46 64 46 74 (GPa) UTS (100°C) (MPa) 65 240 175 250 200 270 240 285 UTS (200°C) (MPa) 45 180 150 240 120 220 200 245 UTS (300°C) (MPa) 30 120 n.d. n.d. 60 130 125 180 Vickers hardness 40 75 n.d. n.d. 65 140 75 125 HV10 (kp/mm2) CTE (10"6K"')* 26.5 21.5 24.0 18.0 27.0 20.5 26.0 20.0 The second group of discontinuous reinforced composites are particle reinforced magnesium alloys. The high range of properties is achieved by the limitless vari-ation possibilities of alloys, type of particle and produc-tion techniques. In general only a modest improvement in the strength by addition of particles is observed. But with the increase in hardness, wear resistance and Young's modulus together with the reduction of the CTE the material becomes interesting for commercial applica-tion17. The Tables 5 and 6 show the propety profiles of different produced particle reinforced magnesium composites. The SiC particles used for composite materials in Table 5 have irregular blocky shape. These particles were treated to achieve a smooth surface without sharp tips. The result are composites with high strength and very good ductility combinated with high hardness, Young's modulus and low CTE values. The P/M produc-tion technique unfluences the properties of the particle reinforced composites, as shown in Table 6. The highest strength but with low ductility is measured for spray formed and extruded composites. The best properties were achieved for direct powder forged composites, a near net shape production technique. With a special pre-form technique it is possible to produce particle or hy-brid reinforced composites by squeeze casting. The properties of material system investigated are listed in Table 5. As reinforcement a SiC-particles-fibre hybrid preform and aluminia platelets were used. The material shows lower strength and ductility due to the solidification mi- Table 5: Properties profile of P/M produced or squeeze čast QE 22 composites with different additions of reinforcement (SiC-particles, hybrid SiC-AbCb-preforms, Al203-platelets in vol.%)17 0.2 UTS Elonga- Young's Brinell CTE rt-yield (MPa) tion to modulus hardness 30Q°C strength fracture (GPa) HB31,25/(10"6K"') _(MPa)_(%}_2J_ Powder metallurgy produced composites (T6) condition_ P/M QE 22 - T6 175 260 18 43 70 27.1 QE 22 + 10% SiC 200 265 10 48 87 21.4 QE 22 + 15% SiC 210 290 10 58 95 20.0 QE 22 + 20% SiC 225 315 6.5 66 120 18.2 QE 22 + 25% SiC 245 325 4.0 73 108 16.6 Sgueeze čast composites Sq/C QE 22 - T6 185 262 5.2 69 48 27.0 QE 22+20%SiC 265 285 2.4 74 120 18.9 hybrid QE 22+25%SiC 270 282 1.0 80 125 17.5 hybrid QE 22 + 20% 177 250 1.0 85 110 19.8 AhO? platelets crostructure which is different to the rapid solidified structure by use of P/M technologies. Table 6: Influence of the production technique on the properties of P/M QE 22+15 vol% SiC-particles-composites Unrein-forced QE22 Spray formed and extruded Extruded powder blends Forged powder blends 0.2 yield strength (MPa) 180 300 250 220 UTS (MPa) 252 320 300 300 Elongation to fracture (%) 16.0 1.0 4.0 4.5 Vickers hardness (HV10 82 92 88 94 kp/mm2) Young's modulus (GPa) 46 69 70 79 CTE (10"6K"') 27.1 20.5 21.1 20.8 5 Possible uses and applications for metallic matrix composites Light metal composites are interesting materials for automobile components in the engine (oscillating parts: valve system, connecting rod, pistons and piston pin; covers: cylinder head, crankshaft main bearing; motor block: partially reinforced cylinder liner). An example for a successful application involving aluminium composites is the partially reinforced short ftbre aluminium pistons in which the combustion chamber is reinforced with AI2O3 short fibres. Comparable component properties are only possible in powder metallurgical produced aluminium alloys or in iron pistons. The reason for the use of composites are, as explained above, improved high temperature properties. Similar considerations ap-ply to partially reinforced cylinder blocks. In this čase the critical areas, the bridges and cylinder surfaces are reinforced. The same applies to the reinforcement of aluminium cylinder heads where cracking in the combustion chamber is the Iife limiting factor. Figure 5 shows the development goal on increasing the component temperature for reinforced aluminium cylinder heads. % 100 80 60 a) I 40 CD ~ 20 I N. \ \ \ a S \ b 220 240 260 280 300 part temperature 340 Figure 5: Component Iife for aluminium cylinder heads for car diesel engines (The Iife fimiting factor is cracking in the combustion chamber area)7 Potential applications can be found also in the pro-pulsive components e.g. transverse link and particle reinforced brake discs. The latter are also employed in rail transport (tube trains, railway trains). In air and space applications, the high strength, the high E-modulus, the low thermal coefficient of expansion, the temperature stabillity and the high conductivity of reinforced light metals compared to polymer materials make composites interesting for stiffening parts, load bearing tubes, rotors, covers, and containers and supports for electronic de-vices. A collection of potential actual applications of the various metal matrix components (MMCs) is given in Table 7. In history, the first technical applications of MMCs were in the fields of energy and information engineering, e.g. carbon brushes (Cu-graphite) or contact materials. There is stili scope for further development in conductor materials, support materials for printed circuits or structures for electronic components. Further economically interesting applications are to be found in leisure applications e.g. extruded and welded particle reinforced alu- Table 7.1: Potential and actual technological applications of metal matrix composites (part 1) Application Reguired propertv Material system Production method automobile and commercial vehicles stiffeners, connecting rod, frames, piston, piston pins, valve spring retainer, brake disks, brake, brake linings, drive shaft. accumulator plate_ high specific strength and stiffness, temperature stability, low coefficient of thermal expansion, wear resistance, thermal conductivity. high stiffness, creep resistance Al-SiC, AI-AI2O3, Mg-SiC, Mg-Al203, discontinuous reinforcements. Pb-C, Pb-A12Q3 melt infiltration, extrusion, forging, gravity casting, pressure die casting, squeeze-casting. melt infiltration militarv and civil aircraft supporting tubes, stiffeners, high specific strength and wings- and gear boxes, ventila- stiffness, temperature stability, tion and compressor blades. fracture toughness, fatigue resistance turbine blade high specific strength and stiffness, temperature stability, fracture toughness, fatigue ____resistant. Al-B, Al-SiC, Al-C, Ti-SiC, AI-AI2O3, Mg-Al203, Mg-C continuous and discontinuous reinforcements. W, superalloys, intermetallics e.g. NijAl, Ni-Ni3Nb melt infiltration, hot pressing, diffusion welding and soldering, extrusion, squeeze-casting. melt infiltration, directional solidification of near net shape components Table 7.2: Potential and actual teehnological applications of metal matrix composites (Part 2) Application Required property Material system Production method space frames, stiffeners, antennas, high specific strength and joints, bolds. stiffness, temperature stability, lovv coefficient of thermal expansion, thermal conductivity Al-SiC, Al-B, Mg-C, Al-C, Al-A];C>3, continuous and discontinuous reinforcements. melt infiltration, extrusion, diffusion bonding and joining (spatial structures) energy eneineering (electrical contacts and conductive material) carbon brushes electrical contacts superconductor high electrical and thermal conductivity wear resistance high electrical conductivity, temperature and corrosion resistance, switch capacity, resistance to burn. superconductivity, mechanical strength, ductility._ Cu-C Cu-C, Ag-Al203, Ag-C, Ag-Sn02, Ag-Ni Cu-Nb, Cu-NbjSn. Cu-YBaCO melt infiltration, powder metallurgy. melt infiltration, powder metallurgy, extrusion, hot pressing extrusion, powder metallurgy, coating technigues._ other applications spot welding electrodes bearings resistance to burn. load bearing capacity, wear resistance. Cu-W Pb-C, bronze-Teflon powder metallurgy, infiltration. powder metallurgy, infiltration minium-mountain bike frames and golf clubs with particle reinforced inserts. Baseball bats are another possible application because the higher damping would result in a completely different striking behaviour. 6 Recycling The necessity of integrating production waste and scrap of newly developed materials is of particular im-portance. Since ceramic materials are used in the form of particles, short fibres or continuous fibres as reinforcement it is not possible to separate the components vvith aim of reutilising of matrix and the reinforcement. But conventional melting techniques can be employed to re-cover the matrix alloy. In the čase of čast or povvder metallurgical^ produced discontinuously reinforced light metals (short fibre or particle) it is possible under certain conditions to reuse the svvarf. This is particular so for particle reinforced aluminium casting alloys where no problems arise by remelting the svvarf and directly use of the čast ingots vvithout modification. The paper18 pro-vides an overvievv of the various recycling concepts for light alloy matrix composites taking into account alloy composition, reinforcement type and the production and vvorking history. 7 Conclusion The development of metal matrix composites can be used to improve critical properties of metal alloys e.g. high temperature strength, stiffness, vvear resistance and thermal expansion. With high variability of materials combination and manufacturing techniques it is possible to produce tailor-made materials. Which combination and production techniques are choosen depends on the requirement of the possible application. The production processes allovv the manufacture of semi-finished prod-ucts or near net shape parts. 8 Literature 'K. U. Kainer (ed.): Metallische Verbundwerksoffe, DGM Informa-tionsgesellschaft. Oberursel, 1994 2 DURALCAN Composites for Gravity Castings, Duralcan USA, San Diego, 1992 3 DURALCAN Composites for High-Pressure Die Castings, Duralcan USA, San Diego, 1992 ""K. U. Kainer: Guss Produkte 91, Verlag Hoppenstedt, Darmstadt, 1991,261-262 5 C. W. Brovvn, W. Harrigan, J. F. Dolowy, Proc. Verbundvverk 90, De-mat, Frankfurt, 1990, 20.1. - 20-15 6Manufactures of Discontinuously Reinforced Aluminium (DRA), DWA Composite Specialities, Inc., Chatsvvorth USA, 1995 7W. Henning, E. Kohler, Maschinenmarkt, 101, 1995, 50-55 8 S. Mielke, N. Seitz, Grosche, Int. Conf.on Metal Matrix Composites, The Institute of Metals, London, 1987, 4/1-4/3 '' K. U. Kainer, Keramische Partikel, Fasern und Kurzfasern fiir eine Verstarkung von metallischen Werkstoffen in Metallische Ver-bundwerkstoffe, K. U. Kainer (ed.), DGM Informationsgesellschaft, Oberursel, 1994, 43-64 10 H. Hegeler, R. Buschmann, I. Elstner: Herstellung, Eigenschaften und Anvvendungen von Kurz- und Langfaserpreforms in Metallische Ver-bundwerkstoffe, K. U. Kainer (ed.), DGM Informationsgesellschaft, Oberursel, 1994, 101-116 1' Lanxide Electronic Components, Lanxide Electronic Components, Inc., Nevvark USA, 1995 12 C. Fritze, K. U. Kainer: Proc. Conf. Verbundwerkstoffe und Werkstoff-verbunde, G. Ziegler (ed.) DGM Informationsgesellschaft, Oberursel, 1996, 483-486 13 A. G. Leatham, A. Ogilvy, L. Elias, Proc. Int. Conf. P/M in Aero-space, Defence and Demanding Applications, MPIF, Princeton, USA, 1993, 165-175 14 Cospray Ltd. Banbury, U.K., 1992 ,5Keramal Aluminium-Verbundvverkstoffe, Aluminium Ranshofen GmbH, Ranshofen, Austria, 1992 16 K. U. Kainer, B. L. Mordike, Metali, 44, 1992, 436-439 17 K. U. Kainer, Proc. Int. Conf. New and Alternative Materials for the Transportation Industries, I SATA, Croydon, 1994, 463-470 18 K. U. Kainer: Konzepte zum Recycling von Metallmatrix- Ver-bundvverkstoffen, in press Proc. Recycling von Verbundvverkstoffen und Werkstoffverbunden, DGM Informationsgesellschaft, Frankfurt Microstructural Considerations Limiting the Mechanical Properties of HSLA Steel Mikrostrukturne omejitve mehanskih lastnosti HSLA jekel L. Parilak1, IMR SAS, Košice, Slovakia Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The influence of chemical composition, grain and subgrain size, and precipitation on yield strength, transition temperature and the work strengthening exponent was analyzed for HSLA (high strength, low alioy) steel. The relationships are quantified and transferred to graphic charts - nomogram for steel with polygonal as well as non polygonal microstructure. The limits of mechanical properties (the highest combinations of yield strength and transition temperature) were quantified for the polygonal HSLA microstructure. Key words: microstructure, mechanical properties, HSLA steels Vpliv kemijske sestave, velikosti zrn in podzrn ter izločanja na mejo plastičnosti, prehodno temperaturo žilavosti in keficienta deformacijske utrditve je bil analiziran za HSLA jekla. Odvisnosti so kvantificirane in zapisane v grafih - nomogramih za jekla s poligonalno in acirkularno mikrostrukturo. Mejne mehanske lastnosti (kombinacije največje meje plastičnosti in prehodne temperature žilavosti) so bile kvantificirane za poligonalno mikrostrukturo. Ključne besede: HSLA jekla, mikrostrukture, mehanske lastnosti 1 Introduction At the development of nevv steel types the key problem is to understand the influence of chemical composition and obtainable parameters of microstructure on strength. plasticity and brittle fracture resistance. It is es-sential to obtain the quantitative description of the rela-tion, and the description should be based on the knowl-edge concerning the nature of the mechanical properties in question. This way a valuable information can be obtained for the production technology, first for the prime chemical composition and heat treatment. In the presented work descriptions of correlations between chemical composition and parameters of microstructure on one side, and yield strength, work strengthening exponent, and transition temperature on the other, are compiled. They are quantified enabling direct application in engi-neering. In the second part of the work the limits of mechanical properties - combinations of strength, plasticity and brittle fracture resistance, are shown for the polygo-nal microstructure. 2 Microstructural essence of mechanical and fracture properties of microalloyed steels Investigated were lovv-carbon microalloyed steels based on Ti, V, Nb, with eventual addition of Mo, in po-lygonal and non-polygonal microstructures. Introductory studies were devoted to the kinetics of precipitation of carbides, nitrides or carbonitrides of microalloying elements from the viewpoint of its intensity and effective-ness. Furthermore, investigated were also questions of ' Ass. Prof. Dr. L udovii PARILAK Institute of Materials Research SAS Watsonova 47. 04353 Košice. Slovakia laws of interphase precipitation and precipitation in austenite and ferrite. The main objeetive was to gain the knovvledge of laws of the effect of precipitation states on strength as well as, plastic and brittle fracture properties. Analyses were carried out on several hundreds of struc-tural states in the state after rolling at hot rolling mili in VSZ JSC Košice, or in the state after thermal processing. Main attention was paid to the yield point, work strengthening exponent and transition temperature of noteh toughness. 2.1 Yield point The analyses were based on the assumption of an ad-ditive character of individual strengthening contributions to the yield point Re and the follovving relationship was proposed for the studies set of steels: Rc = Rpn+Rin+R0+Rsg+Rs+Rpr+Rp+Rd (1) where Rpn - is the contribution of lattice frietion stress; Rin - contribution to strengthening on account of interstitially dissolved atoms of additives; Rd - contribution of dislocation strengthening; Rg - strengthening contribution resulting from the size of grains; Rsg - contribution resulting from the effect of subgrains; Rpr -pearlitic contribution; Rs - substitution contribution; Rp -precipitation contribution. Their quantitative expression is based on relations comprised in Analyses provided a quantitative expression of the substitution effect of manganese RMn and confirmation of the effect of silicon and pearlite on strengthening contributions (Rsi, Rpr). In addition to that a quantitative effect of polygonaI ferrite grains d, or formations delimited by large angle boundaries (dF) in non-polygonal microstructures was described. The quantitative expression of subgrain strengthening with intensity Rsg = Gb-dsG"1 = 0.1 dsG"1 (the size of a subgrain dsG in mm) in a very good agree-ment with the Landford-Cohen relation, was used. The interchangeability of Rg and Rd was demonstrated, with Rd representing a contribution of transformation or of "geometrically inevitable" dislocations. The analyses of the influence of precipitation on pre-cipitation strengthening, employing ali available theo-retical models, were carried out. These analyses resulted in a quantitative relation for precipitation strengthening: Rp = k^"2 (2) where A. is the average planary interparticle distance of precipitates. The physical interpretation of this relation is follovving: Precipitation strengtheningis inversely pro-portional to the mean size of a free sliding area, corresponding to one precipitate (obstacle) standing in the way of the moving dislocation. The strengthening inten-sity constant kpR acquires a force dimension and can represent a mean value of force interaction phenomena between dislocations and precipitations, leading to a critical stress for the passing of dislocations trough ob-stacles. Its value kpR = 76.8 ■ 10"8 N is of the order corresponding to the size of an interaction of an edge dislocation vvith an elastic field of a particle F = 10"7 N). The quantitative behaviour of a thermaly dependent constituent of the yield point (R*) was determined in the range -196 to +20°C, together vvith parameters Ct, B, ap-pearing in the relation: R* = C, ■ exp(-T/B) (3) yield point B (relation (3)). The surface-plastic energy y, shear modulus of elasticity G, parameters ky, kf and the mode of stressing q are connected to the values A and B in the relation A = Bln(^--kf) (6) ATj is the shift of transition temperature and depends from the struetural parameter t or eventually from the chemical composition. The positive effect of grain refining on an improve-ment of brittle fracture resistance has been demonstrated and a direct relationship of its intensity and a thermal change of the yield point has been observed. A good agreement of the parameter B in relations (3) and (5) was deteeted. An embrittlement effect of pearlite and silicon has been demonstrated. The influence of precipitation on the shift of transition temperature vvas demonstrated to follovv the relation ATp = kpT-X"2 (7) An estimate of the barrier effect of grain boundaries against propagation of cleavage cracks (k = 55 Nmm~3/2) vvas provided together vvith a value of surface plastic en-ergy at T k (y = 10"2 Nmm"1)- In case of polygonal micro-struetures analyses did not exclude a positive effect of manganese on the improvement of brittle fracture resistance and the quantitative expression corresponded to results of Pickering. In case of non-polygonal microstruc-tures an absence of significant effect of subgrains on transition temperature changes vvas observed. 2.2 Transition temperature Our analyses vvere based on CottrelFs energetic bal-ance of cohesion of tough/brittle transition and Petch's condition of equality of the yield point Re and fracture stress Rfr for determination of the transition temperature of brittleness Tk. Contrary to Petch's formulation, we have assumed a general interaction betvveen individual parameters of microstructure and chemical composition and the frietion stress Rofr, appearing in the relation for fracture stress Rpk = R0FR+kf ■ d"1'2 (4) vvhich resulted in the development of a corresponding model and analytical formulation. The kf parameter rep-resents a barrier effect of grain boundaries direeted against the propagation of cracks aeross boundaries of grains. The performed analyses provided the follovving relation for the transition temperature TK05) = A-B-ln(d-U2) + 5;AT1 (5) (i) vvhere A is the so-called threshold value of brittleness, dependent on the intensity of the thermal change of 2.3 Complex relations In our previous vvorks1,2,3 vve presented simplified relations for the evaluation of the influence of microstructure on yield strength Re, transition temperature T35 and vvork strengthening exponent n. For a polygonal microstructure it is expressed as: Re = R0 + Rm„ + AR (8) T35 = A - B ■ ln(d-"2) + C • AR (9) n = a + — (10) AR vvhere Rg = 15 • d~1/2 is the strengthening by ferrite grain size d (mm); RMn = 50 ■ XMn is the strengthening share of manganese XMn (%); AR is the part of embrittlement caused by strengthening, for microalloyed steel in-cluding mainly precipitation strengthening Rp, and also the influence of strengthening by silicon content Rsi, pearlite content Rpr, Peierls-Nabarro stress Rpn, and by interstitial strengthening Rin (AR = Rp + Rsi + Rpr + Rpn + Rin); A = 147°C, B = 110°C, C = 0.4°C/MPa is an embrittlement constant, a, b are regression coeffi-cients. For non polygonal microstructure similar relation vvere derived: Re = R0 + RSG + RMn +AR (11) T35 = A - B • ln(d-"2) + C • AR (12) where Rg = 19 ■ d"1/2; A = 143°C; B = 100°C, C = 0.4 °C/MPa while Rsg = 0.1 ■ dsc"1 is the strengthening contribution of the subgrain size dsG (mm). The graphic interpretation of the relations is shown in Fig.l for the polygonal microstructure (eq. 8-10) and in Fig.2 for the non-polygonal one (eq. 11-12). It is important to note that the yield strength is con-trolled by a set of strengthening contributions with different influences on the brittle fracture resistance. The embrittlement from strengthening AR is resulting for every 100 MPa of strengthening a 40°C shift of the transition temperature into the wrong direction, causing the worsening of plastic properties, too, as shown by the work strengthening exponent. There is an influence of manganese content and subgrain size on strengthening too, though their influence on the transition temperature is not significant. Practically only one microstructural parameter is known, the increasing of the yield point to- gether vvith the increase of brittle fracture resistance. It is the ferrite grain size d, or described more generally the size of the microstructural object limited by large angled borders. It is of prime importance to constitute the chemical composition and microstructure in the way to obtain first this microstructural parameter in the quality reflecting the desired complex of properties. The relations given in the work are simplifted theoretical descrip-tions with coefficients calculated by regression analysis made on more than 300 microstructure types of steel produced in ironvvorks VSŽ, a.s. Košice, Slovakia. 3 Limits of polygonal microstructures We decided to define the limits of the complex of mechanical properties for a steel vvith polygonal microstructure. With this aim the HSLA steel, vvith yield strength from 420 to 700 MPa vvere evaluated. The basic features of the evaluation are shovvn in the graphic chart in Fig.3, vvhich was calculated for a 1% Mn content. The straight lines are representing the yield strength Re. The nomo- Re (MPa 1 Figure 1: A eomplex nomogram for relation betvveen microstructural parameters and mechanical properties of HSLA steels vvith polygonal microstructure Figure 2: A complex nomogram for relation betvveen microstructural parameters and mechanical properties of HSLA steels vvith non-polygonal microstructure Figure 3: Microstructural considerations limiting mechanical properties of HSLA steels with polygonal microstructure gram shovvs the possible combination of embrittlement AR and ferrite grain size d, necessary to obtain the selected yield point. The transition temperature T35 and the work strengthening exponent n are shown also. In Tab. 1 the combinations of embrittlement AR and ferrite grain size d in grades according to ASTM are shown, vvhich are necessary for a steel vvith the desired combination of yield strength Re and transition temperature T35. Table 1: Required ferrite grain size d and embrittlement by strengthening AR necessary for the combination of properties Rc and T35 T35(°C) 0 -20 -40 -60 Re d AR d AR d AR d AR (MPa) (MPa) (MPa) (MPa) (MPa) 420 10 230 11-10 220 11 190 11-12 170 490 11-10 275 11 260 12-11 230 12-13 210 560 11-12 320 12 290 12-13 270 13 250 630 12-13 370 13 340 13-14 320 14 290 700 13-14 420 14 400 14 350 ? ? In ali cases a fine ferrite grain is required. Knovving the manufacturing technology and the limits of the vvide strips hot rolling mili the production of steel vvith ferrite grain size under grade 14 cannot be experted. To obtain the grade 13 is very difficult, grade 12 is demanding, vvhile the more coarse grains are currently obtained. Consequently, Tab. 1 vvas simplified to Tab. 2 vvhich shovv that the elaboration of polygonal steel vvith the yield strength Re = 700 MPa and the transition temperature T35 under -40°C, is not be reliable. It is also not re-alistic to desire expert a limit of elasticity Re = 630 MPa vvith the transition temperature T35 better than -60°C. Table 2: Limits of the polygonal microstructure for different combinations of Re and T35 Rc T35 (°C) (MPa)_0_-20_;40_-60 420 1 1 1 1 490 1 1 2 2 560 1 2 2 3 630 2 3 3 4 700 3 _3_4_4 The possibilities are denoted: 1 - realistic, 2 - demandig, 3 - very difficult, 4 - fiction. In Fig. 3 it can be also seen, that for the mentioned Re and T35 values the ductility is very lovv, the vvork strengthening exponent in the range 0.10 to 0.16 (for the lovver strength) because for high Re values the embrittlement by AR is necessarily high, degrading the ductility and brittle fracture resistance. 