YU ISSN 0372-8633 ŽELEZARSKI ZBORNI K VSEBINA Stran Iskra Kiberne- Brudar Božidar tika, Kranj STRJEVANJE JEKLA V KOKILI 137 Kmetič Dimitrij, J. Ž v o k e 1 j, V. Vodopivec, M. Jakupovič, B. R a 1 i č — Metalurški inštitut Ljubljana F. Mlakar, V. Tucič — Železarna Štore BELE KROMOVE LITINE LEGIRANE Z MOLIBDENOM ZA VALJE 151 Rodič Tomaž — Univerza Edvarda Kardelja v Ljubljani, F NT, VTOZD Montanistika D.R.J. Owen — Dept. of Civil Engi-neering, University of Wales, Swansea, U. K. OSNOVNI KONCEPT NUMERIČNE SIMULACIJE RADIALNEGA KOVANJA Kosec Ladislav — Univerza Edvarda Kardelja v Ljubljani, FNT, VTOZD Montanistika F. K o s e 1 — Univerza Edvarda Kardelja v Ljubljani, Fakulteta za strojništvo NASTANEK IN RAST UTRUJE-NOSTNE RAZPOKE V KOROZIJSKEM MEDIJU Ule Boris, F. Vodopivec, J. Žvokelj, M. Grašič — Metalurški inštitut Ljubljana L. Kosec — Univerza Edvarda Kardelja v Ljubljani, FNT, VTOZD Montanistika ZAPOZNELI LOM JEKLA Z VISOKO TRDNOSTJO DOKTORSKA IN MAGISTRSKA DELA 167 175 183 193 CONTENTS Page Brudar Božidar — Iskra Kiberne-tika, Kranj SOLIDIFICATION OF STEEL IN A MOULD 137 Kmetič Dimitrij, J. Žvokelj, V. Vodopivec, M. Jakupovič, B. R a 1 i č — Metalurški inštitut Ljubljana F. Mlakar, V. Tucič — Železarna Štore WHITE CHROMIUM ČAST IRONS FOR ROLLS, ALLOYED WITH MO-LYBDENUM 151 Rodič Tomaž — Univerza Edvarda Kardelja v Ljubljani, FNT, VTOZD Montanistika D. R. J. O w e n — Dept. of Civil Engi-neering, University of Wales, Svvansea, U. K. BASIC CONCEPTS OF NUMER1CAL SIMULATION OF A RADIAL FOR-GING PROCESS 167 Kosec Ladislav — Univerza Edvarda Kardelja v Ljubljani, FNT, VTOZD Montanistika F. K o s e 1 — Univerza Edvarda Kardelja v Ljubljani, Fakulteta za strojništvo OCCURRENCE AND GROWTH OF FATIGUE CRACKS IN CORROSION ENVIRONMENT 175 Ule Boris, F. Vodopivec, J. Žvokelj, M. Grašič — Metalurški inštitut Ljubljana L. Kosec — Univerza Edvarda Kardelja v Ljubljani, FNT, VTOZD Montanistika DELAYED FRACTURE OF HIGH-STRENGTH STEEL 183 PH. D. AND M. SC. THESES 193 LETO 21 št. 4 - 1987 ŽEZB BQ 21 (4) 137-196 (1987) IZDAJAJO ŽELEZARNE JESENICE, RAVNE, ŠTORE IN METALURŠKI INŠTITUT ŽELEZARSKI ZBORNIK Izdajajo skupno Železarne Jesenice, Ravne, Štore in Metalurški inštitut Ljubljana UREDNIŠTVO Glavni in odgovorni urednik: J. Arh Uredniški odbor: A. Kveder, J. Rodič, A. Paulin, F. Grešovnik, F. Mlakar, K. Kuzman, J. Jamar Tehnični urednik: J. Jamar Lektor: R. Razinger Prevodi: A. Paulin, N. Smajič (English), J. Arh (German), P. Berger (Russian) NASLOV UREDNIŠTVA: Železarski zbornik, SŽ-Železarna Jesenice, 64270 Jesenice, Yugoslavia TISK: TK Gorenjski tisk, Kranj IZDAJATELJSKI SVET Prof. Dr. M. Gabrovšek (Predsednik), Železarna Jesenice Dr. B. Brudar, Iskra, Kranj Prof. Dr. V. Čižman, Univerza v Ljubljani Prof. Dr. D. Drobnjak, Univerza v Beogradu Prof. Dr. B. Koroušič, Metalurški inštitut Ljubljana Prof. Dr. L. Kosec, Univerza v Ljubljani Prof. Dr. J. Krajcar, Metalurški inštitut Sisak Prof. Dr. A. Križman, Univerza v Mariboru Dr. K. Kuzman, Univerza v Ljubljani Dr. A. Kveder, Metalurški inštitut v Ljubljani Prof. Dr. A. Paulin, Univerza v Ljubljani Prof. Dr. Z. Pašalič, Železarna Zenica Prof. Dr. C. Pelhan, Univerza v Ljubljani Prof. Dr. V. Prosenc, Univerza v Ljubljani Prof. Dr. B. Sicherl, Univerza v Ljubljani Dr. N. Smajič, Metalurški inštitut v Ljubljani Prof. Dr. J. Sušnik, Zdravstveni dom Ravne Dr. L. Vehovar, Metalurški inštitut Ljubljana Prof. Dr. F. Vodopivec, Metalurški inštitut Ljubljana Published jointly by the Jesenice, Ravne and Štore Steelworks, and The Institute of Metallurgy Ljubljana EDITORIAL STAFF Editor: J. Arh Associate Editors: A. Kveder, J. Rodič, A. Paulin, F. Grešovnik, F. Mlakar, K. Kuzman, J. Jamar Production editor: J. Jamar Lector: R. Razinger Translations: A. Paulin, N. Smajič (English), J. Arh (German), P. Berger (Russian) EDITORIAL ADDRESS: Železarski zbornik, SŽ-Železarna Jesenice, 64270 Jesenice, Yugoslavia PRINT: TK Gorenjski tisk, Kranj EDITORIAL ADVISORY BOARD Prof. Dr. M. Gabrovšek, (Chairman), Železarna Jesenice Dr. B. Brudar, Iskra, Kranj Prof. Dr. V. Čižman, Univerza v Ljubljani Prof. Dr. D. Drobnjak, Univerza v Beogradu Prof. Dr. B. Koroušič, Metalurški inštitut Ljubljana Prof. Dr. L. Kosec, Univerza v Ljubljani Prof. Dr. J. Krajcar, Metalurški inštitut Sisak Prof. Dr. A. Križman, Univerza v Mariboru Dr. K. Kuzman, Univerza v Ljubljani Dr. A. Kveder, Metalurški inštitut v Ljubljani Prof. Dr. A. Paulin, Univerza v Ljubljani Prof. Dr. Z. Pašalič, Železarna Zenica Prof. Dr. C. Pelhan, Univerza v Ljubljani Prof. Dr. V. Prosenc, Univerza v Ljubljani Prof. Dr. B. Sicherl, Univerza v Ljubljani Dr. N. Smajič, Metalurški inštitut v Ljubljani Prof. Dr. J. Sušnik, Zdravstveni dom Ravne Dr. L. Vehovar, Metalurški inštitut Ljubljana Prof. Dr. F. Vodopivec, Metalurški inštitut Ljubljana Oproščeno plačila prometnega davka na podlagi mnenja Izvršnega sveta SRS — sekretariat za informacije št. 421-1/172 do 23. 