4 Conclusion Starting from theoretical relations the influence of chemical composition and parameters of the microstructure on strength, transition temperature and vvork strengthening exponent vvere investigated. The results are compiled and the limit combinations of strength, plastic properties and resistance to brittle fracture for HSLA steel vvith polygonal microstructure are calculat 5 Acknovvledgment The vvork is supported by Project No. 2/1106/96 of the Slovak Scientific Grant Agency - VEGA.ed. 6 References 'L'. Parilak, M. Šlesar. B. Štefan: Structural Prediction of Mechanical Properties of HSLA Steels. In.: Proc. of Microalloying 88, ASMI, USA, 198B, 559 2B. Štefan: Fyzikalna metalurgia a vyvoj konštrukčnych zvaritetl'nych oceli. Doktorska dizertačnd praca, UEM SAV Košice, 1990 3L'. Parilak: Štrukturna podstata mechanickych a lomovych vlastnosti materialov. In.: Predikce mechanickych vlastnosti' kovovych materialu na zaklade strukturnlch charakteristik. l.dil. Nove Mesto na Morave, 11.-14.5.1993 Brno, P MSVTS VU 070 1993, 125 Predicting of Reactions During Carburization and Decarburization of Steels in Controlled Atmospheres Napovedovanje reakcij, ki potekajo med naogljičenjem in razogljičenjem jekla v kontroliranih atmosferah B. Koroušič1, IMT Ljubljana, Slovenija M. Stupnišek, Faculty of Mechanical Engineering and Naval Architecture, Zagreb Uni-versity, Croatia Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The knovvledge of the thermodynamics of complexe systems consisting of gases and metal should be valuable for the control of industrlal processes. The Gibbs energy minimization model has been implemented in the softvvare program GPftO® and associated with a powerfull and reliable database. The computer package can perform computation of the equilibrium composition in very complex chemical and metallurgical systems. Some examples in this paper illustrate the simplicity of the computation and the use of the program m the field of some typical metallurgical applications. Key words: equilibrium reactions, NOx modelling, combustion of fossil fuels, active gas-atmospheres, decarburizing of non-oriented electrical steels, carburizing of alloyed steels with in situ produced atmospheres Poznavanje termodinamičnih odnosov v kompleksnih sistemih plin - kovina ima lahko izreden oomen za kontrolo industrijskih procesov. Gibbsov model o minimizaciji energije je implementiran v programsko opremo GPRCP, ki mu služi kot osnova močna baza verificiranih termodinamičnih podatkov. Programska oprema omogoča izračunavanja ravnotežnih sestav v zelo kompleksnih kemijskih in metalurških sistemih. Navedeni primeri v tem članku ilustrirajo enostavnost izračunavanj in način uporabe programa na področju metalurških reakcij, ki jih večinoma izvajajo strokovnjaki na tem področju. Ključne besede: ravnotežne reakcije, tvorba NOx, zgorevanje fosilnih goriv, aktivne plinske atmosfere, razogljičenje neorientirane elektropločevine, naogljičenje legiranih jekel 1 Introduction The application of thermodynamics to a system gas/solid enables to calculate the composition at equilib-rium, direction and extent of change vvhich can take plače under specified conditions. Rapid developments have taken plače in recent years in efficiency of thermodynamics in the engineering as thermodynamic can be defined as being the meeting point between physical - chemical principles and practi-cal applications1. In this paper an attempt has been made to demonstrate use of a personal computer softvvare program as an ellegant and sensitive method for numerous metallurgical applications especially for the analysis of gaseous systems. It is hoped, that users of this method will be in a good position to go more deeply into learn-ing thermodynamic laws. 2 Principles of the Gibbs method In the fields of heat treatment of metals like anneall-ing, carburizing, decarburizing, nitrocarburizing and many other operations, the metallurgist is concerned not with the pure gases but with the mixture of various spe-cies (gaseous and solids) vvhich form the atmosphere in the furnace. Prof. Dr. Blaženko KOROUŠIČ Inštitut za kovinske materiale in tehnologije 1000 Ljubljana. Lepi pot 11. Slovenija The thermodynamics of sueh complex systems can be treated by two methods: - The classical method of numerical solution of an equilibrium problem when the equilibrium constant (Kr) or free energy change AG° of the involved reactions are known. - The general Gibbs method for the numerical solution of an equilibrium. The problem is to determine the values of the species vvhich minimize the state of total free energy at the given temperature and pressure. Both treatments are thermodynamically equivalent, hovvever, it seems that the later method has significant advantages for calculating the equilibrium conditions in complex systems, in mixtures containing both gaseous and condensed species. During the last 20 years, SOLGASMIX computer program, as the method of attacking chemical and metallurgical problems, has influenced our approach to the study of a braneh of scientific knovvledge in physical chemistry. There can be no doubt that to attack sueh a complex application of thermodynamics is only possible vvith the use of computer technology. 3 Description of the method used for the calculation of complex equilibrium conditions Several excellent softvvare programs for calculating equilibria reactions at high temperatures, have been de-veloped in the last two decades (SOLGASMIX, THER-MOCALC, FACT, CHEMSAGE...)23. However, most of them are designed and vvritten in a complex form using very strong computer units, while few are intented as a simply a tool to be applied for the purposes of solving practical problems. Therefore, it seemed worthwile to develop a program which would combine these two computer program designs. The new software program, called GPRO is based on the method of free energy mini-mization and extended to systems containing numerous gaseous and condensed phases in accordance with SOL-GASMIX-principles. GPRO-program is dimensioned for 16 elements and 100 species. If necessary, this figure can be increased or new included datasets, vvhich if neces-sary are vvritten by the user (private databases are open and can be easily included also). 3.1 Thermodynamical approaches to the Gibbs-method The povver of Gibbs method energy minimization lies in its simplicity for the description of chemical reactions in complex systems, and its ability to facilitate the determination of the effect, on equilibrium state, of changes in the external influences vvhich can be brought to bear on the system. In our softvvare program, the user needs only to specify the type, the species present and the conditions (for example: temperature of the system) for the calculation. The program vvill automatically perform equilibrium thermodynamic computions typically associ-ated vvith complex chemical equilibria from a defined database. With the aid of the GPRO-program, a user is able to perform most of the follovving operations: 1. The energy for pre-heating the initial mixture from the initial temperature T0 to the reaction temperature T, 2. The reaction heat, 3. The computation of the complex chemical equili-bria in gaseous mixtures and activity of solid compounds, 4. Displaying and printing data for compounds and solutions at selected temperature and composition. An additional scientific and engineering benefit of this program is the softvvare able to develop a more basic understanding of chemical equilibria at high temperatures and its applications. Although the povver softvvare program vvill automatically perform the thermodynami-cal computation (no danger of pluging vvrong numbers in vvrong equations), hovvever, the user must have some knovvledge of the chemical nature of the considered sys-tem. In this paper the attempt is made to demonstrate the breadth and diversity of the modern softvvare program in simple way so that a user may be able to understand the thermodynamical method and apply it to metallurgical problems. Most of the examples are chosen vvith the aim to shovv superiority of the computer program, over tradi-tionally manual methods, vvhich are particularly stressed for the engineers and students. 3.2 Databases associated for the equilibrium thermody- namic comptutations From many excellent standard treatises on thermody-namics it is knovvn, that vvithout reliable thermodynamic data most equations are ineffective and the numerical an-svvers vvill be therefore vvrong. GPRO softvvare program is based on the use of both the expressions for calcula-tions of the standard Gibbs energies of the formation of a selected phase: in the form: AG°t = Y + B + CT + Dj2 + Ej3 + + FTlnT or using thermodynamical data on enthalpy AH°t, en-tropy AS°t and heat capacity CP(T): rT rTC (T) AG°t = AH°298 + J Cp(T)dt -TAS°298 -Tj -^dT t t 'o o Both methods used from the database involve the search for a minimum value of free energy AG of a sys-tem and give an equivalent result. Hovvever, the last method using enthalpy AH°t, entropy AS°t and heat ca-pacity CP(T) has more advantages because it combines heat and equilibrium calculations. A typical example is the determination of the adiabatic flame temperature, vvhere enthalpy of reaction serves as the criterion of the heat balance. 4 Exploiting the GPRO-program for complex equili-bria calculations Modelling Mechanism of Formation Nitrogenous Oxides by the Combustion of Fossil Fuels Modern combustion processes of fossil fuels meet the relevant requirements for cost-effective operation and avoidance of enviromental pollution. In article some results of the basic study of the formation and reduetion NOx in high temperature combustion processes are pre-sented. The obtained results demonstrate the use of the sophisticated methods of thermodynamics as one of the most important tools by the study of the combustion processes for a better understanding of the mechanism of formation of nitrogen oxides, one of the most important pollutants in combustion of fossil fuels8"16. Example 1: In this example is a demonstration of the use of the GPRO-softvvare program as method for predietion of complex combustion reactions and equilibrium gas composition including NO-oxydes formation. The high temperature furnace is fired vvith natural gas and air (no air preheating). The question vvas: calculate Table 1: Results of GPRO-analysis of the natural gas combustion by different air - index and without air preheating Equilibrium data for methane combustion (A,=0,74...2,2)' CH4 + 2^.0 2 + 7,52XN 2 Air index 0.74 0.84 0,94 1,00 1,10 1,30 1,60 1.80 2,00 2.20 CO (%) 6,90 4,44 1.95 0,95 0,30 0,04 0.002 0,0 0,0 0,0 CO2 (%) 4.77 6,29 7.69 8,49 8,40 7,44 6,16 5,51 4,99 4,71 NO (%) 0,0021 0,017 0,112 0,23 0,345 0,334 0,200 0,130 0,087 0,066 H20 (%) 18,09 18,99 19,01 18,52 17,28 14,94 12,32 11,03 9,98 9,42 O (ppm) 1,04 14,20 133 258 282 111 15 5 1 0 02 m 0,00 0,00 0,127 0,53 1,75 4,32 7,28 8,75 9,93 10,56 N2 (%) 64,97 67,79 70,02 70,89 71,78 72,90 74,03 74,58 75,02 75,25 H, (%) 5.27 2,47 0.81 0.37 0,11 0,02 0,00 0,00 0,00 0,00 SfinpuD (mole) 8.04 9,00 9,25 10,52 11,47 13,37 16,23 18,14 20,04 21,24 If«.»ii (mole) 8,56 9,32 10,10 10,59 11,49 13,37 16,23 18,14 20,04 21,24 TadbiK)_2023 2143 2233 2231 2151 1955 1712 1584 1482 1421 1 air index the equilibrium gas composition and the adiabatic flame temperature for the air-index in range 0,74 < X < 2,2 and compare the obtained results of the flame temperature vvith similar reference data knovvn in the literature (nor-mally presented in graphically form). In table 1 and figure 1 the computed values for the gas equilibrium are given. The adiabatic flame temperature calculation shovv values slightly above the compared data. Thermodynamic evaluation of carburizing atmospheres The accuracy of the gaseous atmosphere control in the steel carburizing furnaces has been remarkably im-proved ovving to the application of the computer control system and the development of nevv measuring tech- niques, for example: oxygen and/or carbon sensors. The atmosphere in carburizing furnaces are consists of: air + methane or other hydrocarbons and involves the gases CO, CO2, H2, H2O, N2. The four first gases are interde-pendent in a reversible reaction, commonly called the vvater-gas reaction: CO + H20 = C02 + H2 The ratio: K,„ — Pco; • Ph, Pco ' Ph,o (D is a constant, the value of vvhich depend on the temperature. The carburizing of steel, i.e. carbon content increas-ing on the steel surface, occurs through the reaction: CH4+2n02+7.52nN2 TTiermal NOx . dNO/di (t-o) (»l %/•) air inde* ■ . ............. i 105 CH4+2n02*7.62nN2 Th*rm«l HCM IdNO/dtKt-O) IvdVil air hxtax ■ 105, .............i. / ........ ...........f........ T .... i ....... Q»yo«nWi + (volH) CH4*2n02«7.52nN2 Thflrm«! NOx CH4*2n02*7.62nN2 Thtrmal NO* IdNO/dlKl-O) (votV«) Figure 1: The formation of the nitro-genous oxide NO (model simulation) Slika 1: Tvorba dušičnega oksida NO (modelne simulacije) .dNO/dt (t-O).(vcl v«) no aoo Alf tompTahira C air index ■ 2CO = C + CO, (2) For any given temperature, the coresponding equilib-rium constant of Boudouard's reaction will determine the carburizing potential of the atmosphere: K: Pco ■ (Pcc/Pc (3) The carbon potential of on atmosphere is simple to determine if the partial pressure of CO and ratio (pco/pco:) is known. From EMF measurement (electro-motive force) with the oxygen probe, considering the furnace temperature (pco/pco2) or (ph2/Ph2o) and measuring the CO - content in the atmosphere is possible to en-sure the control of the carburizing process. Example 2: The carbon activity in a steel depends on the content of alloying elements, thus every steel composition will have determined carbon potential which corresponding to the atmosphere composition. In next example three type steels were treated vvith air + methane atmosphere vvith the aim to obtain a constant carbon content near the surface of about 1 wt.% C). Data in table 2 shovv, that small deviations in the gas atmospheres (or the change of air + methane ratio) have a remarkable effect on the carbon activity. This model simulation is in good agreement vvith practical data. Table 2: Influence of the steel chemistry on the process parameters (Simulation made by GPRO programme by T = 1223K) Chemistry (%) Fe+l%C+ l%Si Fe+ 1%C Fe+l%C+ l%Cr CO 19.27 19.25 19.24 co2 0.0629 0.0717 0.078 h2o 0.199 0.227 0.248 ch4 5.17 5.15 5.13 n2 36.83 36.88 36.92 h2 38.46 38.42 38.38 02 (bar) 9.4 10"21 1.16 lO'20 1.46 lO"2« EMF(mV) 1173 1167 1162 Tdp*'(°C) -13 -11 -10 ac 0.818 0.715 0.654 %c 1.03 1.06 1.08 Qgas(m3/h) 1.1 1.1 1.1 Oair(m3/h) 2.094 2.100 2.1045 > Tdp = Dew point temperature Example 3: The carburizing of steel is a continuously process vvithin vvhich - due to the kinetics of various reactions -damming up effects may occur leading to non equilib-rium CH4-contents in the furnace atmosphere. In this čase the carburizing reactions under non-equi-librium conditions are modelled. A mixture of natural gas and air at 1 bar total pressure is introduced into the carburizing furnace heated to 1223 K5. The quantity of natural gas and air are 3,1.10"4 m3/s and 4,2.10"4 m3/s. Calculate the gaseous equilib-rium composition in the furnace atmosphere and the carbon activity assuming graphite as standard state. If the air flovv suddenly increased from 4,2.10~4 m3/s to 5,8.10"4 m3/s by constant natural gas flovv 3,1.10"4 m3/s in the in-let mixture, determine the new gas equilibrium composition and carbon activity! Table 3 shovvs the computed results for gas non-equi-librium composition and obtained energy changes, the preheating energy H°t - H°289, the heat of reaction H°r and H°totai the total heat of the system. Table 3: Example of input and output of a non-equilibrium composition by the production of endothermic gas from methane and air by t = 1223 K _1,1CH4 + 2A.02 + 7,52XN2_ X = 0,157 X = 0,221 Air/ 4,2.10(-4)/3,1.10(-4' 5,8.10('4V3,1.10("4) methane (m3/s) ot1 'melhane = 0,575 aVthane = 0,795 Xinn (mole) Xoul voI.(%) Xinp (mole) X0Ui vol.(%) CO 0,0000 17,58 0,0000 19,25 C02 0.0000 4,28.10"2 0,0000 7,19.10"2 CH4 1,1000 13,40 1,1000 5,80 h2o 0,0000 0,135 0,0000 0,227 HCN 0,0000 3,76. lO"3 0,0000 2,65.10'3 h2 0,0000 35,29 0,0000 38,41 n2 1,1890 33,33 1,6590 36,22 o2 0.3150 (5,45.10"21)2' 0,4410 (1.28.10"20)2' a.-3' 1,0000 0,714 X(mole) 2,604 3,5397 3,200 4,4962 (H2/C)4) 2.00 2,00 2,00 2,00 H°t-H°298 (kJ/mol) 92,95 116,77 H°r (kJ/mol) - 14,79 - 22,24 H°tolyl (kJ/mol) 78,15 94,53 CH4(c) " ^methane = CH4(c) = fully cracked CH4 and CH4(t) CH4(t) = total CH4 2) p02 in bar, 3) ac = carbon activity referred to graphite as standard state 4. H2 H, + H,O + 2CH4 ' — =-, A. = air index c co + co2 + ch4 Calculation of the decarbirization process of silicon al-loyed steels The use of gaseous atmospheres vvith a vvell-defined oxygen potential for the decarburization of low carbon iron-silicon steels in continuous furnaces can be simu-lated using a thermodynamical model. Equilibrium cal-culations and practical measurements shovv that the solu-bility and carbon activity in Fe-C-Si steels depend on the gaseous atmosphere, temperature and steel composition. Silicon-iron alloys containing 1 - 3% Si and 0,3 - 1% Al are typical steels for non-oriented sheets and a strict Fe + C + 3 % Si Dynamo steel 0.03%C, 2.7%Si, T=840 C, 6t/h TdpCC) EMF(mV) (H20/H2)_eq n95tor09 Figure 2: Plot of thermodynamical data for Fe2SiO-i as a function temperature calculated with GPRO-program Slika 2: Diagram termodinamihnih podatkov Fe2SiC>4 kot funkcija temperature izrahunano s GPRO-programom control of the decarburization and surface reactions is re-quired. The optimum properties for an electrical steel normally include high permeability with low core loss and minimal aging effects. An important factor in a process control is the formation of a high quality glassy film which is developed through a complex series of processing steps. Example 4: In order to clarify the relation between the decarburization atmosphere for the carbon removal during the an-nealing, the thermodynamical reactions and formation of different oxide phases in the scale have been studied. The first task was the determination of the carbon activ-ity in the decarburization gas atmosphere containing at the start H2 + N2 + H2O in temperature range 600 -1000°C. The mathematical model GPRO allows an easy use of thermodynamical data to predict the equilibrium carbon content in electrical steels. It is convenient to use the carbon activity in the gas atmosphere by different partial pressures of CO to present the conditions for the formation of FeO and Fe2SiC>4 by the different temperature. Figure 2 shows the results obtained. Having these curves available, it is possible to determine the dew point temperature as the function of the partial pressure ratio H2O/H2. Figure 3: Equilibrium oxide-formation by the decarburization of non-oriented electrical sheets in gaseous atmosphere H2 + H2O + N2 (Tdp - dew point temperature, EMF (mV) = electromotive force, (H20/H2)-eq = equilibrium pressure ratio) Slika 3: Ravnotežni pogoji tvorbe oksidov med razogljičenjem neorintitane pločevine v plinski atmosferi H2 + H2O + N2 (Tdp - točka rosišča, EMF (mV) = elektromotorna napetost, (H20/H2)-eq = ravnotežno razmerje plinov) As shown on Figure 3, the ratio of H2O/H2 at which the formation of fayalite actually disappeared is near H2O/H2 = 0,24 at 840°C. It is obvious that the pressure ratio H2O/H2 and CO2/CO is interchangeable with the partial pressure of oxygen - po3 and finally also by means of the relation: log po, = log A Pco ac v / 11854 - 9,090 (4) which allows the application of the oxygen (carbon) sensor signal (EMF). 5 Conclusions The use of thermodynamic predictive model offers many advantages over conventional gas atmosphere cal-culations because of the simplicity for description of chemical reactions in complex systems, the automatic performance of equilibrium computations, of the avoid-ance plugging wrong numbers in wrong equations and so on. The rational and theoretical basis for the Gibbs en-ergy model used was presented elsewhere5'17-19. To sum-marise, the key features of model calculations for the ni-trocarburizing atmospheres are as follows: - Modern combustion processes of fossil fuels meet strict requirements for cost-effective operation and avoidance of enviromental pollution. This article presents the first results of study into formation and reduction NOx in high temperature combustion processes. - The obtained results demonstrate the use of the so-phisticated methods of thermodynamics as one of the most important tools for the study of combustion processes to understanding better the mechanism of formation of nitrogen oxides, as one of the most important pollutants in fossil fuels combustion. - Little is given in disponible references on use of thermodynamical models in the field of active at-mospheres. Such mixtures containing both gaseous and condensed components for example: Fe + C + O + H + N are extremly complicated for the numerical calculations. Detailed experimental studies are diffi-cult and also thermodynamical results are mostly presented in the graphical form, which are very use-ful in research work but of little effectivness in searching solutions for a current practical operation. - To obtain equilibrium compositions in the real gaseous mixtures by high temperatures, taking into ac-count both energy and material balances, the development of new approaches are strongly required. 6 References 'D. R. Gaskel: Introduction to Metallurgical Thermodynamic McGravv-Hill Book Company, (Washington, D.C.), 1973 2 W. B. Christopher, G. Erriksson: Metallurgical Thermochemical Data-bases, Canadian Metallurg. Quat., 29, 1990, 2, 105-132 3 G. Eriksson, K. Hack: ChemSage - A Computer Program for the Cal-culation of Coinplex Chemical Equilibria, Metallurgical Transactions B, 21B, 1990, 1013-1023 4 F. D. Richardson, J. H. E. Jeffes: J. Iron Steel Inst., 160, 1948. 261-270 5B. Koroušic: Contributions to the Computer Predictions of the Homogene and Heterogene Equilibrium Compositions for Gaseous Atmos-pheres, Veitsch-Radex-Rundschau, 1994, 1-2, 465-542 6R. Hoffmann: Aspekte des Kurzzeitnitrierens, HTM, 31, 1976, 152-157 7H. Henrich, W. Kozlowski, W. Liere: Energieeinsparung und Schad-stoffreduzierung an Beheizungseinrichtungen in der Industrie, Gaswarme International, 41, 1992, 2/3, 89-100 8 F. A. \Villiams: Combustion Theory, Benjamin Cummings, Menlo Park, 1985 'K. Hein: Fossil Fuel Utilisation, Combust. Sci. and Tech., 93, 1993, 61, 27-39 10 R. Haupt, R. Oppenber: Feuerungen in Dampf - und Heisswasserer-zeugern, Gaswarme International, 41, 1992, 10, 445-456 "B. Bonn. H. Baumann: Kenntnisstand der N2O Bildung in ver-schiedenen Feuerungsanlagen, VDI Berichte, 922, 1991, 17, 625-633 12 H. Schuster: Minderung der NO* - Emissionen aus Kraftvverksfeuerungen, VDI-Kolloquium Emissionsminderung bei Feuerungsanlagen SO2 - NOx - Staub Essen, 10-11 November 1983, VDI-Bericht, Nr. 495 13 Y. H. Song, J. M. Beer, A.F. Sarofim: Reduction of Nitric Oxide by Coal Char at Temperatures of 1250 - 1750 K, Comb. Sci. Techn., 25, 1981,237 14 H. Schulz, H. Kremer: Bildung von Stickstoffoxiden bei der Kohlen-staubverbrennung, Brennstoff Warme Kraft BWK, 37, 1985, 1/2, 29-35 15 J. Zel'dovich: The oxidation of nitrogen in combustion and explo-sions, Acta Physicochimica URSS, 21, 1946, 4, 577-628 16 Y. B. Zefdovidh, P. Y. Sadovnik, D. A. Frank-Kamentku: Oxidation of Nitrogen in Combustion, (translated by Shelef), Academy of Sciences ofUSSR, 1947 17 B. Koroušic: Fundamental thermodynamic aspects of the Ca0-Ah03-S1O2 system, Steel res., 62, 1991, 7, 285-289 18 B. Koroušic: Study of equilibrium reactions in gaseous nixtures (Part 1. Protective atmospheres), Rudarsko-Metalurški zbornik, 1993, 1-2, 5-17 " B Koroušic, M. Stupnišek: A thermodynamic evaluation of nitrocar-burizing atmospheres, Steel res., 66, 1995, 8, 349-352 Fusion of Low Carbon Steel Scrap in the Middle Carbon Steel Melt Taljenje niskougljičnog čeličnog otpatka u talini srednje ugljičnog čelika V. Grozdanič,1 Metallurgical faculty Sisak, Croatia Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 A quasi three-dimensional mathematical model of fusion of cylindrical steel scrap In converter melt was developed. The model vvas solved using the implicit alternating direction method. The obtained algorithm vvas programmed in ASCII FORTRAN for the computer SPERRY 1100/72. In the model temperature dependent thermophysical properties of material vvere incorporated. That gives to the model a nonlinearity. On the basis of the model it vvas concluded that the addition of 1% of Iow carbon steel scrap decreases the temperature of a middle carbon steel melt for cca 20°C. This is in good agreement vvith experimental data from literature. The mathematical model was tested for one-dimensional exact solution using the Bessel functions. A good agreement was found. Key vvords: fusion, steel scrap, mathematical model Razvijen je i istražen kvazitrodimenzijski matematički model taljenja valjkastog čeličnog otpatka u talini kod konverterskog procesa. Matematički model riješen je implicitnom metodom promjenljivog smjera. Dobiveni algoritam programiran je u programskom jeziku ASCII FORTRAN za računalo SPERRY 1100/72. U matematički model inkorporirana su temperaturno ovisna toplofizička svojstva materijala, što modelu daje nelinearnost. Na temelju matematičkog modela zaključeno je da 1 % niskougljičnog čeličnog otpatka snizi temperaturu srednje ugljične celične taline za cca 20°C, što se dobro slaže s eksperimentalnim podacima iz literature. Matematički model testiran je na jednodimenzionalnom egzaktnom rješenju pomoču Besselovih funkcija. Konstatirano je njihovo medusobno dobro slaganje. Ključne riječi: taljenje, čelični otpadak, matematički model 1 Introduction From the increase of the share of metalic scrap in the charge, an increased economy of steel manufacturing in the converter processes is expected. The share of steel scrap should approach that in the openhearth process. Steel scrap, usually as low carbon steel refuse, is a very economical means of melt cooling. Iron ore is only better means of cooling. Data in ref.l show that 1% of steel scrap decreases the bath temperature for 12 to 15°C, vvhile 1% iron ore decreases the temperature for 30 to 40°C. Hovvever, production expirience shows that steel scrap is better than iron ore. For example, by using steel the quantity of metal ejected from the converter is dimin-ished, the resistance of refractory lining is increased and a better utilization of excess heat in the bath obtained. The regime of scrap fusion depends in significant degree on scrap size, and affects the bath temperature, the slag forming processes, as well as the oxidation of carbon and the metal desulfurization. For instance, small sized scrap melts faster and cools quickly the bath. This decreases the rate of slag forming, carbon oxidation, desulfurization, and lower the quantity of blovvn oxygen. Very large pieces of scrap don't melt completly during the process-•ng in the oxygen converter. A mathematical model of low carbon steel scrap melting in the carbon steel melt was developed and tested vvith the aim to determine the optimal scrap size. Since the bath temperature is above Dr. Sc. Vladimir GROZDANIČ Univcrsily of Zagreb Metallurgical facully Sisak. Aleja narodnih heroja 3. 40 Sisak. Crnalia 1550°C the so-called diffusion melting is not considered. In figure 1 the investigated system is illustrated. It con-sists of a volume element of melt in which cylindrical steel scrap is immersed. 2 Mathematical model For the start of the melting of a cylindrical scrap piece in a volume element of melt in figure 1 the Fourier's partial differential equation of heat conduction has the form2: dT dt = a V čPT J_3T čPT dr2 + r dr + dz1 (D Since for the horizontal axis of the system r — 0 the equation (1) is modified according to L'Hospital's in the form: Figure 1: Volume element of melt with cylindrieal steel scrap 3T 52T 32T (2) The basic assumption for the validity of the differen-tial equations (1) and (2) is that scrap piece is immersed is physicaly realistic since the difference in density of steel scrap (7860 kg/m3) and melt (7507 kg/m3) is very small. Considering the system in figure 1, it can be con-cluded that mathematical model is quasi three-dimen-sional. In the time t = 0 the temperature of the melt is Tl, and that of the scrap piece is Ts. The initial temperature at the steel scrap/melt boundary interface is obtained by solving the Fourier's differential equation for heat flow through the contact area of two semifinite medias3: Ti = T. + - tl-t, k„, a. (3) On the contact steel scrap/melt area a continuous heat flow occurs with boundary condition of the fourth kind: Km i ~~ Ks -n dn dn (4) In developing the model it was asummed that thermal properties of low carbon steel scrap (0,2% C) and middle carbon steel melt (0,6% C) are temperature dependent4. 3 Implicit alternating direction method The differential heat flow equations (1) and (2) with the corresponding initial and boundary conditions were numerically solved using the implicit alternating direction method5 and dividing the time interval into two steps. In the first half of the time interval the equation is solved implicitly for the z and explicitly for the r direction. The procedure is reversed in the second half of time interval. Consequently, for the differential equation (1) and first half of time interval At/2 we obtain: T11 _Tn i "T* J 2rAr z a,jn At/2 7 Whereas for the second At/2 we obtain pn+I _'pn+l - 4- i. + i-H-1 i.j-1 + 32 ^ = J_ hA L, 2rAr r ai„i.n At/2 (6) The numerical solution of the differential equation (2) of heat flow for first At/2 is: npn _nrn — + 3; T (Ar)2 and for second At/2 . -LlLzZIi z ,1_a,,„ At/2 Tn+1 "-pn+1 , --pn+1 hp* i,2 Ai.l , rr* 1 i.l i. (Ar)2 + d2rT',=- ai.j.n At/2 (8) The solution of equations (1) and (2) for the net point (i, j) in the melt and scrap piece as well as for net points on their boundary surface are given in Appendix 2. Generally it holds: bivi+civi = di a2Vl+b2V2+C2V3 = d2 a3V3+b3V3+C3V4 = d3 aivi-i+bivj-t-civi+i = di (9) aN-lVN-2+bN-lVN-l+CN-lVN = dN-1 aNVN-l+bNVN = dN where v is the unknown temperature, and N is a real number. On the base of the presented algorithm of fusion a computer program was written in ASCII FORTRAN and solved on SPERRY 1100/72 computer. 4 Discusion The simulation of fusion of a low carbon steel scrap in the carbon steel melt is carried out by space steps Az = Ar = 1 cm and the time step At = 30 s till tmax = 630 s. The initial melt temperatures of 1700°C for the melt, 25°C for the scrap piece and 883°C for the boundary surface were assumed. On the basis of succesive temperatures prints out for particular net points the fusion time of 540 s was obtained for a low carbon steel cylinder of size <|) 50 x 100 mm. The weight of the scrap piece was 4,3% of total vveight of the melt. Thus, it can be con-cluded that 1 % of steel scrap decrease the temperature of steel melt for 20 to 22°C, a value in good agreement with published experimental data1. Three-dimensional mathematical model of fusion of scrap piece was tested through the one-dimensional exact solution of the fusion of the low carbon steel cylinder and a good agreement was established. The derivation of the exact solution using Bessel functions is given in Appendix 3. 5 Conclusions A quasi three-dimensional mathematical model of fusion of low carbon steel scrap piece in carbon steel melt in oxygen converter was developed. The model was checked in the base of experience data. The simulation of the fusion is carried out on cylindrical piece of diameter of 50 mm and length of 100 mm. The fusion time of 540 s was calculated. Also it has been established that the ad-dition of 1% of steel scrap decrease the melt temperature for cca 20°C, a value in good agreement with experimen-tal data, and also with the exact one-dimensional solution of the equations, which are the model base. Appendix 1 Abbreviations used: a - temperature conductivity ai, bi, ci,di - coefficients adjoining to unknowns in tridiagonal system of algebric equations cp - specific heat at constant pressure k - thermal conductivity n - vertical direction r - space coordinate t - time T - temperature vi - unknown in system of simultaneous algebric equations z - space coordinate Appendix 2 Constant whivh appear in tridiagonal coefficients aAt pi =■ 2(Ar)2 aAt P2 " 4rjAr P3 = Pt - P2 P4 = pi + P2 At(kA + kB ps P6 = ' 2c(Ar)2 At(kA + kB 4crjAr kA , kB C =-H-- aA aB qi = = q3 = q4 = q6 = aAt 2(Az)2 kAAt c(At)2 kBAt c(At)2 At (kA + kB) 2c(Az)2 kAAt c(Ar)2 kBAt c(Ar)2 Tridiagonal coefficients 1- Point (i,j) in the melt or scrap piece - first At/2: ai = Ci = -q[ bi = 1 + 2q, di = p3Tni,j_i + (l-2pi)T\j + p4Tni,j+i - second At/2: aj = "P3 bj = 1 + 2p i cj = -p4 ^ dj = qiT*n j + (l-2qi)T*ij + qiT*i+i j (10) 2. Point (i.j) on the boundary surface parallel to r axis separating the materials A (left) and B (right) - first At/2: ai = -qi bi = 1 + q2 + q3 Ci = -q3 di = (p5-p6)Tni,j-i + (l-2p5)Tni,j + (p5+p6)Tni,j+i (12) - second At/2: aj = -(P5-P6) bj = l+2p5 Cj = -(ps+pe) dj = q2T j-ij + (l-q2-q3)T i.j + q3T i+ij (13) 3. Point (i,j) on the boundary surface parallel z axis separating the materials A (down) and B (up) - first At/2: ai = Q = -q4 bi = l+2q4 di = (q5-q6)Tni,j-i + (l-2p5)Tni,j + (q6+p6)Tni,j+i - second At/2: aj = pe - q5 bj = l+2p5 Cj = -(qe;Pe) dj = q4T*i-ij + (l-2q4)T*ij + q4T*i+i,j 4. Point (i,l) out of the boundary surface - first At/2: ai = Ci = -qi bi = l+2qi di = (l-4pi)T*i,i +4piTni,2 - second At/2: bj = l+4pi cj = 4pi * * * dj = qiT i-i,i + (l-2qi)T i,i + qiT i+i,i (14) (15) (16) (17) 5. Point (i,l) on the boundary surface which separates the materials A (left) and B (right) - first At/2: ai = -q2 bi = 2q4 + 1 ci = -q3 di = (l-4p5)Tni,i +4p5Tni,2 - second At/2: bj = 4p5 + 1 cj = -4p5# dj = q2T*i-i,i + (l-2q4)T\i + q3T*i+l,l (18) (19) Appendix 3 One-dimensional mathematical model of fusion of steel cylinder consists of the solution of a differential equation of heat conduction vvith adequate initial and boundary conditions 5T ,52T (11) - 1 3Tv at 00 (20) T(r,0) = To T(l,t) = TL I T(r,t) | < M vvhere M is a positive real number. It is convenient to consider insted of the equation (20) the equation 9T 32T J_ 3T at " 3r2 + r dr and then to repleace t by at. Applying the Laplace transform6, we find (21) d2e 1 dO T^ + T - s6 = _T<> dr r dr (22) 9(1,s) = Tl/s, 0(r,s) is connected. The general solution of this equation is given in terms of Bessel functions as 9(r,s) = c,J0 (ir Vš) + c2Y0 (ir Vš) + - (23) Since Yo(iWs) is unbounded as r —> 0, we must choose ca = 0. Than T 9(r,s) = c,J(1 (ir Všj + — From the boundary conditions we find 9(1,s) = c,J0 = ^ c, = tl-t„ sJ„ (i Vš) (24) (25) (26) (27) T„ J« (ir Vs) Thus 8(r,s)=-f+ (Tl-T„) r s sJ0(iVs) After complex inversion this equation acquires the form 1 r^" estJ0 (ir Vš) T(r,t) = T„ + (Tl-T„) — - J 2n\ i„sJ0 (t Vs) ds and the final solution is obtained as T(r,t) = TL - 2(Tl-T0) ^ r f . (K) (29) where A.2, ..., Xn, ... are positive zeros of equation Jo(^-n) = 0, vvhich are given in Table 1. Table t: Zeros of equation Jo(A.n) = O7. n An 1 2,40482 55577 2 5,52007 81103 3 8,65372 79129 4 11,79153 44391 5 14,93091 77086 6 18,07106 39679 7 21,21163 66299 8 24,35247 15308 9 27,49347 91320 10 30,63460 64684 6 References 1 M. Ja. Madžibožski, Osnovi termodinaraiki i kinetiki staleplaviteljnih processov, Viša škola, Kiev, 1986 2 F. Oeters, Metallurgie der Stahlerstellung. Springer-Verlag, Berlin, 1989 3 V. Grozdanic. Metalurgija, 30, 1991, 1/2, 47-50 4 V. Grozdanic, Livarstvo, 5, 1990, 1, 3-11 5 J. Douglas, H. H. Rachford, Trans. Amer. Malh. Soc., 82, 1956, 421 6 M. R. Spiegel. Laplace Transforms, McGraw-Hill, New York, 1965 1 M. Abramovvitz, I. A. Stegun (eds.), Handbook of Mathematical Functions vvith Formulas, Graphs and Mathematical Tables, National Bu-reau of Standards. VVashington, 1964 Equilibrium Grain Boundary Segregation of Antimony in Iron Base Alloys Ravnotežna segregacija antimona po mejah zrn v zlitinah železa in antimona r. Mast1, H. Viefhaus, M. Lucas, H. J. Grabke, Max-Planck-lnstitute, Dusseldorf, Ger-many Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The equilibrium grain boundary segregation of antimony ivas investigated in iron base aiioys (Fe-Sb, Fe-C-Sb, Fe-Ni-Sb) after anneaiing at temperatures betvveen 550°C and 750°C. Utiiizing Auger electron spectroscopy (AES) the concentration of antimony at intergranular fracture faces was determined as a function of buik concentration and equilibration temperature. The segregation of antimony in Fe-Sb alloys with 0,012 wt.% - 0,094 wt.% Sb was described by the Langmuir-McLean equation. The evaluation leads to the free enthalpy of segregation AGsegr. = -19 kJ/mol - T 28 J/mol K. For Fe-0,93 wt.% Sb and Fe-1,91 wt.% Sb a thermodynamic calculation is not possible because of intergranular antimonides had formed. Scanning electron micrographs (SEM) of fractured samples show that the percentage of intergranular fracture increases with an increasing coverage of antimony at the grain boundaries. The addition of carbon to Fe-Sb alloys results in a higher grain boundary cohesion which is caused by two effects of carbon, displacement of antimony from the grain boundaries by carbon and enhanced grain boundary cohesion. In the Fe-Ni-Sb alloys an additional segregation of nickel was found at the grain boundaries but no enhanced antimony segregation, as expected from previous models of other authors, assuming Ni-Sb cosegregation. Key words: grain boundary segregation, antimony equilibrium segregation, Fe-Sb alloys, Fe-C-Sb alloys, Fe-Ni-Sb alloys, segregation thermodynamics, Langmuir-McLean equation, Auger electron spectroscopy (AES), intergranular fracture, embrittlement, site competition, Charpy impact tests Ravnotežna segregacija antimona po mejah zrn v zlitinah z železnoosnovno (Fe-Sb, Fe-C-Sb, Fe-Ni-Sb) po žarjenju v temperaturnem področju od 550°C do 750°C. Z metodo spektroskopije Augerjevih elektronov (AES) je bila določena koncentracija antimona na interkristalnih prelomnih ploskvah kot funkcija vsebnosti antimona v osnovnem materialu in ravnotežne temperature. Segregacija antimona v Fe-Sb zlitinah z 0.012 ut.% - 0,094 ut.% Sb je opisana z Langmuir McLeanovo enačbo izračunana je bila prosta entalpija segregacije AGsegr. = -19 kJ/mol - T 28 J/mol K. Za zlitini Fe -0,93 ut.% Sb in Fe -0,91 ut.% Sb termodinamični izračuni niso mogoči zaradi tvorbe interkristalnih antimonidov. Posnetek z vrstičnim elektronskim mikroanalizatorjem (SEM) prelomljenih vzorcev kaže, da odstotek interkristalnega preloma narašča z naraščajočo segregirano plastjo antimona na mejah zrn. Dodatek ogljika v Fe-Sb zlitino povzroči večjo kohezijo med posameznimi zrni, ogljik namreč izrine antimon z mej zrn in zviša kohezijo kristalnih mej. V Fe-Ni-Sb zlitinah je bila določena še segregacija niklja na mejah zrn ne pa tudi povečana koncentracija antimona kot je bilo pričakovati po prejšnjih modelih nekaterih avtorjev, ki so predvideli skupno segregacijo Ni-Sb. Ključne besede: segregacija na mejah zrn, ravnotežna segregacija antimona, Fe-Sb zlitine, Fe-C-Sb zlitine, Fe-Ni-Sb zlitine, termodinamika segregacij, Langmuir McLeanova enačba, spektroskopija Augerjevih elektronov (AES), interkristalni prelom, krhkost, tekmovanje za prosta mesta na površini, Charpyjev udarni preizkus 1 Introduction The increased usage of low quality scrap in steel production will lead to a higher content of antimony in steels, which may have a deleterious effect on material properties. The presence of antimony (and/or other tramp elements such as P, Sn, S, As) induces temper embrittlement of low alloy ferritic steels by segregation to the grain boundaries during application at higher temperatures1'2-1. The driving forces for such an enrichment in a range of a monolayer are the decrease of interfacial en-ergy and the release of elastic energy. Especially the lat-ter effect is important for antimony because of its large atom size compared to iron atoms. Many researches have been shown that the amount of antimony segregation de-pends on the total composition of the steel. However, there is no uniform evidence how other alloying components, especially nickel2-4'5, influence antimony segregation. Dr.Sc. Ralph MAST Max-Planck-Inslitut fiir Eisenforschung GmbH Postfach 140 444. 40074 Dusseldorf, Germany Therefore, the equilibrium grain boundary segregation of antimony and its effects on material properties vvere examined in simple iron base alloys to avoid the complex chemistry of multicomponent steels. The degree of coverage vvas determined by Auger electron spectros-copy (AES) on the intergranular fracture faces after fracture by impact inside the UHV chamber. The influence on the mechanical behaviour vvas studied by scanning electron microscopy (SEM) and Charpy impact tests. 2 Experimental procedure The alloys used in this study vvere melted in a vacuum induction furnace. The chemical compositions are listed in Table 1. Small amounts of manganese (0,02 wt.%) vvere added to each alloy to tie up sulfur, vvhich has a strong tendency for grain boundary segregation3 and may hinder antimony segregation. The ingots of the Fe-Sb, Fe-C-Sb and Fe-Ni-Sb al-loys vvere hot forged and then machined into rectangular specimens. The Fe-Sb and Fe-Ni-Sb samples vvere heat treated by austenitizing at 1060°C for 70-90 min, air Table 1: Chemical composition of the Fe-Sb. Fe-C-Sb and Fe-Ni-Sb alloys (wt.%) Alloy Sb C Mn P S Fe-Sb 1 0,012 0,005 0,027 0,0015 0,0013 Fe-Sb2 0,049 0,0048 0,027 0,0011 0,001 Fe-Sb3 0,094 0,0057 0,027 0,001 0,0011 Fe-Sb4 0,93 0,006 0,026 0,0013 0,0012 Fe-Sb5 1,91 0,0039 0,028 0,0014 0,0012 Fe-C-Sb 1 0,056 0,0043 0,025 <0,002 0,0013 Fe-C-Sb2 0,053 0,0085 0,023 <0,002 0,0013 Fe-C-Sb3 0,052 0,0144 0,023 <0,002 0,0014 Fe-C-Sb4 0,094 0,0057 0,027 0,001 0.0011 Alloy Sb Ni C Mn P Fe-Ni-Sb 1 0,049 0,53 0,0035 0,022 <0,002 Fe-Ni-Sb2 0,049 2,85 0,0069 0.024 <0,002 cooling, and then tempering at 780°C for 168 h and water quenching. These two heat treatments were per-formed in flowing wet hydrogen to decrease the bulk carbon concentration below 10 wt.-ppm. The Fe-C-Sb alloys were annealed in flowing dry argon to avoid carbon losses. The samples were homoge-nized at 1060°C for 70 min and air cooled. Afterwards they were recrystallized at 780°C for 2 h and water quenched. Then ali specimens were held at ageing temperatures of 550°C, 600°C, 650°C, 700°C and 750°C for different periods of time, to establish the equilibrium concentration of antimony at the grain boundaries. The time neces-sary for equilibration at each temperature can be as-sessed using an equation proposed by McLean6. AES measurements confirmed that the calculated time was long enough to reach equilibrium segregation; the condi-tions of each exposure are listed in Table 2. Table 2: Conditions for the establishment of segregation equilibria Ageing Temperature/ Ageing Time Exposure Conditions 550°C/600 h vacuum/quenched in vvater 600°C/140 h vacuum/quenched in vvater 650°C/ 50 h flovving argon/quenched in vvater 700°C/ 5 h flovving argon/quenched in vvater 750°C/ 2 h flovving argon/quenched in vvater The amount of grain boundary segregation was to be measured by AES, vvhich is conducted in UHV to avoid surface contamination. After cooling to about -120°C the cylindrical notched specimens were fractured by impact in the UHV chamber of the spectrometer. The fracture surface was then imaged by operating the electron beam in a scanning electron microscope (SEM) mode to distin-guish between intergranular and transgranular areas. Auger spectra were taken from at least 10 individual grain boundary facets using a cylindrical mirror analyzer (CMA) and the results were averaged. The peak-to-peak heights of antimony (454 eV), nickel (848 eV) and carbon (271 eV) were related to the iron peak at 651 eV. The entire analysis of each fracture face had to be completed within approximately 3 h to prevent contamination ef-fects. The operating conditions were as follows: primary beam energy 5 kV, primary beam current 3 x 10"6 A, and primary beam size 10 pm. To estimate the degree of coverage of antimony at the grain boundaries, it can be assumed that antimony is uni-formly distributed on both fracture faces. This supposi-tion was verifted by some AES measurements in which opposite fracture facets vvere investigated7. From LEED studies of surface segregation on Fe-Sb single crystals a calibration factor had been obtained vvhich converts the peak-to-peak height ratio to the degree of coverage8. Supplementary surface analytical methods vvere em-ployed. The binding state of core electrons of segregated antimony was determined by X-ray photoelectron spec-troscopy (XPS), vvhile scanning Auger microscopy (SAM) vvas applied to examine the distribution of segregated elements on grain boundary facets. The fracture type and the mechanical properties vvere investigated using SEM (accelerating voltage 20 kV) and Charpy impact tests (DIN 50115). 3 Results and discussion 3.1 Fe-Sb alloys Typical Auger spectra of transgranular and intergranular areas are represented in Figure 1. On transgranular fracture surfaces of the Fe-0,094 wt.% Sb alloy, no antimony peak vvas observed, since the bulk concentration is belovv the detection limit of the AES method. The oxygen peak is due to adsorption from the residual atmosphere after breaking the sample. The spectrum taken on a grain surface of the same alloy clearly indi-cates the enrichment of antimony vvhich is caused by grain boundary segregation. Figure 2 illustrates that the average coverage of anti-mony at the grain boundaries increases vvith increasing bulk concentration and decreasing equilibration temperature. The scatter of the data indicated by the error bars in one curve is rather large (25% - 30% of the mean value) due mainly to the follovving reasons: a) The segregation of antimony may be strongly de-pendent on grain orientation as indicated by surface segregation studies on Fe-Sb single crystals8. b) The examined areas have different distances and different surface normals to the cylindrical mirror ana-lyzer (CMA). c) The degree of coverage is calculated from measurements on only one side of the intergranular fracture face. It vvas verified by some AES measurements in vvhich opposite fracture facets vvere investigated that the average grain boundary antimony concentration is nearly the same on both fracture faces7. The assumption that an- 100 200 300 400 500 Kinetic Energy [eV] 600 700 100 200 300 400 500 Kinetic Energy [eVj 600 700 .6 AH, ln-- lnxoh = - 1-0 Sb RT a cexs segr ^ _ '-'segr R (D which expresses the relationships between bulk concentration (mole fraction) xsb, temperature T, and degree of coverage 0, at the grain boundaries. The results according to the Langmuir-McLean equation are plotted in Figure 3. The estimation yields the segregation en-thalpy AHsegr. = -19 kJ/mol ± 5 kJ/mol and the segregation entropy ASsegr. = 28 J/mol K ± 6 J/mol K. The free enthalpy of segregation in a-iron can be expressed as follows: AG«gr = -(19 kJ/mol±5kJ/mol) - T(28 J/mol K±6 J/mol K) o\ 18-16-14-12- 10-8 -6 -4 -2 - 0.094.% Sb Fe-Sb 600 650 700 Ageing Temperature [°C] 750 Figure 2: Grain boundary concentration of antimony plotted as a funetion of equilibration temperature for the alloys Fe - 0,012 wt.% Sb, Fe - 0,049 wt.% Sb and Fe - 0,094 wt.% Sb Slika 2: Koncentracija antimona na kristalni meji kot funkcija ravnotežne temperature za zlitine Fe - 0,012 ut.% Sb, Fe - 0,049 ut.% Sb in Fe - 0,094 ut.% Sb Figure 1: Auger spectra of fracture surfaces of Fe - 0,094 wt.% Sb alloy after annealing at 650°C. a) cleavage facet, b) intergranular fracture surface Slika 1: AES spekter prelomnih površin Fe - 0,094 ut.% Sb po žarjenju pri 650°C. a) prelomna ploskev, b) interkristalna prrlomna površina timony is equally distributed is probably not true for each single intergranular area. In spite of the large scatter of the data, a thermody-namic calculation was attempted, applying the Langmuir-McLean equation 6,8 6,6 X 6,4 C 6,2 — ■Zl 6,0 s. 5,8 CD 5,6 5,4 5,2 ■ 0.012 %Sb • 0.049 %Sb * 0.094 %Sb 0,95 1,00 1,05 1,10 l/T [K 'xl0'] 1,15 1,20 1,25 Figure 3: Langmuir-McLean plot of the data in Figure 2 Slika 3: Langmuir-McLeanov diagram podatkov iz slike 2 The segregation enthalpy value is low compared to values for phosphorus (AHsegr = -34 kJ/mol)9 or tin (AHsegr = -23 kJ/mol)10, this indicates the low tendency for grain boundary segregation of antimony in iron. It would be unreasonable in the present thermody-namic calculations to include the AES data for the Fe -0,93 wt.% Sb and Fe - 1,91 wt.% Sb alloys, since un-known antimonides had formed at the grain boundaries. In Figure 4, a typical scanning electron micrograph and the corresponding elemental map for antimony on the same intergranular area of the Fe - 0,93 wt.