1.1974 11229280 ZELEZARSKI ZBORNIK IZDAJAJO ŽELEZARNE JESENICE, RAVNE, ŠTORE IN METALURŠKI INŠTITUT LETO 21 LJUBLJANA DECEMBER 1987 Vsebina B. Brudar Strjevanje jekla v kokili UDK: 669.18:669.112.223:620.192.43 ASM/SLA: N21, M28h, E25n, D9p, 9-69 D. Kmetič, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M. Jakupovič, B. Ralič Bele kromove litine legirane z molibdenom za valje UDK: 669.15'26-194:669.14.018.255 ASM/SLA: M28, N8b, TSk, 5, Cr, W23k T. Rodič, D. R. J. Owen Osnovni koncept numerične simulacije radialnega kovanju UDK: 621.73.045:519.6 ASM/SLA: F22, Q24, 1-66, U4g, U4k L. Kosec, F. Kosel Nastanek in rast utrujenostne razpoke v korozijskem mediju UDK: 620.193.01 ASM/SLA: Rlh, Rle, R2j, Q26p B. Ule, F. Vodopivec, J. Žvokelj, M. Grašič, L. Kosec Zapozneli lom jekla z visoko trdnostjo UDK: 669.14.018.2:539.56:620.192.3 ASM/SLA: Q26s, SGBa, ST, 2-60, EGn, 3-66 Doktorska in magistrska dela Stran Contents B. Brudar 137 Solidification of Steel in a Mould UDK: 669.18:669.112.223:620.192.43 ASM/SLA: N21, M28h, E25n, D9p, 9-69 D. Kmetič, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M. Jakupovič, B. Ralič 151 White Chromium Čast Irons for Rolls, Alloyed with Mo-lybdenum UDK: 669.15'26-194:669.14.018.255 ASM/SLA: M28, N8b, TSk, 5, Cr, W23k T. Rodič, D. R. J. Owen Basic Concepts of Numerical Simulation of a Radial 167 Forging Process UDK: 621.73.045:519.6 ASM/SLA: F22, Q24, 1-66, U4g, U4k L. Kosec, F. Kosel Occurrence and Growth of Fatigue Cracks in Corrosion 175 Environment UDK: 620.193.01 ASM/SLA: Rlh, Rle, R2j, Q26p B. Ule, F. Vodopivec, J. Žvokelj, M. Grašič, L. Kosec 183 Delayed Fracture of High-strength Steel UDK: 669.14.018.2:539.56:620.192.3 ASM/SLA: Q26s, SGBa, ST, 2-60, EGn, 3-66 193 PH. D. and M. SC. Theses Page 137 151 167 175 183 193 no*«1 5 ŽELEZARSKI ZBORNIK IZDAJAJO ŽELEZARNE JESENICE, RAVNE, ŠTORE IN METALURŠKI INŠTITUT LETO 21 LJUBLJANA DECEMBER 1987 Strjevanje jekla v kokili Solidification of Steel in a Mould B. Brudar* UDK: 669.18:669.112.223:620.192.43 ASM/SLA: N21, M28h, E25n, D9p, 9-69 UVOD Proces strjevanja jekla v kokili je opisan že v najrazličnejši strokovni literaturi1-10 s področja metalurgije in matematične fizike. Pritem pa ne gre zgolj za samo opisovanje pojavov, ki nastopajo pri strjevanju. Tu mislimo predvsem na iz-ceje", notranje napetosti12-13 in lunker. Vedno večje poskusov, da te pojave razložimo in jih opišemo z enačbami matematične fizike. Pri računalniški simulaciji strjevanja jekla v kokili, pri kateri smo raziskovali vpliv eksotermnih plošč na obliko primarnega lunkerja, smo namreč dobili zanimive rezultate, ki so nam dali ideje za nadaljnje raziskovalno delo. Tako se nam je z rekonstrukcijo kokile OK 650 posrečilo, da smo bistveno vplivali na porazdelitev izcej v prerezu strjenega bloka. Uspelo nam je zmanjšati sekundarni lunker in prišli smo do novih spoznanj o samem procesu strjevanja — da je namreč mogoče vplivati tudi na strukturo strjenega bloka. V nadaljevanju je natančneje opisan sam poskus in dobljeni rezultati. Gre za novo gledanje na pojav strjevanja, ki je morda nekoliko neobičajno. Prvi poskusi pa so dali odlične rezultate in odpirajo se nove možnosti v prizadevanjih za izboljšanje kvalitete strjenih blokov tudi drugačnih oblik. VPLIV EKSOTERMNIH PLOŠČ NA VELIKOST PRIMARNEGA LUNKERJA Izolacijske eksotermne plošče, ki jih montiramo v zgornji del kokile, služijo za to, da ohranimo zgornji del taline čim dalj v tekočem stanju. Pri strjevanju se pa volumen zmanjša približno za 4 %. Želimo, da bi bilo to zniževanje gladine taline čimbolj počasno in na čim večjem prerezu. Tako bi dosegli najmanjšo globino primarnega lunkerja. Na tržišču se pojavljajo vedno nove kvalitete izolacijskih plošč z vedno boljšimi izolacijskimi sposobnostmi, pa tudi z novo ceno. S pomočjo poenostavljenega modela smo želeli oceniti, kako vplivajo te lastnosti na globino primarnega lunkerja. * ISKRA Kibernetika, Kranj INTRODUCTION The process of solidification of steel in a mould is de-scribed elsevvhere in various textbooks and in profes-sional literature1-10 from the metallurgical science and the mathematical physics. Usually it is not limited only to phenomenological descriptions of the effects observed with solidification. Here we mean especially the appear-ance of segregations11, the internal stresses1213 and the shrinkage holes. There are more and more efforts to ex-plain these effects and to describe them by the equa-tions of the mathematical phvsics. With the computer simulation of solidification of steel in a mould vvhere special exothermic plates were used to reduce the primary shrinkage hole, very interesting results vvere obtained that gave us new ideas for the fur-ther research vvork. By the reconstruction of the mould OK 650 we suc-ceeded in influencing the distribution of various segregations in the middle of the cross-section of the ingot and in reducing the secondary shrinkage hole. So we came to a new knovvledge — how the structure of the solidified ingot could be modified. In the due text the ex-periment itself is described thoroughly together with the most important results. The idea about our vision of the process of solidification is perhaps a little unusual. But our first experiments gave us exellent results and new possibilities in improving the quality of solid ingots of dif-ferent forms vvere opened. INFLUENCE OF EXOTHERMIC PLATES UPON THE PRIMARY SHRINKAGE HOLE The isolating exothermic plates that are usually mounted at the top of the mould are supposed to en-able the upper part of the liquid steel to stay liquid as long as possible. By the process of solidification the liquid shrinks for about 4 %. It is desired to lovver the le-vel of the liquid steel as slow as possible and to keep it in the largest possible cross-section. In this way the smallest depth of the primary shrinkage hole is obtained. On the market there are permanently nevv qualities of isolating plates available with better isolating properties and naturally with nevv prices. V ta namen smo izdelali matematični model, s katerim smo simulirali proces strjevanja, in pri tem izhajali iz naslednjih predpostavk: — kokila ima obliko pokončnega valja z enakomerno debelo steno, — kokila stoji na debeli livni plošči iz podobnega materiala, — velikost kokile, debelino stene in plošče lahko poljubno spreminjamo, prav tako tudi fizikalne lastnosti, kot sta specifična toplota in toplotna prevodnost strjenega bloka in kokile, — talina se pri vlivanju v kokilo dviga enakomerno z določeno hitrostjo, — gladina jeklene taline je idealno izolirana: toplotna prevodnost praška za posipanje je enaka 0, — temperatura taline je enaka temperaturi tausca, strjevati se začne, ko je kokila nalita do vrha, — latentna toplota ni vključena v specifično toploto, strjevaje v celoti poteče pri temperaturi tališča, — toplotni stik med talino in kokilo naj bo ves čas idealen: strjeni blok je ves čas v tesnem stiku s kokilo, — na zunanji steni kokile predpostavljamo ohlajanje s konvekcijo s konstantnim konvekcijskim koeficientom, — v zgornjem delu kokile imamo izolacijske plošče z znanimi lastnostmi, ki segajo od vrha kokile do določene globine. Za takšen poenostavljen primer smo izdelali računalniški program, s katerim smo lahko izračunali temperaturni profil v prerezu kokile in bloka in pri tem spremljali ugrezanje gladine in napredovanje meje med tekočo in trdno fazo. Variirali smo lastnosti materiala, iz katerega so izdelane eksotermne plošče, dimenzije plošč in hitrost uliva-nja. Poročilo o tej raziskavi je shranjeno v strokovni knjižnici Železarne Jesenice. Računalniški program smo uspešno uporabili tudi pri študiju strjevanja valjavniških valjev14, ki so bili uliti v železarni Štore, in se prepričali, kolikšna je upravičenost omenjenih predpostavk. Prišli smo še do naslednjih spoznanj: — če bi ulivali brez eksotermnih plošč, bi dobili zelo globok primarni lunker, ki bi segal skoraj do polovice višine bloka, — pri vsakem ulivanju z eksotermnimi ploščami pa opazimo pojav »mostu« iz strjenega jekla, ki nastane nekje na 3/4 višine bloka. V trenutku, ko nastane most, se pač en del taline nahaja pod njim, drugi del taline pa nad mostom. Delež taline nad mostom je tem večji, čim boljše izolacijske sposobnosti imajo eksotermne plošče in čim hitreje se dviga talina v kokili. Most nastane v vsakem primeru. Pojavlja pa se drugo vprašanje: Kaj se zgodi z mostom in kako se talina pod njim strjuje, če upoštevamo, da se pri nadaljnjem strjevanju volumen ujete taline zmanjša za 4 %. Odgovore na to vprašanje je mogoče najti v različnih člankih, ki govorijo o rahli sredini v bloku, o notranjih razpokah, luknjicah, različnih vrstah poroznosti15, sekundarnem lunkerju in podobno. Mnogi avtorji trdijo, da pri tem pride do ugrezanja strjenega dela mostu. Po naših izračunih in po natančnejšem ogledu Bau-mannovega odtisa prerezanega bloka iz avtomatnega jekla Č399016 smo ugotovili, da je verjetno res šlo za ugrezanje strjenega mostu oziroma za vdiranje nečistoč, ki se nabirajo v glavi, v notranjost bloka. Using the simplified mathematical model we wished to estimate how ali these different properties influence the depth of the primary shrinkage hole. For this purpose the mathematical model is made to simulate the process of solidification based on the follovving assumptions: — the mould has the form of a cylinder with a uniform thick vvall, — the mould is put upon a thick casting plate made of the same material, — the dimensions of the mould, the thickness of the vvall and of the casting plate can be varied together with the physical properties of the material like specific heat, the thermal conductivity of the solid ingot and of the mould, — the level of the liquid steel is isolated perfectly: the thermal conductivity of the isolating povvder on the top is equal zero, — the temperature of the liquid steel is equal to the melting point, the solidification starts to proceed in the moment when the mould gets completely filled, — the latent heat of the liquid steel is not included into the specific heat and the solidification is proceeded at the melting point, — the thermal contact betvveen the liquid and the mould is ideal ali the tirne, even the solid ingot stays at-tached to the vvall of the mould, there is no air gap, — at the outer surface of the mould the convective heat transfer with a definite coefticient of convection is assumed, — inside in the upper part of the mould the exother-mic isolating plates are mounted extending from the top of the mould to a certain depth into the liquid steel. For such a simplified čase the computer program is made for calculation of the temperature profile in the cross-section of the mould and in the ingot. It is possi-ble to follovv the lovvering of the liquid metal and to study the improving of the boundary betvveen the solid and the liquid phase. The properties of the isolating plates were varied together vvith the dimensions and vvith the casting speed. The detailed report of this research work could be obtained at Strokova knjižnica Železarne Jesenice. The computer program was also successfully applied to the study of solidification of the steel cylinder14 for the rolling mili, čast in Železarna Štore, and the validity of the suppositions mentioned above could be verified. The follovving conclusions were found: — casting vvithout exothermic plates vvould cause very deep primary shrinkage hole extending nearly to the half height of the ingot, — vvith any čase vvhere exothermic plates vvere used, the formation of a "bridge" of the solid steel could be observed at about 3/4 of the ingot height. In the moment of the formation of the bridge one part of the liquid steel remains above the bridge, the second part gets caught under the bridge. The amount of the liquid metal under the bridge is somehovv propor-tional to the isolating properties of the isolating plates and to the speed of raising of the liquid metal in the mould. The bridge is formed in any čase. Novv a new question arises. What happens to the bridge and how is the solidification improving when the reduction of the volume (4 %) of the caught liquid metal is taken into account. The answer to this question could be found in the ar-ticles describing the formation