% Sb alloy indicate star shaped antimonides. In spite of the low tendency for grain boundary segregation, antimony has a strongly embrittling effect. The relationship betvveen the percentage of intergranular fracture and the grain boundary coverage of antimony is demonstrated in Figure 5. With increasing enrichment of antimony at the grain boundaries the fracture mode at 100 <=\ 80 i 1-1 PH S 3 fe 40 (D Figure 4: Intergranular antimonides observed in Fe - 0.93 wt.% Sb after annealing at 650°C; a) scanning electron micrograph. b) corresponding scanning Auger image of Sb Slika 4: Interkristalni antimonidi opaženi v Fe - 0,93 ut.% Sb po žarjenju na 650°C; a) posnetek z vrstičnim elektronskim mikroanalizatorjem, b) vrstični Augerjev posnetek low temperatures (about -120°C) changes from trans-granular to intergranular already at rather low grain boundary concentrations. The influence of antimony segregation on the mechanical properties vvas also studied by Charpy impact testing. The transition temperature determined Tt is a measure of the embrittlement of iron base alloys. Tt is defined as the temperature vvhere half of the difference value betvveen the impact vvork necessary for ductile fracture and the impact vvork for brittle fracture is reached. For the Fe-Sb alloys a shift of the impact transition temperature to higher values is expected vvith in-creasing antimony concentration at the grain boundaries. This supposition is verifted in Figure 6. For each of the tvvo investigated alloys a higher transition temperature is obtained vvith increasing coverage of antimony at the grain boundaries. Hovvever, the Fe - 0,094 wt.% Sb alloy tempered at 750°C has a lovver transition temperature than the Fe - 0,049 wt.% Sb alloy annealed at the same temperature. The observed phenomenon can be ex-plained by the different average grain size of these materials (Fe - 0,049% Sb: 0,21 mm; Fe - 0,094% Sb: 0,08 mm), vvith increasing antimony concentration the grain 0,1 0,2 0,3 0,4 "0,5 Peak-to-Peak Height Ratio [I(Sb)/I(Fe)] Figure 5: Percentage of intergranular fracture versus peak-to-peak height ratio I(Sb)/I(Fe) Slika 5: Odstotek interkristalnega preloma v odvisnosti od razmerja višine vrhov I(Sb)/I(Fe) size decrease vvhich leads to a higher strength of the material. One possible way to explain the embrittling behav-iour of antimoy is to apply quantum mechanical mod-els"12. The main conclusions of these calculations can be summarized as follovvs: The segregated antimony atoms are electronegative vvith respect to the host metal iron. Consequently electronic charge is transferred from iron to antimony. This charge transfer leaves fevver electrons to participate in 100 p ai l— ■D 03 d) O. E C o +-J tn C ni 80 60 40 20 0.049 wt.% Sb 0.094 wt.% Sb 82 °C 52 °C 25 °C 18 °C 0.024 0.035 0.042 0.064 Peak-to-Peak Height Ratio [I(Sb)/I(Fe)] Figure 6: Dependence of the transition temperature on the grain boundary antimony concentration for Fe - 0,049 wt.% Sb and Fe - 0,094 wt.% Sb alloys Slika 6: Odvisnost koncentracije antimona na mejah zm od prehodne temperature za zlitine Fe - 0,049 ut.% Sb in Fe - 0,094 ut.% Sb Figure 8: SEM of a faceted grain boundary in Fe - 0,094 wt.% Sb after annealing at 650°C Slika 8: SEM posnetek facetirane meje v zlitini Fe - 0.094 ut.% po žarjenju na temperaturi 650°C Figure 9: Pore at a grain boundary facet of Fe - 0,094 wt.% Sb after annealing at 600°C; a) SEM. b) corresponding scanning Auger image of Sb Slika 9: Razpoka v kristalni meji Fe - 0,094 ut.% po žarjenju na temperaturi 600°C: a) SEM posnetek, b) odgovarjajoči SAM posnetek Sb surface energy and such pores will intensity the observed embrittlement of the material. 2,5x1 o6 — 2x1o6 e jz I l,5xl06 Sb (segregated) 5x10'" 542 540 538 536 534 Binding Energy [eV] Figure 7: Photolines of pure Sb and segregated Sb in Fe-Sb alIoys Slika 7: XPS krivulje čistega Sb in segregiranega Sb v Fe-Sb zlitini Mg K,, Sb 3d3n Sb (elcmcntal) 537,40 eV the iron-iron bonding and these bonds at the grain boundary will be weakened. XPS measurements on a large area of intergranular fracture of Fe - 0,93 wt.% Sb alloy after annealing at 600°C show that the energies of the Sb 3d electron levels of segregated and pure antimony are distinctly different (Figure 7). The energy shift of about -0,5 eV in com-parision to pure antimony indicates an electron transfer to segregated antimony, as expected in the above model. It is also possible to explain the embrittling behaviour of antimony in another way by taking into consideration that the grain boundaries often are facetted, as illustrated in Figure 8. The segregation of antimony induces a re-construction of the grain faces vvhich results in a de-crease of grain boundary cohesion. On some intergranular areas pores were detected with an average diameter of 2 pm as can be seen in Figure 9. An antimony map recorded for the same area, shovvs an-tirnony enrichment vvithin this pore. Segregated anti-mony certainly favours the formation of such pores since its surface segregation causes a pronounced decrease of 3.2 Fe-C-Sb alloys Samples vvith different antimony and carbon contents vvere investigated to study the effect of carbon on anti-mony grain boundary segregation. The fracture faces of the Fe-C-Sb alloys vvith 0,049 wt.% Sb shovv transgranu-lar fracture caused by the carbon content. The higher cohesion of these materials compared vvith corresponding Fe-Sb alloys is due to the fact that antimony is displaced from the grain boundaries by carbon, according to the equation C(dissolved) + Sb(segregated) = C(segregated) + Sb(dissolved) (2) The mutual displacement of these tvvo elements corresponding to the displacement equilibria in the systems Fe-C-P9 and Fe-C-Sn10, vvas proven for the Fe-C-Sb al-loy vvith 0,094 wt.% Sb, as shovvn in Figure 10. The average grain boundary concentration of antimony de-creases vvith increasing grain boundary and bulk concentration of carbon. Simultaneously the percentage 0.14 | 0.12 .g" 0.10 "H (2 £ 0.08 nb '5 SE 0.06 ■S U 0.041 o "I 0.02 0 ▲ i / C/Fe ............ • ■ / ' /s......................... ■ 700°C --•-- 650°C —i— 600°C /s r ■'•»■'■ ' 1 ■ 40 10 20 30 40 50 Carbon Content [wt.-ppm] 60 Figure 10: Dependence of the Sb and C grain boundary concentrations on the bulk concentration of carbon in Fe - 0,094 wt.% Sb Slika 10: Odvisnost koncentracij Sb in C na mejah zrn od koncentracije ogljika v osnovnem materialu Fe - 0,094 ut.% Sb 20 O 0 Dpv17. At lower temperatures (~300°C), C segregated to the surface due to very high diffusion coefficient in compari-son to Si and P, although the bulk concentration was at very low 15 ppm. At higher temperatures, C atoms were displaced by Si atoms'8. The P and S atoms displaced the silicon at higher temperatures17. Their bulk diffusion co- Steel Fe-Si-Sn(0.05%) Grain orientation (100) Grain 1 2 3 4 5 7 PHR Sn/Fe 0.23 0.31 0.29 0.40 0.30 0.40 Orientation (144) (025) (118) (111) (5913) (236) free surfaces betvveen inclusion (A1N, AbCh) and matrix clearly indicated that the considerable tin segregation oc-curs at the interface. The degree of tin segregation at the interface is five times larger than at the grain boundaries. 3.2 Surface segregation Scanning Auger image (SAM) of non-oriented electrical steel heated to 800°C for 10 minutes was taken. The orientation of individual grains vvas determined by the method described in our previous publication15. Figure 2 shovvs SEM and SAM images of surface. A differ- Figure 2: a) SEM image of 0.2 mm thick non-oriented electrical steel alloyed with 0.1% Sn. b) a SAM image Sn-MNN transition recorded on a same area, c) table shows a relation between grain orientation and Sn PHR Slika 2: a) SEM posnetek površine neorientirane elektro pločevine legirane z 0.1% Sn, b) SAM posnetek Sn MNN prehoda posnet na istem mestu, c) tabela podaja zvezo med orientacijami zrn in RVV T (°C) Figure 3: Temperature dependence of surface segregation of C, Si, P, S and Sn of electrical steels alloyed vvith 0.05% Sn a) (100) oriented grain and b) (111) oriented grain Slika 3: Temperaturna odvisnost površinske segregacije C, Si, P, S in Sn za elektro pločevino legirano s 0.05% Sn a) zrno (100) orientacije in b) zrno (111) orientacije efficient was rather lovv, but their segregation enthalpy vvas very high, so tin started segregating significantly above 600°C. The kinetics study confirmed the orientation dependence of tin surface segregation as well as thickness of segregated layer. It was ascertained" that on (100) and (111) faces, the segregation of tin was beyond one monolayer, due to the strong decrease of surface energy. On a surface with a (111) orientation FeSn intermetallic compound of one unit celi thickness was found. Our measurements showed that tin surface coverage dependence on tin bulk concentration and © value approached one for (100) and (111) orientation. 3.3 Texture measurements The textures of 0.5 mm thick electrical steels were measured on the surface and in the middle plane after the half of the sheet thickness were removed. Taking into ac-count that approximately six crystal grains constitute the 0.5 mm thick cross-section steel sheet and the fact that penetration depths of x-rays vvere less than 0.1 mm one might conclude that there vvere analysed some grains vvhose grovvth vvas not affected by the surface segregated tin. Nevertheless, there vvere not more than 10% of such grains. {001} {110} {111} {111} {001} <100> {011} o oo> {001} <1 to> {110} {111} {m} {001} {011} <0"ii> ooo> f (g) t s h x 0V.Sn a-fibre - O 0,0 5 7. Sn 0° 30" 60* 90 Figure 4: Fibre diagram of recrystallized texture for electrical steels measured a) in the middle plane and b) on the surface Slika 4: Diagram vlaken rekristalizacijske teksture za elektro pločevine merjen a) v sredini in b) na površini The orientation distribution functions (ODF) f (g) vvere calculated from the (200), (110) and (211) pole figures. The textures vvere presented as a, y and r| fibres. Figure 4 shovvs texture fibres in the middle plane (a) and on the surface (b) of electrical steels alloyed vvith and vvithout tin. The volume fraction of grains vvith the (100) planeš measured on the surface and in the middle plane increased to the order of tvvo compared the steel vvithout ■ tin vvith the steel alloyed vvith 0.05% tin. Less hard magnetic orientations vvere found on the surface. Texture development during the recrystallization vvas. Steel alloyed vvith 0.05% Sn, vvhich had previously been aged 25 hours at 550°C, compared to the steel vvithout tin, shovved an increase of (100) planeš parallel to the rolling direction to the order of three. 4 Conclusions Grain boundary and surface segregation of tin in non-oriented electrical steels vvere determined. Maximum equilibrium segregation on the surface vvere reached at 750°C and approached for majority of orientations one monolayer. One iron atom on the surface corresponds to one segregating tin atom. It vvas proved that thickness of tin segregating layer depended of tin bulk concentration. Tin segregation vvas controlled by bulk diffusion; thus, the equilibrium enrichment of tin on the surface vvas slightly faster for a specimen vvith higher tin contents. The tendency for tin surface segregation vvas much higher compared to grain boundary segregation. At equi-librium grain boundary segregation only 7 and 3% of tin atoms vvere found on a grain boundary for steel alloyed vvith 0.1 and 0.05% Sn, respectively. Different crystallographic orientations can provide different sites for segregating tin atoms. During the re-crystallization tin atoms segregated on the surface and also at the grain boundary and so decreased the surface energy of crystal grains selectively. The obtained results confirmed our supposition. Tin segregation took plače during the recrystallization and decreased the surface energy of crystal grains vvith (100) and (110) plains parallel to the sheet surface. Textures represented as sections through three-dimensional orientation distribution space in fixed directions shovved that volume fraction of magnetically soft grains increased for tvvo times compared to steel vvithout tin. Slightly better textures vvere obtained near the surface than in the middle plane of 0.5 mm thick steel sheet. The best results vvere obtained for steel alloyed vvith 0.05% Sn. We sup-pose that only a certain level of segregation promotes de-sired selective grain grovvth. 5 References 1 H. Viefhaus, Analytica Chimica Acta, 297, 1994, 43-53 2 H. J. Grabke, V. Leroy and H. Viefhaus, 1SIJ International, 35, 1995, 2, 95 3 Lyudkovsky, P. K. Rastogi and M. Bala, Journal of Metals, 1. 1986, 18 4 H. Shimanaka, T. Irie, K. Matsumura, K. Nakamura, J. Magn Magn. Mat., 19, 1980, 63 5 M. Jenko, F. Vodopivec, H. J. Grabke, H. Viefhaus, B. Praček, M. Lu-cas and M. Godec, Steel Research, 65, 1994, 11 6M. Jenko, F. Vodopivec, B. Praček, M. Godec, D. Steiner, J. Mag. Mag. Mat., 133, 1994, 229 7M. Jenko, F. Vodopivec, H. J. Grabke, H. Viefhaus, M. Godec and D. Steiner Petrovič, Journal De Physique IV, 5, 1995, C7-225 8 S. Nakashima, K. Takashima, J. Harase and K. Kuroki, J. Japan Inst. Metals, 55, 1991, 12, 1392-1399 "Lyudkovsky and P. Rastogi, Metallurgical Transactions A, 15 A, 1984, february, 257 10 K. Iwayama, K. Kuroki, Y. Yoshitomi, K. Homma and T. Wada: J. Appl Phys„ 55, 1984, 2134 11 H. Viefhaus and M. Rusenberg, Surface Science, 159, 1985, 1-23 12R. Mast, H. J. Grabke, M. Jenko and M. Lukas, in print 13 V. Rusenberg, H. Viefhaus; Surf. Sci., 172, 1986, 615 14 Beguinot and P. Lesbats, Metalography, 10, 1977, 115-119 15 M. Godec, M. Jenko, F. Vodopivec, M. Ambrožič, D. Mandrino, L. Kosec, M. Lovrečič Saražin, Kovine, zlitine, tehnologije, 28, 1994, 1-2, 105-109 16 W. Jager, H. J. Grabke, R. Moller, 4th International Conference, Portorož Jugoslavija, 1985 17H. J. Grabke, V. Leroy and H. Viefhaus, ISIJ International, 35, 1995, 2, 95-113 18 H. De. Rugy and H. Viefhaus, Surface Science, 173, 1986, 418-438 impol industrija metalnih polizdelkov slovenska bistrica E&i v imm rasmcN IfflIFCfflO BCOES EN 29001 /ISO 9001/BS 5750 APPROVED BY BVQI Izdelki iz aluminija: pločevine, trakovi, rondele, rondelice, prometni znaki, folije, palice, cevi, profili, žice, mreže, varilni materiali Telefon: 817-521, 817-421 Telefax: 811-219 Telex: 33-113 Corrosion Resistance of NdDyFeB Basic Alloys Korozijska obstojnost osnovnih zlitin NdDyFeB S. Kobe Beseničar1, IJS Ljubljana, Slovenija L. Vehovar, IMT Ljubljana, Slovenija B. Saje, Magneti d.d. Ljubljana, Slovenija Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 Nd-Dy-Fe-B-X (X = Zr, Hf) alloys were exposed to severe corrosion conditions and the corrosion rates vvere followed by various technigues (eiectrochemistry, Tafei extrapolation method). The vveight loss vvas measured over a period of 10 weeks in a wet corrosion chamber. Corrosion products were anaiysed using X-ray diffraction and the microstructures were investigated by optical microscopy and on SEM - EDS. In aggresive media, such as diluted NaCI or H2SO4, the differences between the corrosion rates vvere small. The lowest potential difference between the anodic phase (corrosion products) and the matrix, acting as cathode, vvas observed in Nd-Dy-Fe-B-Zr alloys. Corrosion rates in fresh vvater vvere 0,30 mm/year for Nd-Dy-Fe-B alloy and 0,02 mm/year for Nd-Dy-Fe-B-Zr alloy. The same trend vvas shovvn on samples exposed to conditions of simulated condensed atmospheric humidity. The highest cumulative vveight loss occurred vvith pure Nd-Dy-Fe-B alloys and the lowest vvith the alloy improved by Zr02 addition. The corrosion rates for three different alloys were 0.089 mm/year for Nd-Dy-Fe-B alloy, 0,072 mm/year for Nd-Dy-Fe-B-Hf alloy and 0,063 mm/year for Nd-Dy-Fe-B-Zr alloy. Key words: corrosion, Nd-Fe-B alloys, permanent magnets Osnovne zlitine Nd-Dy-Fe-B-X (X = Zr, Hf) smo izpostavili agresivnim korozijskim pogojem in zasledovali korozijski proces z različnimi metodami (elektrokemija, Taflova ekstrapolacijska metoda). V vlažni komori smo merili izgubo teže v obdobju desetih tednov. Korozijske produkte smo analizirali z uporabo X -žarkovne difrakcije ter opazovanjem mikrostrukture z optično mikroskopijo in elektronskim mikroskopom opremljenim z EDS. I/ agresivnih medijih kot sta NaCI in H2SO4 so bile razlike y korozijski hitrosti med različnimi zlitinami majhne. Najmanjšo razliko potenciala med anodno fazo (korozijski produkt) in matrico, ki deluje kot katoda, smo opazili pri zlitini Nd-Dy-Fe-B-Zr. Korozijska hitrost v vodi je bila 0,30 mm/leto pri zlitinah Nd-Dy-Fe-B in 0,02 mm/leto pri zlitinah Nd-Dy-Fe-B-Zr. Enako tendenco smo opazili pri eksperimentih, pri katerih so bile zlitine izpostavljene pogojem, ki so simulirali nasičeno zračno vlago. Najvišja kumulativna izguba teže je bila dosežena s čistimi Nd-Dy-Fe-B zlitinami in najnižja z Nd-Dy-Fe-B-Zr zlitinami. Korozijske hitrosti za različne zlitine so bile 0,089 mm/leto za zlitino brez dodatkov, 0,072 mm/leto za zlitino z dodatkom Hf02 in 0,063 mm/leto za zlitino z dodatkom cirkon oksida. Ključne besede: korozija, Nd-Fe-B zlitine, trajni magneti 1 Introduction Among the rare earth based permanent magnets, Nd-Fe-B magnets have assumed an important position due to their outstanding magnetic properties1,2 and their use is stili on grovving in different fields of application3. Hovvever, cotTosion has been a problem vvith Nd-Fe-B magnets, because phases rich in rare earth elements are easily oxidised in air, especially in humid air4-5. Since corrosion can deteriorate seriously the magnetic properties and on the other hand, can also be detrimental to magnetic cir-cuits, much effort has been made to improve the corrosion resistance of Nd-Fe-B magnets. Even coating and plating are not the perfect solution to this problem, because they can be imperfect and allovv the penetration of reacting species such as moisture to the magnet surface6. Searching for a better resistance of the material itself, various referred possibilities have been studied. Narasimhan et al.7 reported that raising the oxygen content to betvveen 0,6 to 3,5% significantly improved the corrosion resistance; Kim and Jacobson reported that the addition of Al, Dy or Dyz03 improved the corrosion resistance in humid air4, vvhile Tenaud, Vial, Sagavva8 and Hirosavva et al.9 used V and Mo to improve the basic 1 Dr. Spomenka KOBE BESENIČAR Inslitul Jožef Štefan. Jamova 39 1001 Ljubljana. Slovenija corrosion resistance of Nd-Fe-B magnets. Kobe et al. reported on the beneftcial influence of ZrC>2 addition not only to the inereased coercivity, but also to the corrosion resistance of the Nd-Dy-Fe-B magnets10. Previously Nakamura11 attained better corrosion resistance of the Nd-rich phase by the substitution of Fe vvith Co and Zr, and Sagavva et al.12 improved the corrosion resistance by addition of Co and Al. Kim et al.6 influenced the corrosion resistance by varying the amount of O, C and N in the basic composition of Nd-Fe-B magnets. On the basis of the promising results in our previous work10, vve continued our studies on the influence of Zr02 and Hf02 additions on improving the corrosion resistance of the basic Nd-Dy-Fe-B alloy vvith the composition Ndi5DyiFe76B8. The corrosion resistance vvas fol-lovved over experimental periods during vvhich the samples vvere exposed to various severe corrosion conditions. 2 Experimental The basic alloys used fot the corrosion experiments vvere prepared by are melting the alloys NdFe, DyFe, FeB and Fe povvder in a pure Ar atmosphere. In order to prevent the oxidation Ti sponge vvas used as a getter for oxygen. Three different batehes vvere prepared: A - samples vvithout other additives, B - 1 wt.% hafnia vvas added before are melting, C - I wt.% of zirconia was added prior to are melting. Samples were remelted three times in order to attain a better homogeneity. Buttons of melted alloys were sliced and polished to dises, dimen-sionally appropriate for the corrosion tests. The investigations were focused on general corrosion resistance, based on electrochemical determinations of the possible passivity of electrode surfaces, or aetive corrosion. Moreover, service conditions were simulated by exposing the test specimens in a wet corrosion facility (DIN 5017), with the aim of establishing the effect of chemical composition and microstructure on the corrosion rate and the form of corrosion. The potentiodynamic anodic polarisation measurements were performed using an EG and G-PAR poten-tiostat and "Softcorr 352" softvvare. Experiments were carried out in fresh water and in various aqueous test-so-lutions containing low concentrations of aggressive ions sueh as Cl" and SO42*. Sueh media could only represent approximative atmospheric conditions in the industrial environment. Electrochemical determination of corrosion rates were performed by the Tafel plot technique. After exposing the samples to various corrosion conditions they were characterised by optical and electron microscopy (SEM/EPMA JEOL, JXA 840 A). Phases in corrosion products were identified using EDS and WDS analysis facilities and an X-ray diffractometry (Philips 1710). 3 Results and discussion 3.1 Effect of the HfC>2 and ZrC>2 additives on the corrosion rate of Nd-Dy-Fe-B alloys The example of the anodic polarisations curves pre-sented in Figure 1 indicates that ali of the three materials cannot achieve passivity. The overall shape of the curves indicates that the materials undergo aetive corrosion. It is evident that the potentiodynamic scans did not reveal any significant feature, sueh as a passive region where passi-vation is spontaneous, the pitting potentional or the eriti-cal anodic current. The conclusion from the anodic po- tentiodynamic scans of the materials carried out in different solutions was that no significant passivation oc-curred. Due to sueh polarisation behaviour of the materials, corrosion rate measurements were performed by the Tafel plot technique. The corrosion rates of the materials tested vvhen exposed in various media are presented in Table 1 and graphically in Figure 2. From these results it can be concluded that chloride ions drastically promote corrosion. As their concentration inereases, so does the rate of corrosion. The corrosion process is also particularly dramatic in acid solutions containing SO42" ions, which represent very aggressive industrial atmosphere. The corrosion rates of ali materials in fresh water are relatively favourable. In addition, the results of this investigation showed that a defined trend which favours a NDFB-Zr02 material ex-ists (Figure 3, Table 1). The same trend among the materials was observed by exposure in a wet corrosion facility, but a substantial im-provement of the corrosion properties by addition of ZrOa was not achieved. Results are presented in Table 2. 6 .386 ...........■ n-rrm].......... ................ - NFB e .iee----- NFB-Hf NFB-Zt S -e.iae - tn 3 -8.388 u -6.986 ..........^ ........1 ................—.........—....... -6.886 -7.888 -6 888 -5.888 BB8 3 B8B -Z.886 I/area (18** Figure 1: Potentiodynamie polarisation curves for three types of alloys tested in fresh water, 20°C Table 1: Corrosion rates of alloys in different media at 20°C Material Media Corrosion rate (mm/year) NFB fresh vvater 0,300 NFB-HfO: fresh water 0,530 NFB-ZrCb fresh water 0,022 NFB 0,09 M NaCl 2,120 NFB-Hf02 0,09 M NaCl 2,710 NFB-Zr02 0,09 M NaCl 2,650 NFB 0,17 M NaCl 2,650 NFB-HfCh 0,17 M NaCl 3,260 NFB-Zr02 0,17 M NaCl 3,150 NFB 0,5 M H2S04 303,0 NFB-HfOj 0,5 M H2SO4 274,0 NFB-ZrOj 0,5 M H2SO4 237,8 546 E E Figure 2: Corrosion rates of the materials exposed in various media presented graphically Figure 5: Microstructures (cross sections) of sample with Hf02 (B) and sample vvith ZrC>2 (C) (385 x) proceeds in samples A. In samples B and C the corrosion products are located mainly on the surface, especially in samples C, vvhere no deep corrosion in the bulk material was observed. The reason for such local corrosion is sup-posed to be the presence of particular phases. More detailed analyses of the phases present vvere obtained by electron microscopy. Figure 6 shows the combined BS/SE image of an SEM micrograph of sample A and spectra of phases Pi and P2. The phases present in the corrosion products of sample A vvere found to be combined Nd, Dy and Fe oxides. The ratio betvveen Nd and Fe oxides differs in the phases Pi and P2. The results of standardless quantitative analyses (ZAF correc-tion program) are presented in Table 3. Table 3: The results of standardless quantitative analyses of the oxide phases Nd203 (wt.%) Dy203 (wt.%) FeO (wt.