of the "soft middle" in the ingot, the internal cracks, the holes and various types of 500 600 700 900 1000 1100 1200 1300 1400 Slika 1: Izoterme v prerezu bloka v kokili. Toplotna prevodnost ekso-termnih plošč 1.563 W m-1 K-1, hitrost dviganja taline 200 mm/min, stanje po 106 minutah. Fig. 1: Isotherms in the cross-section of the ingot and the mould. Thermal conductivity of the exothermic plates 1.5 Wm~1 K~1, the speed of raising of the liquid steel 200 mm/min, the situa-tion after 106 minutes. Izračunani potek meje med trdno in tekočo fazo v simulaciji strjevanja bloka pa nas je napeljal še na novo idejo. Slika 1 prikazuje izoterme v trenutku, ko je pri omenjeni simulaciji nastal most. Ali se morda meja med trdno in tekočo fazo ne ujema s črtami, ki na odtisih po Baumannu običajno pomenijo izceje MnS v obliki črke A? Izredna podobnost v poteku linij izcej MnS v obliki črke A in izračunanih izoterm nas je silila v iskanje dokaza za tako predpostavko. Po nekaterih razlagah17 18 naj bi bila za to kriva ko-ničnost kokile. V svojih izračunih pa smo predpostavljali, da kokila ni konična. Kasneje smo našli poročilo, ki opisuje enake oblike izcej pri kokili, ki je celo širša v zgornjem delu'9 20. To je bil dokaz, da so razlage, ki se pojavljajo tudi v metalurških učbenikih, včasih zelo pomanjkljive. Če so izceje v obliki črke A zares slike trenutne meje med tekočo in trdno fazo v prerezu bloka, ki naj bi nastale ob strjevanju, bi bilo mogoče na to porazdelitev vplivati, če bi lahko vplivali na hitrost strjevanja. Z omenjenim računalniškim programom smo naredili simulacijo, pri kateri smo pogoje ohlajanja spreminjali. Tako smo si »izmislili« dodatno plast iz šamotne opeke, ki naj bi bila pritrjena na zunanji strani v zgornji polovici kokile. Izračunali smo, kako bi se spremenil potek strjevanja v takem primeru. Ugotovili smo, da bi to prav nič ne vplivalo na hitrost strjevanja in da bi bile razlike v temperaturni porazdelitvi v trenutku, ko bi nastal most, zanemarljivo majhne. Seveda nismo naredili nobenega praktičnega poskusa, saj ni bilo potrebno. Druga ideja je bila, da bi kokilo »postavili« na vodno hlajeno bakreno livno ploščo. Tudi ta simulacija je pokazala, da s tem ne bi prav nič vplivali na tvorbo mostu. V vsakem primeru je bila stena kokile predebela, da bi bilo mogoče kakorkoli vplivati na potek strjevanja v notranjosti oziroma na potek meje med trdno in tekočo fazo v trenutku, ko nastane omenjeni most. porosity15, secondary shrinkage holes and similar things. Many authors are suggesting the lovvering of the solidifi-ed bridge. According to our calculations and after a thorough examination of the sulphur prints in the cross-section of the ingot16, it was found out that in fact the lovvering of the bridge and the penetration of impurities from the top into the middle of the ingot could be assumed. The calculated course of the boundary betvveen the liquid and the solid phase in the simulation of solidifica-tion of a steel ingot led us to a new idea. Fig. 1 is represeting the isotherms in the moment of the formation of the bridge according to our simulation. Is it possible to assume that the boundary betvveen the solid and the liquid phase vvere equal to the lines that correspond to the segregations of MnS in the sulphur prints in the form of the letter A? The outstanding similarity in the course of the lines corresponding to segregations of MnS in the form of the letter A vvith the calculated isotherms forced us to look for the confirmation of our assumptions. According to some interpretations1718 about the pro-cess of solidification this effect could be explained sim-ply by the conicity of the mould. In our calculations, hovvever, it was assumed that the mould was not conical. Later we found the articles vvhere the same form of segregations vvere reported19-20 even in the moulds that vvere vvider at the top. It was the proof that explanations appearing in metallurgical text-books are sometimes too superficial. If the so called A-segregates are really corresponding to the boundaries betvveen the liquid and the solid phase in the cross-section of the ingot that should be formed at solidification it vvould be possible to influence them if we could influence the speed of solidification. With the computer program different simulations vvere made vvith different cooling conditions. An additional layer of recovery that vvould be "fixed" at the outer side of the upper half of the mould was si-mulated. It vvas calculated how the solidification vvould be changed in such čase. It vvas found out that it would Slika 2: Izoterme pri simuliranem primeru brez mosta. Toplotna prevodnost zunanje obloge 1.5 W ml| K-1, hitrost dviganja taline 200 mm/min, stanje po 109 minutah Fig. 2: Isotherms in the cross-section of the simulated čase. The ther-mal conductivity of the isolating layer outside 1.5 W m-1 K-1, the speed of raising of the liquid steel 200 mm/min, the situa-tion after 109 minutes. Ostala nam je torej samo še ena možnost: v zgornjem delu »stranjšati« steno kokile in jo še dodatno izolirati z zunanje strani. V omenjenem programu za simulacijo strjevanja smo postopoma »tanjšali« steno kokile in »dodajali« izolator toliko časa, da smo prišli do zaželenega rezultata. Rezultat j?a je bil takle: (slika 2). Jeklena talina se je strjevala tako, da se most sploh ni pojavil. Meje med tekočo in trdno fazo so dobile obliko črke U. V sredini je sicer prišlo do primarnega lunkerja, o sekundarnem lunkerju, ugrezanju mostu oziroma o vdiranju nečistoč iz glave v notranjost pa ni bilo sledu. Seveda je bilo vse to izračunano na matematičnem modelu. Model sam je temeljil na razmeroma hudih poenostavitvah21 glede stika med kokilo in talino, vendar pa nam je dal ideje za nadaljnje raziskovalno delo. Na osnovi rezultatov omenjene simulacije smo naredili praktični poskus s tako imenovano rekonstruirano kokilo. REKONSTRUIRANA KOKILA Odločili smo se za stanjšanje stene kokile tako, kot prikazuje slika 3. Zgornjo polovico smo konično posneli z zunanje strani in preostalo debelino nadomestili z izolatorjem. Potrebna je bila posebna konstrukcija, ki je omogočala stripanje in polnjenje zgornjega dela z izolacijskim sredstvom. Na ta način smo hoteli zmanjšati toplotno kapaciteto zgornjega dela kokile v primerjavi s spodnjim delom in dodatno zmanjšati hitrost strjevanja v glavi bloka. Zaradi primerjave rezultatov smo ulili en blok v klasično kokilo. Odločili smo se za format OK 650 (650 x 650 x 2000) in za jeklo, kvalitete Č3990. Pri tem jeklu je namreč mogoče opazovati izredno intenzivne izceje MnS in je zato tudi primerjava med različnimi bloki lažja. Napravili smo 4 poskuse: 1. jeklo ulito v klasično kokilo (A) 2. jeklo ulito v kokilo d = 40mm izolacija: livarski pesek (B) not influence the solidification at ali and that the temperature differences in the moment of the formation of the bridge vvould be negligibly small. The second idea vvas to "put" the mould upon a wa-ter-cooled casting plate made of copper. Also the re-sults of this simulation shovved that this vvould not influence the formation of the bridge at ali. In ali cases the wall thickness seemed to be too large to be able to influence the process of solidification inside, especially the course of the bundary betvveen the solid and the liquid phase in the moment of the formation of the bridge. There vvas only one possibility stili left. To lessen the thickness of the vvall in the upper part of the mould and to isolate it additionally from the outside. In the program for the simulation of solidification the thickness of the vvall vvas gradually "thinned" and an isolator vvas "added" till the final result vvas obtained. The final result vvas the follovving (Fig. 2). The liquid steel became solid in such a way that the bridge vvas not formed at ali. The boundaries betvveen the solid and the liquid phase got the form of the letter U. In the middle the primary shrinkage hole could be calculated, but there vvas no secondary shrinkage hole, no lovvering the bridge or penetrating impurities from the top into inside. Naturally ali this vvas calculated according to mathe-matical model. The model itself based on rather severe simplifica-tions concerning the thermal contact21 betvveen the mould and the liquid steel but it gave us anyway the ideas for the further research work. On basis of these simulations we performed the experiment vvith the so-called reconstructed mould. RECONSTRUCTED MOULD We decided to thin the mould vvall according to Fig. 3 The upper part vvas taken off conicaly and the re-maining thickness vvas replaced by the thermal isolator. notjcijstce ploitt Slika 3: Prerez rekonstruirane kokile: A izolator B kokila C livna plošča Fig. 3: The cross-section of the reconstructed mould: A isolator B mould C casting plate 3. jeklo ulito v kokilo d = 40 mm izolacija: perlit (C) 4. jeklo ulito v kokilo d = 20 mm izolacija: perlit (D). (V primeru 4 je bil perlit posebej sušen.) Na slikah 4, 5, 6 in 7 so prikazani odtisi po Bauman-nu za vse 4 primere obenem z ustrezno sliko jedkane površine prereza. POJASNILO K POSAMEZNIM SLIKAM Slika 4 je zelo podobna sliki iz leta 197316. Prav lepo se vidijo izceje v obliki črke A in v sredini izceje MnS v obliki črke V. Na fotografiji jedkane površine se zelo lepo vidijo vzdolžne razpoke, ki potekajo praktično po vsej dolžini bloka. Blok B na sliki 5 je bil pa ulit v rekonstruirano kokilo. Očitno je število razpok v področju sekundarnega lunkerja bistveno manjše, potek izcej v obliki črke A pa ni bistveno drugačen, kot pri bloku A. Ta ugotovitev nas je v prvem trenutku nekoliko razočarala, saj smo pričakovali znatnejše razlike v poteku teh izcej. Pojav smo si razložili s tem, da so izolacijske sposobnosti livarskega peska verjetno razmeroma majhne. Pri bloku C smo namesto livarskega peska uporabili perlit (U2), ki se uporablja za izolacijo fasad na zgradbah. Pokazalo se je, da smo vendarle na pravi poti. S slike 6 se jasno vidi, da je izcej v obliki črke V v sredini prereza manj in praktično tudi ni več razpok v sredini prereza v področju sekundarnega lunkerja. Glava je nekoliko bolj čista, robne izceje v obliki črke A potekajo nekoliko bolj pokončno, kot pri bloku A, nečistoče v obliki črke V v sredini prereza so manj izrazite. Opazili pa smo nekaj, kar nam prej ni zbudilo pozornosti. Osrednji del v glavi, v katerem sicer vidimo mnogo izcej V, (primerjaj blok A!), je tu najširši, če med seboj primerjamo prereze blokov A, B in C. To področje je omejeno nekako z dvema paralelnima navpičnima črtama. A special construction vvas necessary to enable the stripping of the ingot and the filling of the isolating material. Tha thermal capacity of the upper part of the mould in comparison with the bottom part was reduced and we hoped that the speed of solidification in the top vvould be additionally reduced too. To make the comparison among different experiments easier one ingot was čast into the ordinary mould. For our practical experiment the mould OK 650 (65 x 650 x 2000) and the quality of the free-cutting steel C 3990 were chosen. With this type of steel very intense segregates of MnS could be detected so that the comparison among different ingots is easier. The follovving experiments were performed: — steel čast into an ordinary mould (A) — steel čast into the reconstructed mould with D = 40 mm and vvith the casting sand used as the isolating material (B) — steel čast into