%) Phase P, 43,35 30,18 24,47 Phase P2 07,69 - 92,31 Phase P n 37,55 24,73 37,72 Phase P|2 11,99 - 88,01 In samples B a Hf-Fe rich phase vvas detected. The combined BS/SE image of the SEM micrograph of sample B and the corresponding spectrum of phase P6 are shovvn in Figure 7. Other phases present are the matrix phase P5 (REjFenB) and RE -rich phase P7. In samples C a Zr-Fe -rich phase vvas found, mostly on the phase boundaries betvveen the hard magnetic RE2Fei4B phase (P5) and the RE -rich phase (P7). A combined BS/SE image of the SEM micrograph of sample C and the corresponding spectrum of Zr-Fe -rich phase P9 are shovvn in Figure 8. The SEM micrograph of the same sample shovving different phases in the cor-roded area and the corresponding spectra of these phases are presented in Figure 9. In samples C, the barrier based on the Zr-Fe -rich phase, vvhich exists betvveen the 0123466789 10 Time (weeks) Figure 3: Cumulative mass-loss of different alloys during 10 vveeks of exposure in a wet corrosion chamber Table 2:Corrosion rates of alloys exposed in a wet corrosion cabinet Material Environment Corrosion rate _(mm/year) NFB Wet corrosion chamber 0,089 NFB-HfO: Wet corrosion chamber 0,072 NFB-ZrO;_Wet corrosion chamber_0,063 3.2 Microstructural study Cross section of the samples A, B, C vvere ground and polished vvith diamond paste. The polished surfaces vvere examined by optical microscopy and electron mi-croscopy (SEI and BSEI). The phases present vvere ana-lysed using EDS standardless quantitative analyses. Figure 4 shovvs a comparision of the microstructures (cross sections) of sample vvithout any addition (A) and sample vvith 1 wt.% of Hf02 addition (B). Figure 5 shovvs the cross sections of the polished surfaces of samples vvith Hf02 (B) and Zr02 (C) addition. There is an obvious difference in the level of corrosion attack betvveen the three samples. The most aggressive corrosion Figure 4: Microstructures (cross sections) of sample without any addition (A) and sample vvith 1 wt.% of Hf02 (B) (385 x) Phase P9 44,22 54,02 05,76 Phase P|»_37,28_54,35_08,38 4 Conclusion The results of the corrosion experiments and analyses of RE-Fe-B-X aIloys, as well as analyses ot the corrosion products and microstructural observation and analyses show, that zirconia addition gives the most promising re- Figure 7: Coinbined BS/SE image of SEM micrograph of sample B and corresponding spectra of Hf-Fe -rich phase Ps p ODO VFS ■ 4036 10 240 Figure 8: Back scattered image of SEM micrograph of sample C and the corresponding spectrum of Zr-Fe -rich phase Pio corrosion products (in the RE -rich phase) and the hard magnetic matrix phase, prevents the propagation of corrosion. Phase Pio shovvn on SEM micrograph (Figure 9) illustrates this tentative explanation. The results of stan-dardless analyses of the Zr-Fe -rich phases found are presented in Table 4. Table 4: The results of standardless quantitative analyses of Zr-Fe -rich phases Figure 6: Combined BS/SE image of SEM micrograph of sample A and spectra of phases Pi and P2 sults in corrosion protection of the basic material. Corrosion rates in fresh water were 0,30 mm/year for Nd-Dy-Fe-B alloy and 0,02 mm/year for Nd-Dy-Fe-B-Zr alloy. The same trend vvas shovvn vvhen the samples vvere ex-posed to conditions vvhere condensed atmospheric hu-midity vvas simulated. The highest cumulative vveight loss occured vvith pure Nd-Dy-Fe-B alloys and the lovv-est vvith the aIloy improved by ZrO: addition. A tentative explanation for the difference is that the change in microstructure is obviously responsible for im-proving the corrosion resistance of Nd-Dy-Fe-B-Zr a!loy. The reason for local corrosion is the presence of particu-lar phases, (Fe-Hf, Fe-Zr) acting as an anode, vvith con-siderable potential difference betvveen these and the ma-trix. A tentative explanation for the formation of Fe-Hf and Fe-Zr -rich phases is that in the samples vvith HfC>2 and ZrOa addition, during the are melting process most probably Nd from Nd -rich phase reduces both oxides and Hf or Zr -rich phases are formed. They act as the barrier betvveen the corrosion products (in the RE -rich phase) and the hard magnetic matrix phase and to some extent prevent the propagation of corrosion. The improvement of the corrosion resistance of basic material itself can contribute significantly to the stability of coated magnetic material. j O 0CB VF5 ■ 1024 10 24« p„ ] 1 (j . k_________________jM. 1 LA::.- 0 eea v^s ■ zo*9 10 2*0 Figure 9: SEM micrograph of sample C showing various phases in the corroded area and the corresponding spectra of phases Pjo, Pu and P12 5 References 1 M. Sagawa, S. Fujimura, N. Togawa, H. Yamamoto, Y. Matsuura, J. of Appl. Phys„ 55, 1984, 2083 2 J. Croat, J. Herbst, J. Lee, F. Pinkerton, J. of Appl. Phys., 55, 1984, 2078 3 M. Hersch, Permanent Magnet Demand Fuels Continued Market Grovvth, PCIM, July 1994, p 26 4 A. Kim, J. Jacobson, IEEE Trans, on Mag., 23, 1987, 5,2509 5 A. Kim, Journ. of Material Eng., 11, 1989, 1, 95 6 A. S. Kim, F. E. Camp, S. Constantinides, Corrosion of Electronic and Magnetic Materials, ASTM STP 1148, Ed. P. J. Peterson, Am. Soc. for Testing and Materials, Philadelphia, 1992, p 68 7 K. Narasimhan, C. VVillman, E. Dulis, US Patent, No. 4588439, 1986 8P Tenaud, F. Vi al, M. Sagawa, IEEE Trans, on Magn., 26, 1990, 5, 1730 9 S. Hirosawa, S. Tomizawa, S. Mino and A. Hamamura, IEEE Trans, on Mag., 26, 1990, 5, 1960 10 S. Kobe Beseničar, J. Holc, G. Dražič, B. Saje, IEEE Trans, on Mag., 30, 1994, 2, 693 " H. Nakamura. A. Fukuro, T. Yoneyama, Proc. of the IOth Inter. Work-shop on Rare Earth Magnets and Their Application, Kyoto, Japan, 1989, p 315 12 M. Sagawa, S. Fujimura, H. Yamamoto, S. Hirosavva, Japanese Patent, No. 6338555, 1988 Lahkota prihodnosti TALUM, d.o.o., KIDRIČEVO Tovarniška ulica 10 2325 Kidričevo, Slovenia Telephone: +386 62/79 61 10 Telex: 33116 Telefax: +386 62/79 62 69 Some Aspects of lmpurity Grain Boundary Segregation in Low Alloy Cr-Mo-V Steels Segregacije nečistoč v nizko legiranih Cr-Mo-V jeklih J. Janovec1, Institute of Materials Research, Košice, Slovakia V. Magula, VVelding Research Institute, Bratislava, Slovakia P. Sevc, Institute of Materials Research, Košice, Slovakia Prejem rokopisa - received: 1996-10-04; sprejem za objavo - accepted for publication: 1996-11-01 The present work is focused on theories of grain boundary segregation. An overvievv of different approaches to solution of surface enrichment phenomenon is given in the first part. The second part is devoted to the verification of introduced theories by means of experimental results. Key words: low alloy steels. phosphorus, grain boundary segregation, non-equilibrium segregation, kinetics V članku so predstavljene teorije različnih avtorjev o segregaciji po mejah zrn. V prvem delu je podan pregled različnih razlag obogatitve prostih površin. V drugem delu smo obravnavane teorije verificirali z eksperimentalnimi rezultati. Ključne besede: nizko legirana jekla, segregacija po mejah zrn, neravnotežne segregacije, kinetika segregacije po mejah zrn 1 Introduction Enrichment of solute or solvent atoms from bulk at the grain boundaries is referred to as grain boundary segregation. Segregation is mostly attributed to the grain boundary weakening due to lowering the interface cohe-sion. As a consequence, an intergranular embrittlement occurs. Because segregation phenomenon decisively in-fluences properties of commertial materials, the grain boundary segregation has been intensively studied in last decades1"8. The present vvork deals vvith segregation theories9-15 and their experimental verification. By use of multicom-ponential alloys in the verification, the introduction of some simplifications is necessary because the segregation theories vvere mostly derived for binary or ternary solid solutions. For example, lovv alloy steels containing Fe, Cr, Mo, V, Mn, Si, C and P vvere considered to be binary Fe-P or ternary Fe-Mo-P systems13"15. 2 Segregation theories 2.1 Non-interactive equilibrium segregation The theory of equilibrium segregation for dilute bi-nary solid solution Fe-I (Fe - solvent, I - solute impurity) vvas derived by McLean to be the grain boundary analo-gous of Langmuir adsorption at free surfaces1,16. The Langmuir-McLean isotherm yields: CFCtc (11) [4D,(t-tc)] [4D,(t-tc)] The process in vvhich the desegregation is dominant can only occur vvhen CiN(tc) > CiEeci for a given temperature. It means the desegregation is limited by reaching the equilibrium grain boundary concentration. The mi-gration of grain boundaries during austenitizing and re-crystallization can also contribute to the non-equilibrium segregation in term of a svveep effect. The nature of this phenomenon resides in embedding and subsequent drag-ging of solute species by moving grain boundary. As a consequence the grain boundary enrichment of solute species occurs31'32. 2.4 Segregation under stress Stress and thermal energy does not affect the equilib-rium grain boundary concentration of impurities during the tempering (aging) significantly, but it influences segregation kinetics. Grain boundary segregation of impurities vvith higher diffusivity can be enhanced effectively by applied stress. Atoms of some impurities (e.g. carbon, nitrogen, boron) fastly occupy the convenient sites on grain boundaries and they prevent subsequently due to competition effect the segregation of other elements33"35. Shinoda and Nakamura36 studied the grain boundary segregation of phosphorus in lovv carbon steel during long-term tempering and subsequent aging under stress at the same temperature. In the first step of aging under tension (compression) phosphorus grain boundary concentration increases (decreases), then its value approxi-mate to the initial one36. Changes in impurity concentration at the grain boundaries oriented normal to the applied stress ACis can be calculated as follovvs37: 4<|)CfDIpCTAt "f RT (12) where 0 is a numerical factor of the order of unity, Ciso is the initial grain boundary concentration of impurity, p is the specific volume of alloy, c is the stress related to the grain boundary, and At is the aging time under stress. 3 Verification of segregation theories To verify the above described theories the phosphorus grain boundary segregation in five low alloy steels was investigated, Table 1. Schedules of heat treatment and phases identified in individual investigated steels termed 1, 2, 3, 4, and 5 are given in Table 2. Grain boundary concentrations of relevant elements were calculated after Daviš et al.38 from Auger spectra. Peaks of Pl20eV, Sl52eV, M0l86eV, C272eV, N379eV, V473ev, Cr529eV and Fe703eV were used in calculation. The peak of oxygen was not considered because of additional adsorption of this element on freshly fractured surface. Parameters, at which Auger spectra were achieved are given in Ref.13,15. Table 1: Chemical composition of investigated steels in wt.% Steel C Mn Si Cr Mo Ni 1 0.110 0.004 0.525 0.385 2.685 0.694 0.355 - 0.010 2 0.100 0.014 0.700 0.270 2.620 0.690 0.330 - 0.007 3 0.110 0.027 0.665 0.340 2.700 0.733 0.357 - 0.010 4 0.060 0.013 0.650 0.290 2.660 0.700 0.310 - 0.009 5 0.160 0.014 0.460 0.290 2.700 0.640 0.300 0.060 0.015 Table 2: Schedules of heat treatment and phases identified in investigated steels Steel Heat treatment Phases identified 1 1250°C/0.75h, water quenching. 680°C/20 h, water cooling, Ferrite+M7C3+MC aging at 500°C for 0.33h,lh,5h,150h 2 3 4 1250°C/0.16h, water quenching, 680°C/20h, water cooling, aging 580°C for 5 min and 150h Ferite+M7C3+MC 5 welding cycle:Tmax=1300°C At«/5=30s, 580°C/100h Ferrite+MjC+IvbCi welding cycle:Tmas=l 300°C,Atn/5=30s, 600°C/120s under stress(strain rate 300mm.h"1) Ferrite+M3C Phosphorus grain boundary concentrations measured for steels 1, 2, 3 aged at 500°C for different times shovved the best fit with McLeans non-interactive kinetic equation (2), (Figure 1). The segregation can be characterized as slow, because after 150 h aging the equilibrium vvas not reached for any of the steels. A completely different situation vvas observed for steel 4 aged at 580°C 50 45 40 35 30 25 20 15 10 5 0 50 45 40 35 30 25 20 15 10 5 0 50 45 40 35 30 25 20 15 10 5 0 iT....... a Steel 1 ♦ T 0,001 0,01 0,1 10 m,-1 100 1000 ™r b Steel 2 0,001 0,01 0,1 10 100 1000 c Steel 3 i- 0,001 0,01 0,1 1 time [h] 10 100 1000 Figure 1: McLean's non-interactive equilibrium kinetic equation fitted to values of phosphorus grain boundarv concentration for steels: 1 (a), 2 (b), and 3 (c), aged at 500°C (after15) -th -t/-l ■ calculated after (4) and (7) experiment I ' 1 ' I 1 I 1 I 1 I 1 I // 1 1 I 1 0,0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 500 550 600 time [ks] Figure 2: Interactive kinetic equations (4) and (7) fitted to values of phosphorus grain boundary concentration for steel 4 aged at 580°C (after13) (Figure 2). Here, the measured values of phosphorus grain boundary concentrations correlate vvith the curve, calculated according to equations (4) and (7). The equi-Iibrium vvas reached after 5 min, and that indicates to very rapid segregation process. The obtained results shovved that both rate of equilibrium segregation and also participations of interactions in this process are temperature dependent. McLean's non-interactive kinetic theory seems to be available for the description of segregation kinetics at lower aging temperatures (slovver segregation rates) and interactive equations are more con-venient for higher ones (accelerated segregation rates). In Figure 3 Auger spectra for the steel 5 after tem-pering (a) and short-term aging under stress (b) are shown. For loaded state the peaks of C, S, N, and Cr vvere evidently higher than for tempered one (Table 3). Differences betvveen carbon, sulphur and phosphorus grain boundary concentrations for the loaded state can be explained by different diffusivity of these elements in iron at 600°C20'39'40. Atoms of carbon and sulphur dif-fuse faster than phosphorus atoms and occupy earlier convenient sites at grain boundaries. Site competition betvveen P-C and P-S5'41 make impossible an additional phosphorus enrichment at grain boundaries. With pro-longing the aging an inerease in phosphorus and a de-crease in carbon grain boundary concentrations occur because of carbide precipitation5'42,43. Table 3: Experimentally measured grain boundary concentrations of C, S. P, N, Cr. Mo, and V for steel 5 in at.% C S P N Cr Mo V tempered 9.5±1.0 - 4.411.6 - 4.9±0.5 1,4±0.5 2.7±0.4 stressed 26.9 17.7 5.4 9.3 9.6 2.1 2.5 Higher grain boundary concentrations of chromium and nitrogen in the first period of aging under stress are probably caused by Cr-N interactive segregation. Misra and Balasubramanian34'35 supposed a Cr-N co-segrega-tion (stressing up to 5 min at 580°C) due to strong chemical interaetion betvveen these elements. After reaching the maximum coverage (depending on aging temperature), a continuous decrease in Cr and N grain boundary concentrations occurs. The shape of carbon peaks (Figure 3) indicates the occurence of carbide particles on the grain boundaries44. Then, also peaks of alloying elements, preferentially Cr, must originate partially from these particles45"47. Re-flexes originating from intergranular carbide particles mostly influence the achieved spectra and they can not be neglected in interpretation of grain boundary segregation in multicomponential alloys. 4 Concludig remarks An overvievv of the theones of grain boundary segregation is given in the present vvork. The verification of E[eV] Figure 3: Characteristic Auger spectra taken on intergranular facets of steel 5: a) tempered at 580°C for 100 h, b) aged under stress for 120 s at 600°C the theories for multicomponential Cr-Mo-V lovv alloy steels leads to the follovving findings: 1. The McLean's non-interactive equation is the most convenient for the description of equilibrium segregation kinetics at lovver temperatures (500°C), vvhile the interactive equations are more suited for the description of equilibrium segregation kinetics at higher temperatures (580°C) 2. In comparision vvith unstressed aging, the higher rates of C, S, N and Cr grain boundary segregation in the first period of the aging under stress (600°C) vvere observed 3. In the investigated multicomponential steels, an influence of carbide particles on achieved Auger spectra can not be neglected. Ackno\vledgment - This study vvas supported by the Grant Agency of Slovak Republic under grant No. 2/2001/96. 5 References 1 D. McLean, Grain Boundaries in Metals, Chap. V., Oxford Univ. Press, London 1957 2 M. P. Seah, Proc. Roy. Soc. Lond. A, 349, 1976, 535 3 C. J. McMahon, Jr„ Mater. Sci. Engng., 25, 1976, 233 "M. Guttmann, Surf. Sci., 53, 1975, 213 5 H. Erhart and H. J. Grabke, Metal Sci., 15, 1981, 401 6R. G. Faulkner, J. Mater. Sci., 16, 1981, 373 1 C. Uebing, Surface Segregation of Nonmetallic Solutes on Metals and Alloys, HRC Revievv, Wiley, New York 1996 8 P. Lejček and S. Hofmann, Solid State Mater. Sci., 20, 1995, 1 9 J. Yu and C. J. McMahon, Jr„ Metali. Trans., 11A, 1980, 277 10 C. L. Briant and H. J. Grabke, Mater. Sci. Forum, 48, 1989, 253 " M. Jenko, F. Vodopivec, H. J. Grabke, H. Viefhaus, B. Praček, M. Lu-cas and M. Godec, Steel. Res., 65, 1994, 500 12 B. Ule and V. Leskovšek, Kovine zlitine tehnologije, 29. 1995, 417 13 P. Ševc, J. Janovec and V. Katana, Scripta Metali. Mater., 31, 1994, 1673 14 P. Ševc, J. Janovec, M. Koutnik and A. Vyrostkova, Acta Metali Mater., 43, 1995, 251 15 P. Ševc, J. Janovec, M. Lucas and H. J. Grabke, Steel Res., 66, 1995, 537 16M. P. Seah and E. D. Hondros, Proc. R. Soc. Lond. 335, 1973, 191 17 M. Guttmann and D. McLean, in Inteifacial Segregation (edited by W. C. Johnson and J. M. BIakely), American Society for Metals, Metals Park, Ohio 1979, p. 261 18 M. P. Seah, Acta Metali, 25, 1977, 345 19 W. R. Tyson, Acta Metali, 26, 1978, 1471 20G. Luckman. R. A. Didio and R. W. Graham, Metali. Trans., 12A, 1981. 253 21 M. Mackenbrock and H. J. Grabke, Mater. Sci. Technol., 8, 1992, 541 22 M. Militzer and J. Wieting, Acta Metali, 37, 1989, 2585 23 M. Militzer and J. Wieting, Scripta Metali. Mater., 28, 1993, 1043 24 J. du Plessis and G. N. van Wyk, J. Phys. Chem. Solids, 50, 1989, 237 23 K. T. Aust, J. S. Armijo, E. F. Koch and J. H. Westbrook, Trans. Am. Soc. Metals, 60, 1967, 360 26T. R. Anthony, Acta Metali, 17, 1969. 603 27 Xu Tingdong, J. Mater. Sci., 22, 1987, 337 28 Song Shenhua, Xu Tingdong and Yuan Zhexi, Acta Metali, 37, 1989, 319 29Xu Tingdong and Song Shenhua, Acta Metali, 37, 1989, 2499 30Song Shenhua and Xu Tingdong, J. Mater. Sci., 29, 1994, 61 31 M. Menyhard and L. Uray, Scripta Metali, 17, 1983, 1195 32T. Abe, K. Tsukada, H. Tagawa and I. Kozasu, ISIJ Int., 30, 1990, 444 33C. L. Briant, Acta Metali, 36, 1988, 1805 34R. D. K. Misra and T. V. Balasubramanian, Acta Metali Mater., 38, 1990, 1263 35R. D. K. Misra and T. V. Balasubramanian, Acta Metali Mater., 38, 1990, 2357 36T. Shinoda and T. Nakamura, Acta Metali, 29, 1981, 1631 37T. Shinoda and T. Nakamura, Acta Metali, 29, 1981, 1637 38L. E. Daviš, N. C. McDonald, P. W. Palmberg, G. R. Riach and R. E. Weber, Handbook of Auger Electron Spectroscopy, 2nd edn., Phys. Electronics Industries, Minnesota 1976 39G. Seibel, Mem. Sci. Rev. Met., 61, 1964, 413 40 J. Kučera and K. Stransky, Mater. Sci. Engng., 52, 1982, 1 41 R. D. K. Misra and P. Rama Rao, Mater. Sci. Technol., 9, 1993, 497 42 S. Suzuki, Z. Metallk., 82, 1991, 883 43 H. J. Grabke, R. Moller, H. Erhart and S. S. Brenner, Surf. Interf. Anal., 10, 1987^ 202 44 J. Janovec, P. Ševc and M. Koutnik, Kovine zlitine tehnologije, 29, 1995, 40 45 J. Janovec, V. Magula and A. Holy, Kovove Mater., 30, 1992, 44 (In Slovak) 46J. Kočlk and E. Keilova, Mater. Sci. Forum, 97-99, 1992, 337 47 R. C. Thomson and H. K. D. H. Bhadeshia, Mater. Sci. Technol., 10, 1994, 193 slovenske železarne Jt # ZELEZABNA IESENICE ACROISli ■t-" M | i 4 -.. ■ * IZDELUJE: □ nerjavna jekla □ jeklo za elektro pločevino □ nelegirana in legirana jekla - za poboljšanje - za cementacijo □ nelegirana, mikro in malolegirana konstrukcijska jekla □ toplo valjane pločevine, trakove in lamele □ hladno valjane pločevine, široke in vzdolžno razrezane trakove □ hladno oblikovane profile □ kovinske podboje za vrata □ izsekance □ varnostne ograje NUDIMO TUDI STORITVE: □ prevaljanje □ izsekovanje (štancanje) □ krojenje □ ravnanje □ toplotne obdelave pločevin SZ ZJ ACRONI d.o.o. Cesta železarjev 8, 4270 Jesenice, tel. centrala: +386 64 861 -441, tel. direktor: 861 -443, tel. komerciala: 861 -474, fax: 861 -379, telex: 37219 ZELJSN SI Sl<|»tenija Mechanical Properties of High Temperature Vacuum Brazed HSS on Structural Carbon Steel vvith Simultaneous Heat Treatment Mehanske lastnosti visokotemperaturno vakuumsko spajkanih in istočasno toplotno obdelanih spojev V. Leskovšek1, D. Kmetic, B. Šuštaršič, IMT Ljubljana, Slovenija Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 The high temperature vacuum brazing process, at the HSS austenitization temperature makes it possible to carry out simultaneousiy the brazing of HSS on structural carbon steel and heat treatment. The advantages of this process are: increased strength of brazed joints and toughness of the part, optimum hardness and cutting edge strength for a given combination vvorking part/cutting tool. The process is economical when used in modern mass production methods, irrespective of the number of metals to be joined and heat treated. The adaptability makes the process so economical. Key vvords: high temperature vacuum brazing, hardness, microstructure, shear strength, tensile strength, vacuum heat treatment Postopek visoko temperaturnega vakuumskega spajkanja v enokomorni vakuumski peči s homogenim plinskim ohlajanjem pod visokim tlakom vodimo v območju avstenitizacije hitroreznih jekel. Prednost tako izdelanih rezilnih orodij je predvsem v doseganju želene žilavosti nosilnega dela iz konstrukcijskega jekla, v doseganju optimalne trdote rezila izdelanega iz hitroreznega jekla ter njegove odpornosti proti otopitvi pri dani kombinaciji del/orodje. Trdnostne lastnosti vezne plasti so odvisne od dodajnega materiala, tehnologije izdelave in pogojev vakuumske toplotne obdelave. Uporaba tega postopka je ekonomična, če moramo spojiti in vakuumsko toplotno obdelati le nekaj ali pa večje število orodij. Ključne besede: visoko temperaturno vakuumsko spajkanje, trdota, mikrostukture, strižna trdnost, natezna trdnost, vakuumska toplotna obdelava 1 Introduction High temperature vacuum brazing is a method of joining of metals by means of heat and filler metal in vacuum at temperatures above 900°C, yet below the melting point of the joined metals, and vvith no use of fluxes. The products are defect-free joints vvith very high bonding strength that can even reach the strength of the joined metal in many cases (e.g. steel, nickel or cobalt alloys). The high temperature vacuum brazing of HSS on structural carbon steel vvith simultaneous heat treatment is performed in single chamber vacuum furnaces, vvith uniform high-pressure gas quenching at the austenitization temperature of HSS. In this vvork high temperature brazed joints of HSS and structural carbon steel vvith simultaneous heat treatment vvere investigated. Tvvo brazing alloys based on Ni-Cr-Si and copper vvere applied as filler metals. The shear strength of an overlap joint and the tensile strength of a but joint as vvell as, the microstructure and fracture surface vvere investigated. The advantages of the process are, the requested toughness of the carrying part from structural carbon steel and the optimum hardness and cutting edge strength of HSS for the given combination of vvorking part/cutting tool. Such mechanical properties of cutting tools 1 Vojteh LESKOVŠEK. dipl.inž. Inštitut za kovinske materiale in tehnologije 1000 Ljubljana. Lepi pot 11. Slovenija manufactured in the conventional way from HSS can on!y be obtained by an additional tempering operation. Other advantages of the high temperature vacuum process are energy savings, the omittance of expensive tool steels and their cleaning, as vvell as, fevv parts are to be joined or hundreds of thousands vvhen it is economical to use vacuum brazing vvith modern mass production methods. The adaptability makes vacuum brazing of in-creasing use in the metal-joining processes. 