the reconstructed mould vvith D = 40 mm and vvith the pearlite used for the thermal iso-lation (C) — steel čast into the reconstructed mould vvith D = 20 mm and vvith the pearlite used for the thermal iso-lation (D) In the last čase the pearlite vvas specially dried. The Figs. 4, 5, 6 and 7 show the sulphur prints for ali the four cases together with the corresponding photos of the etched surfaces of the cross-sections. COMMENTS TO THE FIGURES Fig. 4 is very similar to the figure from the year 19731B. The segregations in the from of the latter A are clearly seen and the V-segregations of MnS in the middle of the cross-section are evident. From the photos of the etched surface the longitudinal cracks are shovvn, being practically distributed ali along the length of the ingot. The ingot B from Fig. 5 vvas čast into the reconstructed mould. It is evident that the number of cracks in the region of the scondary shrinkage hole is significantly smaller and the course of the segregates in the form of the letter A is not much different from that one of the ingot A. This effect disappointed us at the first moment since considerable differences in the course of these segregates were expected. This could be explained by the fact that the isolating properties of the casting sand were not good enough. With the casting of the ingot C instead of casting sand the pearlite (U2) vvas used. This is the same material that is so often used for the thermal isolation of fa-pades of houses. It vvas shovvn at once that we were nevertheless on the right way. From Fig. 6 it can be easily seen that there are not so many V-segregates just in the middle of the cross-section and that there are not so many cracks in the region of the secondary shrinkage hole. The top of the ingot is a little cleaner, the A-segregates are a little more steep than they are in the čase of the block A. The impu-rities in the form of the letter V are less expressed. But we observed something vvhat vvas not evident at the first moment. The middle top region (Fig. 6), vvhere there are normally many V-segregates, is the largest if the ingots A, B and C are compared. This region is somehovv limited by two vertical parallel lines of segregates. We vvished to increase the isolating properties of the upper part of the mould. So the thickness of the wall Slika 4: Ingot A — odtis po Baumannu sulphur print Fig. 4: Ingot A — jedkana površina prereza the etched surface of the cross-section mmkitim^ JppIptNllJii: ■"> Slika 5: Ingot B — odtis po Baumannu sulphur print Fig. 5: Ingot B — jedkana površina prereza the etched surface of the cross-section Slika 6: Fig. 6: Ingot C — odtis po Baumannu Ingot C — jedkana površina prereza sulphur print the etched surface of the cross-section Slika 7: Rg 7. Ingot D - odtis po Baumannu Ingot D - jedkana površina prereza sulphur print the etched surface of the cross-section Hoteli smo še bolj povečati izolacijske sposobnosti v zgornji polovici kokile, zato smo dodatno stanjšali steno kokile in spet uporabili perlit. V tem primeru smo namenoma šli v skrajnost glede debeline stene kokile. Po pričakovanju se je ta stena zelo močno ogrela, ponekod celo nad tališče perlita, zaradi česar smo morali perlit med samim strjevanjem dodajati. (Končno nam ga je celo nekoliko zmanjkalo). Stena kokile pa se je v zgornjem stanjšanem delu toliko ogrela, da se je nekoliko usločila (napihnila se je) in smo jo morali končno razrezati, da smo dobili blok iz kokile. Rezultat na sliki 7 pa je tisto, kar smo pričakovali. V tem primeru je področje homogene strukture med obema paralelno potekajočima izcejama še širše, kot v primeru C. Če gledamo od roba proti sredini prereza, potem končne izceje, ki so še jasno izražene, nimajo več oblike črke A, ampak opazujemo namesto konice pri A dvoje paralelnih navpičnih črt. Izceje bi laže opisali z dvema polovicama na glavo postavljene črke Y, ki sta nekoliko razmaknjeni. Razlika med slikama 4 in 7 je očitna. Nekoliko pa nas je vseeno razočaral »V« v sredini prereza, saj ga v tem primeru nismo pričakovali. Pozneje smo našli razlago tudi za ta pojav. NAŠA HIPOTEZA STRJEVANJA JEKLA V KOKILI Na osnovi rezultatov predhodnih računalniških obdelav in proučevanja slik 4—7 smo postavili naslednjo hipotezo o poteku strjevanja: Problem prehoda toplote med talino in kokilo opisuje več avtorjev. Nekateri govorijo o tanki plasti strjenega jekla, ki se tvori ob stiku taline s hladno kokilo. Ker se strjena »srajčka« skrči, odstopi od kokile. To pa povzroči, da se toplotni tok iz taline zmanjša in zato se »srajčka« ponovno pretali itd. To naj bi se periodično ponavljalo toliko časa, dokler se ne bi »srajčka« toliko zdebelila, da se ne bi več pretalila in bi že vzdržala hi-drostatični tlak. Težko verjamemo, da bi prišlo do take periodične tvorbe »srajčke«, ki se enkrat dotika kokile, drugič pa spet ne. Gre za neke vrste neidealen kontakt. Po naše narašča koeficient prenosa toplote od zgoraj navzdol zaradi večjega ferostatičnega tlaka. Iz taline torej odteka toplota v stene kokile, tako da se stene intenzivneje ogrevajo spodaj kot pa zgoraj. Talina se pri tem meša zaradi temperaturnih razlik. To pomeni, da se, po našem mnenju, vsa toplota, kije shranjena v talini zaradi pregretja, odteče v stene kokile in jih neenakomerno ogreje (spodaj bolj, zgoraj manj), tako da temperatura stene kokile približno linearno narašča od zgoraj navzdol. Pričakujemo, da ni bistvenih razlik v debelini »srajčke« zgoraj in spodaj v trenutku, ko ingot odstopi od stene kokile v celoti. Zgoraj se sicer zelo hitro naredi, vendar je tanka, saj je tudi ferostatični tlak manjši. Spodaj pa mora biti debelejša, saj blok kasneje odstopi od stene. Ko blok po vsej višini odstopi od stene kokile, pa zaradi nadaljnjega odtekanja toplote narašča debelina strjene »srajčke«. To pa poteka v vsakem primeru hitreje v zgornjem delu kot pa v spodnjem. Če namreč pomislimo, kolikšna je temperatura onstran zračne reže v steni kokile v zgornjem delu, je takoj jasno, da mora potekati strjevanje zgoraj hitreje. Spodnji deli kokile so se le precej bolj ogreli v tistem času, ko je bilo odtekanje toplote v steno kokile zaradi nastajajoče »srajčke« intenzivnejše. Zato trdimo, da v nada- was additionally reduced and the pearlite was used for the thermal isolation again. In this čase we deliberately decided for an extreme situation concerning the thick-ness of the vvall. As expected the vvall became very hot and the temperature raised somevvhere even above the melting point of the pearlite so that pearlite had to be added continously during the process of solidification. (Before the end of the experiment practically ali the isolating material available vvas used). Due to rather high temperatuers the vvall of the mould became deformed in the upper thinner region so that the ingot had to be cut out of it. The results shovvn in Fig. 7 are something what vvas expected before. The region of the homogenous structure betvveen both parallel flovving segregates is stili larger than it is in the čase C. Looking tovvards the center the final stili re-cognizable A-segregates are no more of the form of the letter A. Instead of the top of A, two parallel vertical lines could be observed. The segregations could be better described by two separated halves of the inverted letter Y. The differences betvveen the Figs. 4 and 7 are evi-dent. Hovvever the "V" in the middle of the cross-section in such an extreme čase vvas not expected any more. An explanation also for this effect vvas found later. OUR HYPOTHESIS OF SOLIDIFICATION OF STEEL IN A MOULD On the basis of the results of the computer simula-tions made before and from the careful examination of Figs. 4 to 7 the follovving hypothesis seems to be valid for the process of solidification. The problem of the heat transfer betvveen the mould and the liquid metal was described by several authors. Some of them suggest the formation of a very thin layer of solid steel (shell) formed at the contact of the liquid steel vvith the cold mould. Because the solid shell shrinks, an air gap is formed and it prevents the further heat flux from the inside of the liquid pool. The shell melts again and this process is supposed to be continu-ed so long until the shell gets so thick that is does not melt any more and until it can endure the ferrostatic pressure of the liquid steel inside. It can be hardly believed in the periodic formation of the shell that once sticks to the mould and then melts again. We mean that there must be some kind of non-ideal thermal contact. According to our vision we have to do vvith an increasing coefficient of the heat transfer vvhen looking from the top to the bottom of the mould due to increasing ferrostatic pressure. From the liquid metal the heat flux flovvs into the mould so that the vvalls become more intensively heated at the bottom than at the top. We suppose that the liquid is mixing ali the tirne due to temperature differences. It means that according to our idea ali the superheat that is stored in the liquid metal, flovvs to the vvalls of the mould and vvarms them non-uniformly (more at the bottom and less at the top) so that the temperature of the mould vvall increases ap-proximately linearly from the top to the bottom. We expect that there are not great differences betvveen the thickness of the shell formed at the bottom and at the top in the moment of the formation of the air gap ali along the length of the ingot. At the top a thin shell is formed very quickly together vvith the air gap because of very small ferrostatic pressure. At the bottom ljevanju strjevanja debelina stene narašča linearno od spodaj navzgor. Tudi izceje v obliki črke A so v področjih bliže steni kokile praktično ravne črte, ki so enako strme v vseh štirih primerih. To si razlagamo tako, ker mislimo, da v začetku strjevanja različna debelina stene v posameznih področjih kokile še ne pride do izraza. Kasneje pa se slika spremeni. Zmanjšana toplotna kapaciteta in povečana izolacija v primerih blokov B, C in D lahko znatno vplivata šele proti koncu strjevanja. Pri rekonstruirani kokili dobi torej talina proti koncu strjevanja obliko prisekanega stožca, ki se nadaljuje v valj. Širina tega valja se veča od slike 4 proti sliki 7. Kljub temu, da pri sliki 7 nismo pričakovali mostu, pa je vseeno mogoče videti, da se je del izcej v obliki črke A v sredini ugreznil v obliki črke V, čeprav ne posebno globoko (manj, kot je pa to razvidno s slike 4). Tudi za to smo našli razlago, ki je opisana v nadaljevanju. KAKO SI ZAMIŠLJAMO NASTANEK IZCEJ MnS Ob fronti dendritov, ki rastejo pravokotno na mejo med tekočim in trdnim, se bogati talina z vsebnostjo MnS, tako da ni mogoče več »tolerirati« tolikšne koncentracije. Čisti kristali potiskajo pred seboj talino, ki postane prenasičena. Zato naenkrat pride do izločanja MnS v obliki izcej, kar pa verjetno sprosti nekaj toplote. To pomeni, da se v nadaljevanju prodiranje dendritov proti sredini nekoliko zaustavi, saj odtekanje te reakcijske toplote ne povzroči rasti strjene plasti. V tem času ima tisti del taline ob strjeni steni možnost, da se v njem ponovno izenači koncentracija MnS. Pri tem ima odločilno vlogo temperatura oziroma konvekcijski in difuzijski procesi. Ko reakcijska toplota odteče, se vse skupaj ponovi. Tako si razlagamo nastanek izcej MnS v obliki črke A — nastanejo paralelne črte, pri katerih se medsebojna razdalja manjša, če gremo od roba