2 Basic factors affecting the mechanical properties of the brazed joint The strength of the filler metal is one of the main factors influencing the strength properties of the brazing joint, since it is a direct measure for the strength properties of the joints. Therefore, joints brazed vvith nickel-base filler metal are stronger than those brazed vvith cop-per-base filler metal. The narrovv joint clearance causes a high capillary filling pressure; therefore, the gap should be parallel over the vvhole length of the joint. Only in this way by increased capillary filling pressure the filler metal can be aspired into the gap. The most favourable joint clearance for high vacuum temperature brazing is approximately 0 - 100 pm, vvhen measured at the brazing temperature. Figure 1 shovvs schematically the relation betvveen joint clearance and the tensile strength of the joint for flux brazing and high temperature brazing1. O) C 153 (152.87) >40 jL 60 >40 > C; bal. Ni Figure 6: Tensile test specimen vvith but joint Slika 6: Natezni preizkušanec s čelnim spojem To get a higher strength of the joint or to make the fixturing of parts to be brazed easier, a lap joint should be selected. This joint should be designed to obtain the same stability under load of the joint and of the base metal. The lap length is then function of the tensile strength of the base metal and the shearing strength of the joint: U = ^ (1) T vvhere U = length of the lap in mm, Rm = tensile strength of base metal in Nmm"2, t = shearing strength of the joint in Nmm"2, t = thickness of base metal in mm. If, in addition, a safety factor and an impairment of the joint caused by small brazing errors is taken into ac-count, then the length of the lap should be 3 to 6 times the thickness of the base metal. Generally, three times the base metal is sufficient for metals of low tensile strength; six times should be used for metals of high tensile strength1. The but joint is used for thicker parts (t > 2 mm) if a lap joint is not possible1. In contrast to soldering, the sta-bility under load of this type of joint is often sufficient for practical use if the parts are brazed. Experiments3 vvere performed on shear specimens vvith single and fourfold overlap, (Figure 5). The lamel-lae from HSS and structural carbon steels vvere, finely ground after rough machining. Measurements shovved that the surface roughness Ra = 0.44 ^im in the longitudi-nal direetion vvas equal for both surfaces. The test specimen vvith but joint shovvn in Figure 6 vvas used for the tensile test. For the brazing of the shear and tensile test specimens vvith the but joint the clearance of 80 |4m vvas chosen. The brazing temperature vvas 1120°C for specimens brazed vvith the filler metals LM and Cu, and 1160°C for those brazed vvith the filler metal 30. After diffusion heat treatment, the specimens vvere cooled in nitrogen flovv at the pressure under 5 bar abs, and than double tempered at 550°C, (Figure 7). The brazing vvas performed in a vacuum 5 x 10~2 mbar. Shear and tensile specimens vvith but joints vvere used for metalographical and mechanical research. For the investigation of endurance of brazing joint, tvvo paper knives vvith the dimensions of 425x117x10 mm and one knife vvith the dimension 560x117x10 mm, vvere manufactured from HSS W. No. 1.3343 steel and their bearing parts from the steel W. No. 1.7131 (DIN) steel, (Figure8). The filler metal marked LM vvas used for these knives and considering the knives' shape, a lap joint vvith 80 pm clearance vvas chosen. The brazing temperature vvas 1190°C. After diffusion heat treatment, the knives vvere cooled in a nitrogen flovv at a pressure under 5 bar abs, follovved by double tempering at 540°C, (Figure 7). The brazing vvas performed in a vacuum, 5 x 10"2 mbar. 4 Results and discussion 4.1 Mechanical tests Next to the required properties of structural carbon steel and HSS, the most important property is the bond strength betvveen them. Mechanical tests vvere performed on fourteen shear specimens vvith a single and four-fold overlap and length of the lap of 2 to 6 times the thickness of the base metal and three tensile test specimens brazed vvith LM and Cu filler metal. The joint clearance for high temperature vacuum brazing vvas among 50-70 |im for the specimens brazed vvith fillers LM and 30, and 20-50 um for the specimens Breazing temperature Figure 7: High temperature brazing with simultaneous heat treatment process model Slika 7: Model visoko temperaturnega vakuumskega spajkanja z istočasno toplotno obdelavo Figure 8: Paper knives manufactured by high temperature vacuum brazing vvith simultaneous heat treatment process to achieve a hardness of 64 Hrc Slika 8: Noža za rezanje papirja izdelana po postopku visoko temperaturnega vakuumskega spajkanja in istočasno toplotno obdelana na 64 HRc brazed with the copper filler. Data regarding specimens characteristics and the shear strength obtained by the In-stron tensile testing machine are summarised in table 2. Table 2: Specimens characteristics and the shear strength Tabela 2: Karakteristike preizkušancev in strižne trdnostip rekrovnih spojev Sample Filler metal Overlap Length of the lap Shear strength Nmnr2 A/l LM four-fold 3 x t > 30 A/2* LM four-fold 3 x t 27 A/3 LM four-fold 6 x t > 30 A/4* LM four-fold 6 x t 18 A/5 LM single-fold 3 x t > 71 A/6 LM single-fold 2 x t > 210 B/l 30 four-fold 3 x t > 30 B/2* 30 four-fold 3 x t 27 B/3 30 four-fold 6 x t > 20 B/4 30 single-fold 3 x t > 60 C/l Cu four-fold 3 x t > 32 C/2 Cu four-fold 6 x t > 62 C/3 Cu single-fold 3 x t > 66 C/4 Cu single-fold 2 x t > 205 * Samples fractured in bond layer; C/l- the middle lamellae made from W. No. 1.1141, end lamellae made from W.No. 1.3343; C/2 aH lamellae made from W.No. 1.3343, because of gliding in the chucks, there was no destruction of the sample; C/3- ali lamellae made from W. No. 1.1141. Results in table 2, show that rupture of samples, in general, appeared in the structural carbon steel and not in the bond layer, (Figure 9), since the shear strength of brazed joints was greater than the tensile strength of the structural carbon steel. The sample where the middle lamellae were from the steel W. No. 1.1141, was an excep-tion since the fracture appeared simultaneously on both middle lamellae. The shear strength is dependent upon the overlap shape and the lap length. The maximal shear strength was obtained on samples with a single-fold overlap and with the lap length 2 times the thickness of the base met- M_ i* j^Mf^tfr """ ! jr J < .1 .1 J "4 Figure 9: Shear specimens B/l and C/l vvith a four-fold overlap after the tensile test Slika 9: Strižna preizkušanca B/l in C/l s štirikratnim prekritjem po trgalnem preizkusu al. On samples brazed vvith filler metal LM slightly higher values vvere obtained. After vacuum heat treatment that corresponded to austenitizig and tempering temperatures for HSS M15 (AISI), the strength of the tensile test specimen vvith but joint vvas a little lovver than that for structural carbon steel. The fractures propagated mostly vvithin the bond layer and partly also in structural carbon steel and HSS. By tensile tests, the strength of specimens vvith but joint vvas strongly influenced by defects in the bond layer (sample C/8*). During tensile tests vve did not notice any elongation or reduction of area on the samples. Results of tensile tests are presented in table 3. Table 3: Strength of the tensile test specimen with but joints Tabela 3: Natezne trdnosti čelno spajkanih preizkušancev Sample Filler metal Rc (Nmnr2) Rm (Nmm"2) A/8 LM 330 445 C/7 Cu 340 475 C/8* Cu 325 345 * defects in the bond layer After mechanical tests, a metallographical examina-tion vvas performed. On the single or four-fold overlap Figure 10: Initial microcrack area propagating through the eutectic phase is in the microporous regions, sample A/2 Slika 10: Inicial za nastanek mikrorazpok, ki potekajo po eutektični fazi, so mikroporozna mesta, preizkušanec A/2 after hardening after hardening and tempering W. No. 17131 "(DIN) Figure 11: Microstructure of the bond layer in specimen C/7 Slika 11: Mikrostruktura vezne plasti na preizkušancu C/7 specimens, where fractures appeared in the structural carbon steel, only sporadic microcracks were found in the bond layer. On specimens with fracture in the bond layer, areas with microporosity were noticed, vvithout ex-ception, vvhere microcracks initiated. On specimens brazed vvith the fillers LM and 30, the fracture cracks propagated through the eutectic phase of the bond layer, (Figure 10). As mentioned above, the diffusion of carbon from HSS to structural carbon steel took plače; and conse-quently, the microstructure along the bond layer/struc-tural carbon steel consisted of pearlite and bainite. On specimens brazed vvith copper, cracks appeared at the bond layer/structural carbon steel, respectively, (Figure 11). Tensile test specimens fractured in this region, as vvell. Although carbon is not soluble in copper, the diffusion of carbon from HSS throughout the copper bond layer to structural carbon steel cannot take plače, the mi- - i = M2 (AISI) 2 1.5 1 0.5 0 0 0.5 1 2 Distance in mm Figure 12: Vickers microhardness on transition from the bond layer to HSS and structural carbon steel Slika 12: Potek mikrotrdote HV na prehodu iz vezne plasti v hitrorezno in konstrukcijsko jeklo Figure 13: Fracture through an area of eutectic and austenite phase. sample A/8 Slika 13: Prelom preko eutektika in avstenitne faze, preizkušanec A/8 crostructure along the bond layer/structural carbon steel consisted of ferrite and bainite with traces of pearlite. On the paper knife, the microhardness vvas measured across the bond layer to HSS and the structural carbon steel. The diffusion annealing was carried out with the aim to affect hardness at its transition across the bond layer and Figure 12 shows the microhardness profile ob-tained. It shows that the HSS hardness is decreased, while it is increased in the structural carbon steel. The morphology of fracture surfaces is very hetero-geneous. On the specimens brazed vvith the fillers LM and 30, it vvas possible to identify fracture surfaces that propagated in dendrit's area from those propagated in the eutectic phase and in austenite, (Figs 13 and 14). Ductile fracture on specimens brazed vvith copper propagated mostly vvithin bond layer, (Figure 15). Inclu-sions of copper oxide vvere found in the dimples. 4.2 Microstructural characterisation The used filler metals, structural carbon steel and HSS vvere examined by optical and scanning electron microscopy. The microstructure of the W. No. 1.1141 (DIN) structural carbon steel consisted of ferite-pearlite and bainite vvith a hardness of 145 HV10. The microstructure of the W. No. 1.3343 HSS consisted of a matrix of tempered martensite containing small carbide precipi-tates. The size of austenite grains vvas among 17 and 13 SG depending on the austenitization temperature and the hardness 64 HRc. Figure 16 shovvs the microstructure of the bond layer betvveen the HSS and the structural carbon steel on hard- Figure 14: Fracture surface of the specimen B/2 Slika 14: Prelomna površina preizkušanca B/2 Figure 15: Fracture surface of the specimen C/7 brazed vvith Cu Slika 15: Prelomna površina preizkušanca C/7 spajkanega s Cu - v jamicah so vključki bakrovega oksida ened and tempered specimens A/l and B/2. The specimens vvere brazed vvith the fillers LM and 30. In the bond layer polygonal grains formed because of the diffusion during brazing. The diffusion at the HSS/bond layer border seams to be quicker; therefore, more of this phase is found in the bond layer along the HSS. Along the structural carbon steel/bond layer, the bond layer vvas homogenous. The specimen B/2 vvas ex-amined by SEM, (Figure 17). Figure 16: Microstructure of the high temperature brazed and simultaneously heat treated joints of HSS and structural carbon steel, samples A/l and B/2 Slika 16: Mikrostruktura vezne plasti na preizkušancih A/l in B/2 A detailed investigation in EPMA showed that the larger polygonal grains present along the central line of the bond layer were a phase solidification grains rich in Cr, containing also Ni and Si with traces of W, Mo and V. The smaller grains were carbides, (Figure 18). The in-termetallic phase was hard. The measured microhardness was 500-600 HV. The average matrix microhardness was 195 H V. In the microstructure at the HSS/bond layer border, the effects of the diffusion processes were clearly notice-able. In the thin layer of HSS only carbides particles were noticed, martenzite matrix was transformed be- Figure 17: SEM micrograph of high temperature brazed joint, specimen B/2 Slika 17: Mikrostuktura vezne plasti preizkušanca B/2 posneta s SEM Figure 18: Distribution elements in the bond layer, sample B/2 Slika 18: Porazdelitev elementov v vezni plasti na preizkušancu B/2 cause of diffusion into austenite. This microstructure was very similar to that in the bond layer, (Figure 17). At the bond layer/structural carbon steel border, diffusion of Cr, Ni and Si to structural carbon steel oc- curred. The hardened and tempered samples brazed with filler metal LM and 30 showed along this border a thin layer rich in carbon, (Figure 16) vvith microstructure consisting of a small amount of pearlite and bainite. The diffusion of carbon vvas more rapid on the samples brazed vvith the filler LM. The microstructure of W. No. 1.7131 (DIN) struetural carbon steel used for bearing part of paper knives con-sisted of tempered martensite and bainite. Austenite grains vvere coarse, due to the high austenitization temperature. The microstructure of the W. No. 1.3343 HSS consisted of a matrix of tempered martensite containing small carbide precipitates and the size of austenite grains of 14 SG. The microstructure of the bond layer vvas iden-tical as in hardened and tempered samples brazed vvith filler metal LM. 5 Conclusion Mechanical tests and metallographic observations vvere carried on high temperature vacuum brazed and si-multaneously heat treated shear specimens vvith single and four-fold overlap and tensile test specimens vvith but joint. Tvvo Ni-Cr-Si brazed metals as vvell as copper served as filler metal. During the heat treatment, rapid diffusion processes occurred betvveen the liquid and the hard phase, especially along the HSS border. By use of Ni-Cr-Si based filler metal the formation of intermetallic phases, eutectic phases and carbides in the bond layer, and a net of eutectic carbides and voids on the austenite net along the bond layer/HSS border, vvere observed. The mechanical properties of the bond layer depend on specimen design, manufacture and heat treatment conditions. The bond layer must be as thin and as ho-mogenous as possible and must shovvs no porosity or mi-crocracks. Intermetallic phases and carbides cannot be eliminated, due to the speed of the diffusion processes, which are very high on the HSS heat treatment temperature. 6 References 1 J. VV. Bouvvman, High Temperature Vacuum Brazing, lpsen Instruc-tion Manuals 2 V. Leskovšek, D. Kmetič, J. Gnamuš and G. Rihar: High temperature vacuum brazing of HSS on construction steel vvith simultaneous heat treatment, Vuolo, 20, 1990, 2, 512- 515 3 D. Kmetič, V. Leskovšek, J, Žvokelj and J. Gnamuš: Trdnostne lastnosti visokotemperaturno spajkanih spojev v vakuumu, KZT, 27, 1993, 1-2, 69-73 Discontinuous Al-SiC Composites Formed by a Lovv Cost Chemically Activated Infiltration Technique Pridobivanje in kemijska infiltracija poroznih SiC vzorcev z Al-Si talino V. M. Kevorkijan1, zasebni raziskovalec, Maribor, Slovenija Prejem rokopisa - received: 1996-10-01; sprejem za objavo - accepted for publication: 1996-11-04 In this vvork, the preparation of porous SiC preforms from SiC particles, piatelets and whiskers have been demonstarted. Near net shape preforms, prepared by vacuum casting, vvere sintered and then covered by S1O2 layer using a cost effective oxidation in air at 1175 K for 10h. Surface engineered SiC preforms vvere than pressureless infiltrated in nitrogen atmosphere (96 vol% N2 + 4 vol% Ar) by a Al-Si melt containing 0.5-3 wt% Mg. Based on this, a mathematical model of spontaneous infiltration of a porous ceramic preform has been suggested. The roie of magnesium and nitrogen atmosphere vvas quantitatively evaluated among the other important processing parameters (porosity of preform, the specific surface area, etc.) collected in a new term named preform infiltrability. Moreover, the influence of the above listed parameters on the infiltration rate (expressed as infiltration length and function of time) has also been demonstrated. The optimal conditions for spontaneous and cost effective pressureless infiltration of porous SiC preforms by molten aluminium alloy has been selected and experimentaty confirmed. Key vvords: porous SiC preforms, vacuum casting, pressureless infiltration, infiltration kinetics, infiltrability of porous preforms V delu je opisana izdelava poroznih SiC vzorcev z vakuumskim vlivanjem in njihovo sintranje do poroznih predoblik, sestavljenih iz SiC delcev različne oblike: okrogli, heksagonalne ploščice in kratka vlakna. Z oksidacijo na zraku smo prevlekli površino poroznih SiC predoblik s tanko plastjo S1O2 in izboljšali omočljivost med SiC in Al talino. V nadaljnjem delu smo infiltriraii porozne keramične vzorce z Al-Si-Mg talino v dušikovi atmosferi (96 vol% N2 + 4 vol% Ar) pri normalnem tlaku. Na podlagi pridobljenih rezultatov smo razvili matematični model infiltracije, ki opisuje kinetiko procesa v funkciji poroznosti in specifične površine pripravljenih poroznih vzorcev, vsebnosti dušika v atmosferi in sestave Al zlitine. Model je osnova za nadaljnji razvoj tehnologije priprave Al-SiC kompozitov s spontano oz. nizkotlačno infiltracijo poroznih SiC vzorcev. Ključne besede: porozne SiC predoblike, vakuumsko vlivanje, spontana infiltracija, kinetika infiltracije, infiltrabilnost poroznih predoblik 1 Introduction The need for high strength. lightweight, and high stiffness materials has, in recent years, attracted much in-terest to the development of the manufacturing processes of metal matrix composites (MMCs)1. The most important limitation of the fabrication of MMCs by liquid-phase processes resides in the compatibility betvveen the reinforcement and the matrix2. This compatibility is par-ticularly important in the čase of aluminium-based composites, because Al is usually covered vvith a thin oxide layer that prevents vvetting, and vvhen uncovered, it read-ily reacts vvith most ceramics to form intermetalics. In particular, liquid aluminium reacts vvith SiC to produce aluminium carbide and free silicon. Wettability and reac-tivity determine the quality of the bond betvveen both materials and, therefore, greatly influence the final properties of the composite. In many instances the properties of a reinforced metal have been shovvn to provide a performance advantage over a monolithic metal, but the high cost of producing the composite has prohibited vvidespread commercial use. Liquid-metal processes have the potential to be more economical; hovever, the non-vvetting nature of many ceramics by molten aluminium, vvhich results in 1 Dr. Varužan M. KEVORKIJAN Lackova 139 2341 Limbuš. Slovenija poor ceramic/metal interfaces and incomplete infiltration, has been an obstacle. Melt infiltration is a popular technique for fabricating MMCs, as it allovv near-net shape fabrication of components and material vvith a high reinforcing phase content. The molten metal may penetrate the porous preforms either under the action of an external force (pressure casting3 and vacuum assisted Iiquid infiltration process4) or through a capillary pressure vvhich is created once the molten metal vvets the ceramic surface (pressureless infiltration5). Several pressure casting methods have been used for preparing MMCs. The operating principle of a hydro-static pressure infiltration device6 is to use pressurised gas to force molten metal into an evacuated die. Another pressure casting technique is relatively simple7: pre-heat-ing the particle aggregate in a special mould and then adding 3 MPa pressure to the molten metal poured on the particle aggregate so as to encourage penetration vvhich results in a metal-particle composite. Recently, a bottom mixing process has also been suggested, vvhere an evacuated packed bed in the bottom of a crucible is covered vvith a melt, and than stirrer shears the interface betvveen the particles and the melt, resulting in incorporation8. Different fabrication methods using vacuum tech-niques for cast-in-place hardfacing of casting vvere also described9. In these processes, aluminium poured into a sand mould is dravvn by vacuum into a porous layer of reinforcing phase (named - preform) placed on a wall of the mould cavity. A recent molten metal process is the Lanxide Corp. Primex™ pressureless infiltration process10 ". In this process a packed bed of ceramic powder is infiltrated by an Al-Mg alloy, without any applied pressure, in a nitrogen atmosphere. The resulting composite, vvhich has a packed bed density of about 55 vol.-%, can than be di-luted in the appropriate matrix alloy. Ceramic particles of SiC and AI2O3, vvith particle size as fine as about 1 pm have been infiltrated in this way, and at infiltration rates of up to the order of centimetres per minute under specific processing conditions. Processing details of the Primex™ route are proprietary, but it vvould appear to be a very competitive process for higher volume fraction composites. The Lanxide Corporation has made exten-sive efforts to protect this very valuable technology and has vvell over 100 U.S. patents and over 1500 foreign patents pending, vvith nearly 50 U.S. patents and over 100 foreign patents being issued or allovved by the mid-dle of 1989. In this vvork, the preparation of porous SiC preforms made by SiC particles, platelets and vvhiskers have been demonstrated. The surface of SiC preforms has been cov-ered by SiCb layer using cost effective oxidation in air. Chemically treated preforms vvere than pressurelless infiltrated by an Al-Mg alloy in nitrogen atmosphere. The conditions for spontaneous (as used, spontaneously means vvithout the aid of any externally applied pressure or vacuum), pressureless infiltration, vvhich include the use of a magnesium containing alloy and a nitrogenous atmosphere have been already vvell documented in literature, by the inventors12. Hovvever, the offered explana-tion is semi-empirical based on the vvell knovvn role of magnesium vvhich decreases the surface tension of a molten aluminium alloy. As stated by inventors12, this alone does not induce spontaneous infiltration, but a nitrogen atmosphere may cause a further reduction in the surface tension, thus promoting vvetting. Additionally, the reactivity of magnesium induces interfacial reactions vvith solid ceramic surfaces. These reactions typically are not sufficient to promote spontaneous vvetting, but again in combination vvith a nitrogen atmosphere they may change or be altered, thus allovving the observed infiltration. These results clearly demonstrated that the combination of magnesium in the alloy and a nitrogeneous atmosphere leads to the spontaneous infiltration of aluminium alloy into ceramic fillers. Hovvever, little in-formation is available on the effect of a nitrogen atmosphere on vvetting. Some authors13 found that vvhen fabri-cating aluminium alloy matrix composites via compocasting, the use of a nitrogen atmosphere and a bubble-degassing step vvith nitrogen yielded composites vvith much lovver porosity than those produced similarly vvith argon, but these results may not be assciated vvith enhanced vvetting. In the present paper a mathematical model of spontaneous infiltration of a porous ceramic preform has been suggested. The role of magnesium and nitrogenous atmosphere vvas quantitatively evaluated among the impor-tant processing parameters (porosity of the preform, the specific surface area, surface tension and the contact angle). Moreover, the influence of above listed parameters on the infiltration rate (expressed by the infiltration length as a function of time) has been also demonstrated. In this way, the optimal conditions for spontaneous and cost effective pressureless infiltration of porous SiC preforms by molten aluminium alloy vvere selected and ex-perimentaly confirmed. 2 Pressureless infiltration - theoretical considera-tions A. Capilllar}> Law Spontaneous infiltration of a liquid into a porous medium takes plače vvhen the liquid vvets the solid. Other-vvise, a minimum external pressure should be applied. This threshold pressure P (also called capillary pressure) is related to the contact angle 0 and the particle size through the so-called capillary law or Laplace equation: P = 6^ylv cos0 Vp/((l-Vp)D) (1) vvhere y\v is the liquid-vapor surface tension, X a factor vvhich depends on the geometry of the particles, D the mean diameter of the particles, and Vp the particulate volume fraction. Note that product (- yiv cos 0) is the vvork of immersion Wi defined as the change in the free energy on immersing the solid in the liquid. The vvork of immersion can be vvritten in terms of the threshold (or capillary) pressure through the follovving expression: W, = P(l-Vp)/SsPVp (2) vvhere Ss is the specific surface area (the surface area per unit mass of porous preform) and p is the density of the particulate. Unfortunately, the Laplace equation de-scribes the situation for a cylindrical tube, a very crude model for the types of porous media under considera-tion here. This model, for example, cannot be applied to irregularly shaped pores vvhere the effect of both pore geometry and netvvork cooperatively combine vvith contact angle hysteresis14. Hovvever, White15 derived a spe-cialized expression based on the Laplace equation relat-ing the pressure, P required to prevent capillary rise in porous media for vvhich the specific surface area Ss, solid density p, surface tension yiv, contact angle 0, and porosity a, are knovvn: P = (l-£) p Ssylv cos0 /e (3) B. Darcy's Law The flovv of an incompressible fluid through a porous medium is governed by Darcy's lavv16. For unidirectional flavv, and neglecting any effect of gravity, Darcy's lavv can be vvritten as v0 = - (k/fi) (dP/dx) (4) where v0 is the superficial velocity of the fluid (the ve-loeity of the fluid as measured by the volumetric flovv rate per unit cross sectional area vvhere the cross section is taken perpendicular to the average direction of flow), ji the viscosity of the liquid, dP/dx the pressure gradient at the infiltration front, and k the intrinsic permeability. It has been found empirically that the intrinsic perme-ability k of a porous medium is proportional to the square of the mean particulate diameter17 k = aD2 (5) where the constant a must be determined experimen-tally. The superficial velocity v0 can be related to the actual velocity in the porous medium (dx/dt) by means of the particulate volume fraction Vp: v0 = (1-Vp) dx/dt (6) Combining Eqs. (4) and (6) and integrating, the ex-pression for the infiltration length, L as a function of time and the pressure drop in the liquid metal #9P can be written as: L = [2ktAP/)d(l-V )]' On the other hand, for pressureless infiltration the pressure drop should be at least equal to the threshold (or capillary) pressure (Eq. 3). Under conditions of constant permeability and constant capillary pressure, Eqs. (3) and (7) can be combined to obtain the following relation-ship between infiltration length L, time t, and other proc-essing parameters: L = (1/e) [2ktWjS,p(l- e)/p.]"2 (8) Note that W; (work of immersion) is equal - yiv cosG. The Eq. (8) can be simplifted introducing that e"'V2kSsp(l- e) is the infiltrability of porous preform Q: L = Q [Wi/n]"2t1'2 (9) Again, it's important to note that Eq. (9) is valid under conditions of constant infiltrability and porosity of ceramic filler, constant work of immersion and, ftnally, constant viscosity of the melt, which is very difficult to obtain in practice. In spite of this considerable limitation, Eq. (9) can be successfully used in combination with Eq. (3) in order to designe the simple mathematical criterion for an early stage of pressureless infiltration of porous ceramic preform. Moreover, using this procedure, the parameters of pressureless infiltration can be selected to satisfied both proceesing requirements: spontaneous infiltration at ac-ceprtable infiltration rate. 3 Materials and experimental procedures Preparation of porous SiC preforms For the purpose of this study, three basic SiC mor-phologies - particles, platelets and vvhiskers in several size ranges (Table 1) vvere used for preforms preparation. Photomicrographs of used povvders are compared in Fig. 1. A diagram outlining the preform production process is shovvn in Fig. 2. Table 1: Characteristics of SiC phases used Particles Platelets Whiskers HSC 1200 SiC Platelets M-Grade SiCw Micro grits Millenium Advanced Superior Graphite Materials. Inc. Refractory Technologies, Inc. Chemistry: Stoichiometric SiC Stoichiometric SiC Stoichiometric SiC Crystalographic Primary phase Beta Primary phase Beta Primary phase Beta Structure: Diameter range 2-12 35-40 whisker length (Hm): 15-20 Thickness (um): / 3-5 1-2 Aspect ratio: / 8-10 10-12 Purity: 97-99 wt% SiC <1000 ppm of <1000 ppm of metallic impurities metallic impurities Particulate Content -100 5-10 5-10 (%)■■ Oxygen (%) by 1.0 0.68 1.1 Lečo Free Carbon(%): 1.0 0.01 0.53 Specific Gravity (g/cm3): 3.21 3.21 3.21 (7) Preform infiltration The experimental lay-up used in this vvork consisted of an aluminium alloy ingot, measuring about <|)50 x 30 mm, placed on the top of a porous ceramic preform. The filler material had a height that vvas great enough to pre-vent full infiltration under the process conditions (i.e. more-or-less infinite column of filler material). After processing, the amount of infiltration (distance from al-loy/filler interface) vvas measured, and the composite vvas sectioned and examined both macro- and micros-tructurally. The alloy/filler pairs vvere than placed into a controlled atmosphere furnace vvithin a refractory vessel (a 99.9% sintered alumina). The furnace vvas evacuated to -1 Pa at room temperature and back-filled vvith an ni-trogen-containing atmosphere until a positive flovv vvas obtained. Note that ali experiments vvere conducted under a slight positive pressure that vvas achieved by bub-bling the exit gas through a 25 mm column of oil. Fol-lovving the procedure developed in Lanxide, the furnace vvas ramped to temperature at a rate of 200°C/h, held at temperature for the specified time (e.g. at 800-1000°C for 10 to 24 h for full infiltration of the speciments) and allovved to cool to 675°C, at vvhich time the samples vvere removed from the furnace and cooled to room temperature. Various combinations of magnesium-containing aluminium alloys, silicon carbide porous preforms, nitro-gen-containing gases, and temperature/time conditions vvere employed to study the effect of various process variables on the infiltration kinetics. Because the infiltration of the porous preforms oc-curs in a nitrogenous atmosphere (at least about 10 volume percent nitrogen and the balance a non-oxidizing gas under the process conditions), aluminium nitride pre-cipitates may form vvithin the aluminium alloy matrix. Figure 1: Photomicrographs of used SiC morphologies: a) HSC 1200 Mierogrits Superior Graphite, b) SiC Platelets, Millenium Materials, Inc. and c) M-Grade SiC whiskers, Advanced Refractory Technologies, Inc. Slika 1: SEM fotografije SiC uporabljenih delcev: a) SiC prah -HSC 1200 Mierogrits Superior Graphite, b) SiC ploščice, Millenium Materials, Inc. in c) SiC vvhiskerji - M-Grade, Advanced Refractory Technologies, Inc. The per cent weight gain provides a measure of the amount of aluminium nitride that forms during processing. For comparison, the total conversion of pure aluminium to aluminium nitride produces a weight gain of 52%. Moreover, because this experimental arrangement produced a constant volume of composite in ali cases where full infiltration occured, the vveight gains of different experiments could be directly compared. 4 Results and discussion Preform preparation SiC preforms, containing vvhiskers, platelets or particles, vvere fabricated by vacuum casting in a variety of shapes and vvith a uniform microstructure. The character-istics of these preforms are listed in Table 2. Table 2: Characteristics of SiC- vvhiskers, platelets and particles preforms made by vacuum casting method CHARACTERISTIC PREFORM Particle's grade Platelefs grade Whisker's grade Average bulk density (g/cm3) 1-2.25 1-2.25 1-2.25 Preform diameter (cm) 3-10 3-10 3-10 Preform height (cm) 2-5 2-5 2-5 BET-Specific surface area (nr/g) 1.5-5.9 2.0-2.5 3.5-3.8 Porosity (vol%) 30-70 30-70 30-70 Infiltration experiments The critical process conditions for pressureless infiltration of porous SiC preforms vvith molten aluminium Ultrasonic homogenisation of raw materials (SiC phase + 5 wt% polyphenilene - corresponding to 4% cxcess carbon) in toluene i Evaporation of toluene vvith constant stining 4 _Addition of 1 wt% amorphous boron in the slip by stining_ _l_ Treatment of mould surfaces with an ammoniumalginate solution to enhance _mould release_ _i_ 1 Vacuum casting in plaster mould | _i_ L_Removal of preform from mould | _i_ f Diying of Preform at 75°C I _i_ | Pyrolysis of polyphenylene in Argon Flow ( 4h at 450"C ) | _i_ | Sintering of Porous Body at 1900 - 2000°C for 0.5 h in Argon ( 100 KPa) | _i_ Ojddation of Preform in Air at 900°C for 4 hours ____or Other Surface Treatment_ Figure 2: Preform fabrication process Slika 2: Proces pridobivanja poroznih predoblik Figure 3: Relationship betvveen magnesium content in an Al-lOSi-Mg alloy and iniltration distance (process conditions - 5 h dwell at 1175 K, nitrogen atmosphere vvith 4 vol% Ar) measured in SiC particle grade preform Slika 3: Odvisnost globine infiltracije od vsebnosti magnezija v Al-lOSi-Mg zlitini (eksperimentalni pogoji - 5 h pri 1175 K, atmosfera dušika s 4 vol% Ar) za porozne SiC predoblike pripravljene iz SiC prahu alloys vvere found to be: (i) the alloy composition, (ii) the atmosphere composition, (iii) the process temperature and time and (iv) the infiltrability of the preforms. The influenece of alloy composition (specifically the magnesium content) on infiltration distance is plotted in Fig. 3. The collected results are in agreement vvith data previously reported by Aghajanian et al.12 The nevv data also confirm the linear relationship betvveen magnesium content and amount of infiltration proposed by Aghajanian et al.12. The effect of nitrogen content of the atmosphere on the infiltration process vvas determined by conducting experiments in atmospheres ranging from 100% N2 to 100% Ar. It vvas found that no infiltration occured in 100% Ar, only partial infiltration occured in 10% N2+90% Ar and full infiltration occured vvhen the nitrogen content exceeded 20-30 vol%. As reported", at high percentages of N2, vvhere infiltration vvas rapid, little ni-tride formed, vvhereas in dilute atmospheres, vvhere infiltration vvas slovv, observable levels of A1N formed. In a similar fashion, the process temperature significantly af-fects the quantity of nitride that forms vvithin the aluminium alloy matrix. Figure 4 plots unit vveight gain versus process temperature for samples using alloy Al-10Si-3Mg, preforms made by SiC grit and process conditions of a 5 hour dvvell at temperature in 95% N2/4% Ar. Results also demonstrate that increased process temperature results in increased nitride formaton vvhichin-crease becames significant and nearly linear for temperatures higher than 1125 K. At a constant magnesium level and a fixed nitrogen content, several others process variables can affect the infiltration behaviuor. Fig. 5 plots the infiltration distance against temperature for otherwise constant process conditions. It is evident that infiltration increased in an approximatively linear manner vvith the temperature. Ad-ditionally, the data shovv that there is a threshold temperature required to initiate the pressureless infiltration for a given set of process parameters. Although limited, the data presented in Fig. 6 suggest that the threshold temperature changes vvith the process conditions (the preform infiltrability, the alloy composition and the nitrogen content in the processing atmosphere). One can also conclude that the process temperature affects the quantity of A1N that formed vvithin the aluminium alloy matrix. Fig. 7 plots the unit vveight gain against temperature for different SiC grade preforms. The results demonstrate that as the process temperature increases, the quantity of A1N that forms also increases. These results are in vvell agreement vvith reported data12 and confirm that the increase in A1N content is approxi-mately exponential over the temperature range investigated. 0.5 4 . O J-t-t--1-1-1-1-1-}-1 10 20 30 40 50 60 70 80 90 100 Nitrogen in atmosphere (%) Figure 4: Dependance of unit vveight gain on the content of nitrogen in N2/Ar atmosphere (Al-10Si-3Mg alloy, SiC particle grade preform, 5 h soakat 1075 K) Slika 4: Odvisnost povečanja teže vzorcev od vsebnosti dušika v N2/Ar atmosferi (zlitina Al-10Si-3Mg, predoblike pripravljene iz SiC prahu, 5 h pri 1075 K) eo j 70 i 10 Process temperature (K) Figure 5: Variation of infiltration distance vvith process temperature (Al-10Si-3Mg alloy, SiC particle grade preform, 5 h soak at temperature in a nitrogen atmosphere vvith 4 vol% Ar) Slika 5: Odvisnost globine infiltracije od temperature (zlitina Al-10Si-3Mg, predoblike pripravljene iz SiC prahu, 5 h pri delovni temperaturi v dušikovi atmosferi s 4 vol% Ar) Magnesium contant (wt%) 1400 -r 2 1000 -• 3 ; 800 -a E • Z 600 0 1 f 400 -200 ■• o -1-1-1-1-1-1-I 30 40 50 60 70 80 90 100 Nitrogan In itmoipKen (wt%) Figure 6a: Relationship between threshold temperature and infiltrability of porous SiC particle grade preforms (Al-10Si-3Mg alloy and a 5 h soak at temperature in nitrogen atmosphere with 4 vol% Ar) Slika 6a: Odvisnost temperature začetka infiltracije od infiltrabilnosti poroznih predoblik, pripravljenih iz SiC prahu (zlitina Al-10Si-3Mg, 5 h pri delovni temperaturi v dušikovi atmosferi s 4 vol% Ar) Figure 6b: Relationship between threshold temperature and magnesium content in Al-lOSi alloy (SiC particle grade preform and a 5 h soak at temperature in nitrogen atmosphere vvith 4 vol% Ar) Slika 6b: Odvisnost temperaturnega praga od vsebnosti magnezija v Al-1 OSi zlitini (predoblike pripravljene iz SiC prahu, 5 h pri delovni temperaturi v dušikovi atmosferi s 4 vol% Ar) Figure 6c: Relationship between threshold temperature and nitrogen in processing atmosphere (Al-10Si-3Mg alloy, SiC particle grade preforms and a 5 h soak time at 1125 K) Slika 6c: Odvisnost temperaturnega praga za porozne predoblike pripravljene iz SiC prahu, od vsebnosti dušika v delovni atmosferi (zlitina Al-10Si-3Mg, 5 h pri delovni temperaturi) Figure 7: Relationship between process temperature and A1N formation (unit vveight gain) in aluminium alloy matrix (obtained using alloy Al-10Si-3Mg, SiC particle grade preform and a 5 h soak at temperature in nitrogen atmosphere with 4 vol% Ar) Slika 7: Odvisnost temperature infiltracije od deleža nastalega A1N (izraženega kot povečanje teže analiziranih vzorcev) v Al zlitini (zlitina Al-10Si-3Mg, predoblike pripravljene iz SiC prahu, 5 h pri delovni temperaturi v dušikovi atmosferi s 4 vol% Ar) The effect of the infiltrability of porous preforms (see Eq. (9)) on the infiltration process vvas studied using preforms vvith different porosity and specific surface area. Note that the preform infiltrability, defined as a" 'a2kSsp(l-a), could be expressed as a function of specific surface area (Ss) and porosity (a) taking into ac-count Eq. (5): Q = const. e-'SsVp(l- e) (10) Fig. 8 plots the infiltration distance against preform infiltrability. The changes in the infiltrability of porous preforms vvere obtained by ranging their porosity and specific surface area. In order to meet these require-ments, the preforms vvere prepared using selected sintering conditions. The results demonstrate that ali experi-mental data fit vvell vvith the proposed process kinetics expressed by Eq. (9) for othervvise constant process conditions. Moreover, Eq. (9) seems to be valid for very different morphology of SiC particles. Hovvever, the Eq. (9) also, in some matter, presents a serious problem. There is a very complex correlation betvveen preform porosity and its real specific surface area. Usually, BET technique is used to determine Ss. It should be noted, hovvever, that a method based upon gas adsorption at the surface vvhose area is to be measured may not provide the right value to be inserted in Eq.(9). In fact, as reported14, the specific surface area relevant in the vvetting of particulates by aluminium could be much lovver than that given by the BET technique. 90 80 70 E S 60 0 | BO n 1 40 ■ I 30 20 10 0 80 70 10 0 E 50 E 1.5 2 2.6 Infiltrabilitv, Q 3 3.5 4 1.6 2 2.6 3 Infiltrabilitv. O 3.5 4 0 0.5 1 1 5 2 2.6 3 3.5 4 Infiltrabilitv, Q Figure 8: Dependence of infiltration distance on infiltrability of different porous preforras (Al-10Si-3Mg alloy, a 5 h soak at 1175 K a nitrogen atmosphere with 4 vol% Ar) for - a) SiC particle grade preform, b) SiC platelets grade preform, and c) SiC whisker grade preform Slika 8: Odvisnost globine infiltracije od infiltrabilnosti poroznih predoblik, pridobljenih iz: a) SiC prahu, b) SiC ploščic in c) SiC whiskerjev (zlitina Al-10Si-3Mg, 5 h pri 1175 K, dušikova atmosfera s 4 vol% Ar) 5 Concluding remarks A process for the production of porous SiC preforms consisting particles, platelets or whiskers is reported. It involves the vacuum casting of specialy prepared slip and sintering of green body to the porous specimen. Fol-lowing this procedure, vacuum čast preforms in a variety of sizes, vvith high dimensional and compositional repro-ducibility, and vvith uniform characteristics vvere fabri-cated. Porous preforms vvere sucessfully pressureless infil-trated using the Primex™ method originaly developed by Lanxide, Inc. The results presented in this article demonstrate that the combination of magnesium in the Al alloy, the nitro-geneous processing atmosphere (vvith at least 25 vol% N2) and several porous preform characteristics (specifi-cally the porosity and specific surface area) summarised in term preform infiltrability leads to the pressureless infiltration of molten aluminium alloy into ceramic filler. The collected data have confirmed that no infiltration occured vvithout the correct combination of above listed process variables. This means that the magnesium con-tent in Al alloy and the content of nitrogen in processing atmosphere should be combined vvith correctly designed porous preform characteristics. The results of present vvork demonstrate the infuence of preform porosity and its specific surface area on the infiltration length. The infiltration kinetics vvere shovvn to be strongly affected by preform infiltrability for other-vvise constant process conditions. In addition to kinetics, it vvas found that experimental data fit vvell the proposed expression for the infiltration length as a function of preform infiltrability, vvork of adhesion, viscosity of the melt and processing time. Moreover, the experimental results demonstated that the suggested equation is opera-tive for the different morphology of SiC particulate used in this vvork. Hovvever, the influence of preform surface composition, vvhich should affectsthe vvork of adhesion, and the viscosity of the melt is matter of the further ex-perimental efforts. 6 Acknowledgements The porous ceramic preforms have been infiltrated using the laboratory feasibilities of the partners on Brite Euram project BE96-3925 (proposal). The authors also thank to Slovene Ministry of Science and Technology, Impol from Slovenska Bistrica and Nova Kreditna Banka Maribor for their financial support. This research is the part of basic scientific project: "Preparation and Characterisation of Discontinuously Reinforced Al-SiC MMCs" fully financed by Slovene Ministry of Science and Technology. 7 References 1 D. J. Lloyd, Intern. Materials Reviews, 39, 1994, 1-23 2T. S. Srivatsan, I. A. Ibrahim, F. A. Mohamed, and E. J. Lavernia, J. Mater. Sci., 26, 1991, 5965-78 3 A. Mortnesen and T. Wong, Met. Trans. A, 21A, 1990, 2257 "F. Folgar, Ceram. Engng. Sci. Proc., 9, 1988, 579 5 J. T. Burke, M. K. Aghajanian and M. A. Rocazella, Proc. Int. SAMPE Symp., 34, 1989, 2440 fiJ. A. Comie, A. Mortensen and M. C. Flemings, Proc. of ICCM-VI (ed. F. L. Mathews, et al.), 1987, 2.297-2.319 7 S. Nagata, A. Kitahara, S. Akiyama, and H. Ueno, Trans. AFS. 93. 1985, 49-54 8 S. Caron and J. Masounave: in "Fabrication of particulates reinforced metal composites", (ed. J. Masounave and F. G. Hamel), Materials Park, OH, ASM International, 1990, 107-113 "K. G. Daviš and J. G. Magny„ Trans. AFS.. 89, 1981, 385-402 10 A. W. Urquhart. Mater. Sci. Eng., A144. 1991. 75 "M. K. Aghajanian, M. A. Rocazella, J. T. Burke and S. D. Keck, J. Mater. Sci., 26. 1991, 447-454 12 US Pat. 4 828 008 "J. W. McCoy, C. Jones and F. E. Wamer, Sampe Q„ 19, 1988, 2, 37 l4R. Sharma, Colloids Surf., 16, 1985. 87 15 L. R. White, J. Colloid Interface Sci., 90, 1982, 536 16 J. Bear: in Dinamics of Fluids in Porous Media, Dover Publication, New York, NY, 1988 LETNO KAZALO - INDEX KOVINE ZLITINE TEHNOLOGIJE, 30, 1996, 1-6 Kronološko kazalo Vodopivec Franc, J. Vižintin: Temperatura in temperaturni gradient med fretting preizkusom 1% C in 1.5% Cr jekla ........................................KZT,30,1996,1-2,009-014 Stok Boris, P. Koc: Računalniško podprta identifikacija temperaturne odvisnosti snovnih lastnosti ....................... .................................................KZT,30,1996,1-2,015-018 Zore Borut, L. Kosec: Spajkanje korozijsko obstojnih zlitin z zaščitni atmosferi...........KZT,30,1996,1-2,019-023 Gliha Vladimir, I. Rak, A. Pristavec: Mehanske in lomne lastnosti krhkih delov toplotno vplivanega področja večvarkovnega zvarnega spoja ......................................... .................................................KZT,30,1996,1-2,025-028 Cvelbar Andrej, P. Panjan, B. Navinšek, A. Zalar: Študij pojavov med toplotno obdelavo tankih plasti Ni/Si na osnovi sprotnih meritev električne upornosti ...................... .................................................KZT,30,1996,1-2,029-033 Košir Aleš, B. Sarler: Primerjava lastnosti numeričnih metod kontrolnih prostornin in robnih elementov z dvojno recipročnostjo....................KZT,30,1996,1-2,035-039 Požun Karol, J. Leskovšek, L. Koller, M. Mozetič: Vpliv nečistoč na kontaktno upornost električnih kontaktov ... .................................................KZT,30,1996,1-2,041-043 Požun Karol, B. Paradiž, J■ Leskovšek, L. Irmančnik-Belič: Vpliv naparevanja na odzivni čas senzorja relativne vlažnosti zraka ..................KZT,30,1996,1-2,045-048 Zorko Benjamin, M. Budnar: Merjenje globinske porazdelitve koncentracije vodika v materialih z metodo ERDA .....................................KZT,30,1996,1-2,049-051 Praček Borut: Študij začetne faze oksidacije Pb in zlitine In20Pb80 ................................KZT,30,1996,1-2,053-055 Žlebnik Tatjana, K. Vidmar: Termoforeza ....................... .................................................KZT,30,1996,1-2,057-058 Horvat Mojca, T. Marinovič, A. Sebenik: Vpliv sistema za zamreževanje na reološke lastnosti dinamično zamreženih zlitin PP/EPDM ...........KZT,30,1996,1-2,059-060 Mirčeva Aneta, T. 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Smolej, B. Šuštaršič: Izdelava kompozita SiC/Al-Fe po postopku hitrega strjevanja ..... ..................................................KZT,30,1996,3-4,263-266 Nardin Vladimir, R. Turk, I. Bizjak: Optimiranje števila mehanskih preskusov za določevanje preoblikovalnih lastnosti kovinskih materialov .......................................... ..................................................KZT,30,1996,3-4,267-270 Husič Suhreta: Kvantifikacija mikrostrukture polimernih kompozita analizom slike ......KZT,30,1996,3-4,271-273 Skitek Tanja, R. Cvelbar, M. Samarin, I. Emri: Sočasno merjenje Poissonovega in relaksacijskega modula viskoelastičnih materialov v odvisnosti od časa ...................... ..................................................KZT,30,1996,3-4,275-277 Brodar Maksimiljan, l. Emri: Dušilne lastnosti konstrukcijskih polimerov in kompozitov...................................... ..................................................KZT,30,1996,3-4,279-282 Kralj Aleš, I. Emri, N. W. Tschoegl: Adaptacija visko-elastičnega relaksometra ........KZT,30,1996,3-4,283-285 Barborič F., M. Zigon, F. Rovan: Vpliv vrste pospeševalca na lastnosti bromiranega epoksidnega preprega in laminata ..............................KZT,30,1996,3-4,287-290 Indof Janez, V. lvušič, D. Indof, A. Bejuk: Abrazijsko in erozijsko preskušanje polimernih materialov .................. ..................................................KZT,30,1996,3-4,291-294 Malič Barbara, /. Arčon, M. Kosec, A. Kodre, M. Hribar, M. Stuhec, R. Frahm: Študij alkoksidnih prekurzorjev keramike na osnovi PbZr3-PbTi03 ................................... ..................................................KZT.30.1996.3-4.295-297 Delalut Uroš, M. Kosec: Kristalizacija plasti (Pb,La)(Zr,Ti)03 na platinski plasti in na plasti svinčevega titanata ........................KZT,30,1996,3-4,299-301 Saje Boris, S. Spaič, M. Valant: Mikrostrukturne raziskave v binarnem sistemu Sm-Ti ........................................... ..................................................KZT,30,1996,3-4,303-305 Saje Boris, B. Reinsch, S. Kobe-Beseničar, D. Kolar, I. R. Harris: Nitriranje zlitine Sm2Fen modificirane s Ta...... ..................................................KZT.30.1996.3-4.307-309 Nemec Tomaž, J. J. Rant, V. Apih, B. Glumac: Spremljanje procesov transporta vlage v gradbenih materialih z nevtronsko radiografijo ..........KZT,30,1996,3-4,311-313 Selih Jana: Masni pretoki v betonu med sušenjem ........ ..................................................KZT,30,1996,3-4,315-319 Maček Jadran, B. Novosel, M. Marinšek, V. Francetič: Termična analiza cirkonijevih gelov ................................ .................................................KZT,30,1996,3-4,321-324 Zupan Klementina, J. Maček, B. Novosel: Vpliv temperaturnega režima na termični razkroj gelov za pripravo železo-oksidnih magnetnih materialov ............................ .................................................KZT,30,1996,3-4,325-327 Mozetič Miran, M. Kveder, M. Drobnič, A. Pregelj: Meritve stopnje disociiranosti vodika s katalitičnimi sondami ........................................KZT,30,1996,3-4,329-332 Nemanič Vincenc: Določevanje velikosti stične ploskve med kroglo in kovinsko folijo pri obremenitvi z atmosferskim tlakom .......................KZT,30,1996,3-4,333-337 Sušterič Zoran: Zakaj kavčuki tečejo nenevvtonsko? ..... .................................................KZT,30,1996,3-4,339-342 Milim Milorad: Characterization of Ultrathin Films by Surface Sensitive Methods .... KZT,30,1996,3-4,343-348 Binder Stojan: Razvojna dejavnost v steklarski industriji .................................................KZT.30.1996,3-4,349-351 Vojvodič Gvardjančič Jelena: Lomna varnost jeklenih konstrukcij po različnih merilih ....................................... .................................................KZT,30,1996,3-4,353-363 Mencinger Jure, B. Sarler: Vpliv procesnih parametrov na polkontinuirno ulivanje.....KZT,30,1996,3-4,365-368 Kumer Boris, R. Turk: Ekspertni sistemi in pravila kali-briranja....................................KZT,30,1996,3-4,369-373 Kejžar Rajko, M. Ogrizek: Strženske žice za reparaturno vzdrževanje .............................KZT,30,1996,3-4,375-378 Kejžar Rajko: Navarjanje posebnih Ni-zlitin na konstrukcijsko jeklo .............................KZT,30,1996,3-4,379-382 Kejžar Rajko, L. Kosec: Izdelava rezilnih orodij z navar-janjem .....................................KZT,30,1996,3-4,383-386 Medved Jože, J. Cevka, V^ Gontarev, P. Fajfar: Toplotne lastnosti eksotermno - izolacijskih materialov ................ .................................................KZT,30,1996,3-4,387-389 Kovačevič Mihaela, N. Vižintin: Karakterizacija ognje-vzdržnih materialov - Opredelitev temperature uporabnosti ........................................KZT,30,1996,3-4,391-393 Grum Janez, R. Šturm: Lasersko kaljenje s pretaljevanjem površinske plasti sive in nodularne litine ................. .................................................KZT,30,1996,3-4,395-398 Robič Roman, R. Turk, K Nardin: Vpliv silicija in ana-liznih odstopanj na tehnologijo valjanja dinamo trakov . .................................................KZT,30,1996,3-4,399-401 Lalovič Milisav, M. Bešič: Transient Heat Transfer Process During Heating of Steel .. KZT,30,1996,3-4,403-404 Kralj-Novak Metka, Z. Sušterič, -4. Mesec, M. Žumer: Karakterizacija kavčukov z napravo RPA 2000 .............. .................................................KZT,30,1996,3-4,405-408 Makarovič Matjaž, F. B. Damjanič: Konstruiranje in optimizacija izdelkov iz poliestrskih laminatov .................. ..................................................KZT.30.1996.3-4.409-412 Stadler Zmago: Material za zavorne obloge - jeklena vlakna da ali ne? ....................KZT,30,1996,3-4,413-415 Sventner Kosmos Alenka, L. 1. Belič, D. Sušnik: Sintranje grobozrnate korundne keramike ....................................... ..................................................KZT,30,1996,3-4,417-419 Jenko Monika, F. Vodopivec, F. Marinšek: Površinsko aktivirana rast zrn v neorientiranih elektro pločevinah ...... .....................................................KZT.30.1996.5.431-438 Torkar Matjaž, M. Leban, B. Rjazancev: Razvoj biokom-patibilnih implantatov ...............KZT,30,1996,5,439-441 Spruk Sonja, A. Rodič: Vpliv kemijske sestave jekel na lasersko toplotno obdelavo .......KZT,30,1996,5,443-447 Šuštar Tomaž, B. Ule, T. Rodič: Novi koncepti opisovanja lezenja in preostala življenjska doba kovinskih materialov .......................................KZT,30,1996,5,449-454 Ule Boris, J. Vojvodič-Gvardjančič, M. Lovrečič-Saražin, J. Banovec, F. Kržič, D. Beg, C. Primec: Preostala uporabnost kovičenega železniškega mostu........................... ...................................................... KZT,30,1996,5,455-462 Kejžar Rajko, L. Kosec, A. Lagoja: Perspektive navarjanja v orodjarstvu ....................KZT,30,1996,5,463-466 Gasperič Jože: Vakuumska impregnacija ........................ .....................................................KZT,30,1996,5,467-470 Mozetič Miran, M. Drobnič, S. Spruk, A. Pregelj: Porazdelitev atomarnega vodika vzdolž ravne cevi .................. .....................................................KZT,30,1996,5,471-475 Grabke Hans Jiirgen: Surface and Grain Boundary Segregation of Antimony and Tin - Effects on Steel Properties ..............................................KZT,30,1996,6,483-495 Viefhaus H.: Applications of Surface Analytical Tech-niques in Corrosion Research (Mainly High Temperature Corrosion) ..................................KZT,30,1996,6,497-507 Kainer K. U.: Aluminium and Magnesium Based Metal Matrix Composites ....................KZT,30,1996,6,509-516 Parilak L'udovlt: Microstructural Considerations Limit- ing the Mechanical Properties of HSLA Steel ................ .....................................................KZT,30,1996,6,517-520 Koroušič Blaženko, M. Stupnišek: Predicting of Reac-tions During Carburization and Decarburization of Steels in Controlled Atmospheres .......KZT,30,1996,6,521-526 Grozdanič Vladimir: Fusion of Low Carbon Steel Scrap in the Middle Carbon Steel Melt ..................................... .....................................................KZT,30,1996,6,527-530 Mast R., H. Viefhaus, M. Lucas, H. J. Grabke: Equilib-rium Grain Boudary Segregation of Antimony in Iron Base Alloys................................KZT,30,1996,6,531-537 Godec Matjaž, M. Jenko, R. Mast, F. Vodopivec, H. J. Grabke, H. Viefhaus: Sn Influence on the Recrystalliza- tion of Non-Oriented Electrical Sheet ............................. .....................................................KZT,30,1996,6,539-543 Kobe Beseničar Spomenka, L. Vehovai; B. Saje: Corrosion Resistance of NdDyFeB Basic Alloys ..................... ....................................................KZT,30,1996,6,545-549 Janovec Jožef, V. Magula, P. Sevc: Some Aspects of Im-purity Grain Boundary Segregation in Low Alloy Cr-Mo-V Steel .................................KZT,30,1996,6,551-555 Leskovšek Vojteh, D. Kmetic, B. Šuštaršič: Mechanical Properties of High Temperature Vacuum Brazed HSS on Structural Carbon Steel vvith Simultaneous Heat Treat-ment ............................................KZT,30,1996,6,557-564 Kevorkijan Varužan M.: Discontinuous Al-SiC Composites Formed by a Lovv Cost Chemically Activated Infiltration Technique ..........................KZT,30,1996,565-572 Avtorsko kazalo Anžel Ivan, L. Kosec, A. Križman: Disperzijsko utrjanje hitro strjene zlitine Cu-Zr ......KZT,30,1996,3-4,225-229 Arnšek Aleš, A. Cadež: Merjenje majhnih navorov v vakuumu .................................KZT,30,1996,1-2,155-157 Barborič F., M. Zigon, F. Rovan: Vpliv vrste pospeševalca na lastnosti bromiranega epoksidnega preprega in laminata ..................................KZT,30,1996,3-4,287-290 Binder Stojan: Razvojna dejavnost v steklarski industriji .................................................KZT,30,1996,3-4,349-351 Bizjak Milan, L. Kosec, A. Smolej, B. Suštaršič: Izdelava kompozita SiC/Al-Fe po postopku hitrega strjevanja ..... .................................................KZT,30,1996,3-4,263-266 Brecl Marko, T. Malavašič: Sinteza in opredelitev stran-skoverižnih tekočekristalnih poliuretanov z metoksia- zobenzensko mezogeno enoto .......................................... .................................................KZT,30,1996,1-2,083-085 Brodar Maksimiljan, I. Emri: Dušilne lastnosti konstrukcijskih polimerov in kompozitov ..................................... .................................................KZT,30,1996,3-4,279-282 Cvelbar Andrej, P. Panjan, B. Navinšek, A. Zalar: Študij pojavov med toplotno obdelavo tankih plasti Ni/Si na osnovi sprotnih meritev električne upornosti ...................... .................................................KZT,30,1996,1-2,029-033 Čop Aleš, E. Bricelj, F. Marinšek: Krhkost toplo valjanih trakov višje legiranih dinamo jekel.................................. .................................................KZT,30,1996,3-4,231-234 Delalut Uroš, M. Kosec: Kristalizacija plasti (Pb,La)(Zr,Ti)C>3 na platinski plasti in na plasti svinčevega titanata ........................KZT,30,1996,3-4,299-301 Dime Franc, B. Mušič: Usklajevanje metod za karakteri-zacijo arheoloških materialov z meritvami magnetne susceptibilnosti .......................KZT,30,1996,1-2,111-115 Friedrich Franc, M. Komac, D. Kolar: Priprava SijN4 keramike v visokotemperaturnem avtoklavu ................... ..................................................KZT,30,1996,1-2,103-106 Gasperič Jože: Vakuumska impregnacija ........................ .....................................................KZT,30,1996,5,467-470 Gliha Vladimir, L Rak, A. Pristavec: Mehanske in lomne lastnosti krhkih delov toplotno vplivanega področja večvarkovnega zvarnega spoja ......................................... ..................................................KZT,30,1996,1-2,025-028 Godec Boštjan, L. Vehovar: Mehanizem korozijske odpornosti nerjavne jeklene litine s povečano vsebnostjo Si ..................................................KZT,30,1996,3-4,195-199 Godec Matjaž, M. Jenko, R. Mast, F. Vodopivec, H. J. Grabke, H. Viefhaus: Sn Influence on the Recrystalliza- tion of Non-Oriented Electrical Sheet ............................. .....................................................KZT,30,1996,6,539-543 Grabke Hans Jiirgen: Surface and Grain Boundary Segregation of Antimony and Tin - Effects on Steel Properties ..............................................KZT,30,1996,6,483-495 Grašič Igor, A. Paulin, P. Južina, A. Pregelj: Visokonapetostni napajalnik za ionsko-getrsko črpalko .................... ..................................................KZT,30,1996,1-2,147-149 Grozdanič Vladimir: Fusion of Lovv Carbon Steel Scrap in the Middle Carbon Steel Melt...................................... .....................................................KZT,30,1996,6,527-530 Grum Janez, D. Zuljan: Analiza laserskega procesa rezanja na avstenitnem nerjavnem jeklu in ocenjevanje kvalitete reza...........................KZT,30,1996,3-4,235-240 Grum Janez, M. Kisin: Ocenjevanje integritete površin na osnovi spremembe mikrostrukturnih sestavin pri finem struženju .........................KZT,30,1996,3-4,241-244 Grum Janez, R. Sturm: Lasersko kaljenje s pretaljevan- jem površinske plasti sive in nodularne litine ................. ..................................................KZT,30,1996,3-4,395-398 Hornak Peter, J. Zrnik, F, Kovač: Texture Development of the Hot Rolled Transformer Sheet Steel ..................... ..................................................KZT,30,1996,3-4,201-203 Horvat Mojca, T. Marinovič, A. Šebenik: Vpliv sistema za zamreževanje na reološke lastnosti dinamično zamreženih zlitin PP/EPDM ..KZT,30,1996,1-2,059-060 Husič Suhreta: Kvantifikacija mikrostrukture polimernih kompozita analizom slike ......KZT,30,1996,3-4,271-273 Huskič Miroslav, A. Šebenik: Kemijska modifikacija PVC.........................................KZT,30,1996,1-2,065-067 Indof Janez, V. Ivušič, D. Indof, A. Bejuk: Abrazijsko in erozijsko preskušanje polimernih materialov .................. ..................................................KZT,30,1996,3-4,291-294 Janovec Jožef, V. Magula, P. Ševc: Some Aspects of Im-purity Grain Boundary Segregation in Low Alloy Cr-Mo-V Steel .................................KZT,30,1996,6,551-555 Jenko Monika, F. Vodopivec, F. Marinšek: Površinsko aktivirana rast zrn v neorientiranih elektro pločevinah ...... .....................................................KZT,30,1996,5,431-438 Kainer K. U.: Aluminium and Magnesium Based Metal Matrix Composites.....................KZT,30,1996,6,509-516 Kejžar Rajko, M. Ogrizek: Strženske žice za reparaturno vzdrževanje .............................KZT,30,1996,3-4,375-378 Kejžar Rajko: Navarjanje posebnih Ni-zlitin na konstrukcijsko jeklo .............................KZT,30,1996,3-4,379-382 Kejžar Rajko, L. Kosec: Izdelava rezilnih orodij z navarjanjem .....................................KZT,30,1996,3-4,383-386 Kejžar Rajko, L. Kosec, A. Lagoja: Perspektive navarjanja v orodjarstvu .....................KZT,30,1996,5,463-466 Kevorkijan Varužan M.: Discontinuous Al-SiC Composites Formed by a Low Cost Chemically Activated Infiltration Technique .......................KZT,30,1996,6,565-572 Kobe Beseničar Spomenka, L. Vehovar, B. Saje: Corrosion Resistance of NdDyFeB Basic Alloys ..................... .....................................................KZT,30,1996,6,545-549 Koroušič Blaženko, M. Stupnišek: Predicting of Reac-tions During Carburization and Decarburization of Steels in Controlled Atmospheres ........KZT,30,1996,6,521-526 Kosec Borut: Vpliv temperaturnega polja na jeklo plašča valja pri procesu kontinuirnega litja aluminijevih trakov .................................................KZT,30,1996,3-4,259-261 Košir Aleš, B. Šarler: Primerjava lastnosti numeričnih metod kontrolnih prostornin in robnih elementov z dvojno recipročnostjo....................KZT,30,1996,1-2,035-039 Kovačevič Mihaela, N. Vižintin: Karakterizacija ognje-vzdržnih materialov - Opredelitev temperature uporabnosti ........................................KZT,30,1996,3-4,391-393 Kralj Aleš, L Emri, N. W. Tschoegl: Adaptacija visko-elastičnega relaksometra........KZT,30,1996,3-4,283-285 Kralj-Novak Metka, Z. Šušterič, A. Mesec, M. Žumer: Karakterizacija kavčukov z napravo RPA 2000 .............. .................................................KZT,30,1996,3-4,405-408 Kranjc Andreja, Č. Stropnik: Določanje lastnosti tripsina, imobiliziranega na površino membrane iz celuloznega acetata ............................KZT,30,1996,1-2,079-081 Kumer Boris, R. Turk: Ekspertni sistemi in pravila kalibriranja ....................................KZT,30,1996,3-4,369-373 Lalovič Milisav, M. Bešič: Transient Heat Transfer Process During Heating of Steel .. KZT,30,1996,3-4,403-404 Leskovšek Vojteh, D. Kmetič, B. Šuštaršič: Mechanical Properties of High Temperature Vacuum Brazed HSS on Struetural Carbon Steel vvith Simultaneous Heat Treatment ............................................KZT,30,1996,6,557-564 Lipovšek Nataša, F. Vodopivec, M. Jenko, D. Steiner Petrovič, L. Kosec: Poprava in rekristalizacija legirane neorientirane elektro pločevine ........................................ ..................................................KZT,30,1996,3-4,251-254 Maček Jadran, B. Novosel, M. Marinšek, V. Francetič: Termična analiza cirkonijevih gelov ................................ ..................................................KZT,30,1996,3-4,321-324 Makarovič Matjaž, F. B. Damjanič: Konstruiranje in optimizacija izdelkov iz poliestrskih laminatov .................. ..................................................KZT,30,1996,3-4,409-412 Makovec Darko, Z. Samardžija, D. Kolar: Način vgradnje Ce v strukturo BaTi03 .....KZT,30,1996,1-2,117-119 Makovec-Črnilogar Vesna, /. Anžur, S. Orešnik, A. Gantar: Vpliv izbranih polimernih mastilnih sredstev na lastnosti usnja...............................KZT,30,1996,1-2,095-098 Makovec-Črnilogar Vesna, /. Anžur, S. Orešnik, A. Gantar: Opredelitev izbranih usnjarskih polimernih mastilnih sredstev .............................KZT,30,1996,1-2,163-166 Malič Barbara, I. Arčon, M. Kosec, A. Kodre, M. Hribar, M. Štuhec, R. Frahm: Študij alkoksidnih prekurzorjev keramike na osnovi PbZr3-PbTi03 ................................... ..................................................KZT,30,1996,3-4,295-297 Mast R., H. Viefhaus, M. Lucas, H. J. Grabke: Equilib-rium Grain Boudary Segregation of Antimony in Iron Base Alloys................................KZT,30,1996,6,531-537 Medved Jože, J. Cevka, V. Gontarev, P. Fajfar: Toplotne lastnosti eksotermno - izolacijskih materialov ................ ..................................................KZT,30,1996,3-4,387-389 Mencinger Jure, B. Šarler: Vpliv procesnih parametrov na polkontinuirno ulivanje .....KZT,30,1996,3-4,365-368 Milun Milorad: Characterization of Ultrathin Films by Surface Sensitive Methods .... KZT,30,1996,3-4,343-348 Mirčeva Aneta, T. Malavašič: Sinteza in karakterizacija sulfonatnih poliuretanskih ionomerov.............................. ..................................................KZT,30,1996,1-2,061-063 Mozetič Miran, M. Kveder, M. Drobnič, A. Pregelj: Meritve stopnje disociiranosti vodika s katalitičnimi sondami ........................................KZT,30,1996,3-4,329-332 Mozetič Miran, M. Drobnič, S. Spruk, A. Pregelj: Porazdelitev atomarnega vodika vzdolž ravne cevi .................. .....................................................KZT,30,1996,5,471-475 Nardin Vladimir, R. Turk, I. Bizjak: Optimiranje števila mehanskih preskusov za določevanje preoblikovalnih lastnosti kovinskih materialov .......................................... ..................................................KZT,30,1996,3-4,267-270 Nemanič Vincenc: Določevanje velikosti stične ploskve med kroglo in kovinsko folijo pri obremenitvi z atmosferskim tlakom .......................KZT,30,1996,3-4,333-337 Nemec Tomaž, J■ J■ Rant, V. Apih, B. Glumac: Spremljanje procesov transporta vlage v gradbenih materialih z nevtronsko radiografijo ..........KZT,30,1996,3-4,311-313 Parilak Uudovit: Microstructural Considerations Limiting the Mechanical Properties of HSLA Steel ................ .....................................................KZT,30,1996,6,517-520 Pesek Ladislav: Stable Crack Growth in Microalloyed Steel Sheets ............................KZT,30,1996,3-4,185-190 Petek Marko, B. Kaiserberger: Nestično merjenje raztezka pri trgalnih preskusih ... KZT,30,1996,1-2,137-138 Požun Karol, J. Leskovšek, L. Koller, M. Mozetič: Vpliv nečistoč na kontaktno upornost električnih kontaktov ... .................................................KZT,30,1996,1-2,041-043 Požun Karol, B. Paradiž, J• Leskovšek, L. Jrmančnik-Belič: Vpliv naparevanja na odzivni čas senzorja relativne vlažnosti zraka ..................KZT,30,1996,1-2,045-048 Praček Borut: Študij začetne faze oksidacije Pb in zlitine In20Pb80 ................................KZT,30,1996,1-2,053-055 Pregelj Andrej, M. Drab, J. Novak, M. Mozetič, A. Paulin: Merilni sistem za ugotaljanje sposobnosti ionskogetrske črpalke ............KZT,30,1996,1-2,151-153 Radonjič Gregor, V. Musil: Kompatibilizacija polipropi-lenskih mešanic ......................KZT,30,1996,1-2,075-078 Robič Roman, R. Turk, V. Nardin: Vpliv silicija in ana-liznih odstopanj na tehnologijo valjanja dinamo trakov . .................................................KZT,30,1996,3-4,399-401 Saje Boris, S. Spaič, M. Valant: Mikrostrukturne raziskave v binarnem sistemu Sm-Ti .......................................... .................................................KZT,30,1996,3-4,303-305 Saje Boris, B. Reinsch, S. Kobe-Beseničar, D. Kolar, L R. Harris: Nitriranje zlitine Sm2Fen modificirane s Ta ..... .................................................KZT,30,1996,3-4,307-309 Samardžija Zoran, S. Kobe-Beseničar: Študij defektov v komercialnih magnetih ALNICO z elektronsko mikroanalizo ....................................KZT,30,1996,1-2,107-110 Sešelj Andreja, J. Stražišar: Problematika določevanja velikosti delcev finih materialov ...................................... .................................................KZT,30,1996,1-2,139-141 Sivec Matjaž, M. Drab, M. Cerar, A. Pregelj: Razvoj male suhe vakuumske črpalke - kompresorja za območje 100 mbar do 4 bar .................KZT,30,1996,1-2,159-161 Skitek Tanja, R. Cvelbar, M. Samarin, L Emri: Sočasno merjenje Poissonovega in relaksacijskega modula visko- elastičnih materialov v odvisnosti od časa ...................... .................................................KZT,30,1996,3-4,275-277 Smolej Anton, P. Panzalovič, M. Jelen: Vpliv dodatkov Al-Ti-B in pogojev ulivanja na velikost kristalnih zrn aluminija.................................KZT,30,1996,3-4,255-258 Spruk Sonja, B. Praček, A. Rodič: Laserska toplotna obdelava površine orodnega jekla OCR 12 ......................... .................................................KZT,30,1996,1-2,143-146 Spruk Sonja, A. Rodič: Vpliv kemijske sestave jekel na lasersko toplotno obdelavo ........KZT,30,1996,5,443-447 Stadler Zmago: Material za zavorne obloge - jeklena vlakna da ali ne? ....................KZT,30.1996,3-4,413-415 Šarler Božidar: Numerični postopek za izračun temperaturnega polja brame pri kontinuiranem ulivanju jekla .... ..................................................KZT,30,1996,3-4,217-223 Selih Jana: Masni pretoki v betonu med sušenjem ........ ..................................................KZT,30,1996,3-4,315-319 Spaniček Durdica, Z. Smolčič: Odredivanje deformacij-skog ponašanja uz djelovanje medija na primjeru poliamida 6 .............................KZT,30,1996,1-2,069-070 Stok Boris, P. Koc: Računalniško podprta identifikacija temperaturne odvisnosti snovnih lastnosti ....................... ..................................................KZT,30,1996,1-2,015-018 Šuštar Tomaž, B. Ule, T. Rodič: Novi koncepti opisovanja lezenja in preostala življenjska doba kovinskih materialov ..................................KZT,30,1996,5,449-454 Suštaršič Borivoj, V. Kevorkijan, J. Lamut: Razvoj postopkov izdelave Al/SiC kompozitov ................................ ..................................................KZT,30,1996,3-4,209-216 Sušterič Zoran: Zakaj kavčuki tečejo nenewtonsko? ...... ..................................................KZT,30,1996,3-4,339-342 Sventner Kosmos Alenka, L. I. Belič, D. Sušnik: Sintranje grobozrnate korundne keramike ....................................... ..................................................KZT,30,1996,3-4,417-419 Torkar Matjaž, V. Leskovšek: Obdelava površine zlitine FeAl 12,5 z ionskim nitriranjem v pulzirajoči plazmi .... ..................................................KZT,30,1996,3-4,205-207 Torkar Matjaž, M. Leban, B. Rjazancev: Razvoj biokom-patibilnih implantatov ...............KZT,30,1996,5,439-441 Ulčnik-Krump Manica, T. Malavašič, B. Zerjal: Ter- modinamika mešanic polimerov v raztopini.................... ..................................................KZT,30,1996,1-2,087-089 Ule Boris, J. Vojvodič-Gvardjančič, M. Lovrečič-Saražin, J. Banovec, F. Kržič, D. Beg, C. Primec: Preostala uporabnost kovičenega železniškega mostu........................... ...................................................KZT,30,1996,5,455-462 Vasevska Trajanka: Izdelava žice iz zlitine AlMg5 za kovice in vijake ......................KZT,30,1996,1-2,131-135 Vehovar Leopold, B. Godec: Vpliv silicija na izboljšanje korozijske odpornosti jeklenih legiranih litin .................. ..................................................KZT,30,1996,3-4,245-250 Verko Nerina, C. Stropnik, K. Ribič, G. Jonsson: Lastnosti modificiranih in nemodificiranih membran iz polisulfona ..................................KZT,30,1996,1-2,099-102 Viefhaus H.: Applications of Surface Analytical Tech-niques in Corrosion Research (Mainly High Temperature Corrosion) ..................................KZT,30,1996,6,497-507 Vodopivec Frannc J. Vižintin: Temperatura in temperaturni gradient med fretting preizkusom 1% C in 1.5% Cr jekla.........................................KZT,30,1996,1-2,009-014 Vojvodič Gvardjančič Jelena: Preiskave horizontalne stabilne tlačne posode za skladiščenje utekočinjenega naftnega plina................................KZT,30,1996,1-2,121-124 Vojvodič Gvardjančič Jelena: Lomna varnost jeklenih konstrukcij po različnih merilih ....................................... .................................................KZT,30,1996,3-4,353-363 Zore Borut, L. Kosec: Spajkanje korozijsko obstojnih zlitin z zaščitni atmosferi...........KZT,30,1996,1-2,019-023 Zore Borut, L. Kosec: Kompozitni spajkani spoji........... .................................................KZT,30,1996,1-2,125-130 Zorko Benjamin, M. Budnar: Merjenje globinske porazdelitve koncentracije vodika v materialih z metodo ERDA .....................................KZT,30,1996,1-2,049-051 Zrnik Jožef, P. Hornak, P. Pinke, M. Zitnansky: Creep Fatigue Characteristics of Single Crystal Nickel Base Su-peralloy CMSX 3 ...................KZT,30,1996,3-4,179-183 Zupan Klementina, J. Maček, B. Novosel: Vpliv temperaturnega režima na termični razkroj gelov za pripravo železo-oksidnih magnetnih materialov ............................ ..................................................KZT,30,1996,3-4,325-327 Zupanič Franc, S. Spaič, A. Križman: Interakcija modifikacijske zlitine AlTi5Bl s talino zlitine AlCu6PbBi ...... ..................................................KZT,30,1996,3-4,191-194 Žagar Ema, M. Žigon, T. Malavašič: Lastnosti razredčenih raztopin poliuretanskih ionomerov ........................ ..................................................KZT,30,1996,1-2,091-094 Zigon Majda: Polimerni kompoziti ................................. ..................................................KZT,30,1996,1-2,071-074 Žlebnik Tatjana, K. Vidmar: Termoforeza ....................... ..................................................KZT,30,1996,1-2,057-05 8