proti sredini prereza. Omenjena oblika »taline« v obliki prisekanega stožca, ki se nadaljuje v valj, pa po našem mnenju dobi lastnosti težko se premikajoče »marmelade«, ki se pri nadaljnjem odvajanju toplote pretvori v plastično maso, podobno pudingu in se krči kot celota. Predstavljamo si, da je ta plastična talina nekako »obešena« na stene, ki so se že prej strdile. Ko se sama skrči, potegne za seboj navzdol tudi dele stene, ki so že prej nastali, pa še niso dovolj trdni. Nastnejo izceje v obliki črke V. Na to idejo nas je navedlo dejstvo, ki smo ga lahko opazovali pri vseh odtisih po Baumannu na slikah 4 do 7, da so namreč praktično pri vseh blokih tudi 0.5 metra nad osnovno ploskvijo bloka kristali deformirani, zavihani navzdol, podobno kot se to vidi v področju izcej V. V primeru D, ko imamo nad prisekanim stožcem izrazito širok valj, pa se kljub temu nismo mogli izogniti izcejam v obliki črke V v sredini prereza. Na jedkanem obrusu pa se vidi, da je v tistem področju glave, ki je po navadi kritičen, v primeru D izredno homogena struktura. Trdimo, da smo v zgornjem delu glave dosegli, da je material izredno homogen in čist in zanj ni mogoče več trditi, da je prišlo pri krčenju do vdiranja nečistoč iz glave v sredino prereza. Edino izcejo v obliki črke V v sredini pri bloku D pa lahko razložimo takole: the thickness is larger but the air gap is formed much later. After the air gap is formed completely, the further heat flux into the mould causes the increasing of the thickness of the solid shell. This process is improving more quickly in the upper part than in the lovver part of the ingot. If we just consider the temperature of the mould wall across the air gap in the upper part, it becomes quite evident that the solidification must proceed more quickly there. The bottom of the mould accepted much more heat, because the heat flux into the wall due to higher ferrostatic pressure vvith the formation of shell, vvas more intense. We say that after the air gap is formed in the further process of solidification the thickness of the shell increases linearly from the bottom to the top of the ingot. Even the A-segregations in the region closer to the vvall are practically straight lines vvith nearly the same steepness in ali four cases. This could be explained by the fact that at the beginning of solidification the influence of the different thick vvalls in different heights vvith the reconstructed mould can not play its role yet. In the cases of ingots B, C and D the influence of the smaller thermal capacity of the vvall and the increasing thermal isolation of the mould become considerable only at the end of the solidification process. With the reconstructed mould tovvards the end of solidification the liquid gets the form of a truncated cone that is continuing into a cylinder. The diameter of this cylinder is increasing if Figs. 4 to 7 are compared. Although vvith the ingot D (Fig. 7) no bridge is ex-pected, it is possible to see one part of an A-segregate in the middle of the cross-section is bent down in the form of the letter V, but not so deeply as it could be seen in Fig. 4. The explanation to this effect is also found and it is given later. HOW WE IMAGINE THE FORMATION OF THE SEGREGATIONS OF MnS With the grovvth of the dendritic crystals rectangular-ly to the boundary betvveen the liquid and the solid phase, the liquid gets richer on MnS so that once it is not possible to tolerate such concentration any more. Pure crystals are pushing the oversaturated liquid in front of them. At a certain moment the MnS starts to segregate and it causes most probably the generation of a small amount of heat. This vvould mean that in due course the grovvth of dendrites tovvards the center stops a little, because the reaction heat that flovvs into the mould does not cause further grovvth of crystals. In this moment in the liquid region close to the solid shell the concentration of MnS equalizes again. Different diffusion and convective processes due to temperature differences are playing the most important role. When the flovv of the reaction heat is finished, the vvhole story is repeated. So vve are explaining the formation of the A-segre-gates of MnS. The parallel lines can be observed and their mutual distance is decreasing when moving tovvards the center of the cross-section. The form of the li-quid metal mentioned above in the form of a truncated cone continuing into a cylinder attains according to our assumptions the property of a very slowly moving "jam" that vvith the further cooling gets the properties of a IZCEJE V OBLIKI ČRKE V Mislimo si, da ima talina obliko valja in da so stene tega valja trdne, narejene iz iste snovi. Ker se pri strjevanju zmanjšuje volumen, se pač lahko ugreza le sredina. Predpostavljajmo, da valj razrežemo enakomerno na valjaste plasti po sliki 8. Volumen vsake od teh plasti se pri strjevanju zmanjša za 4 % in zgornja ploskev naj se ugrezne tako, da dobi obliko navzdol obrnjenega stožca. Na sliki 8 so s črtkano črto označene valjaste plasti taline, s polnimi črtami pa plasti, ki bi jih dobili po strjevanju. Takšno sliko bi dobili, če bi želeli, da ostane talina nekako »privezana« na stene valja. Strjevanje taline valjaste oblike Fig. 8: Solidification of a cylindrcal liquid. Drugi primer: Mislimo si, da imamo namesto valjaste taline talino v obliki prisekanega stožca, ki pa naj ima, kot prej, trdno steno. Spet ga razrežemo na enako debele plasti in poskušajmo določiti, kako bi se morale ploskve med posameznimi plastmi deformirati, da bi talina ostala privezana na stene stožca. S slike 9 se vidi, da bi se morala potem ugrezniti v sredini v zgornjem delu (9. plast na sliki 9) zgornja ploskev valjaste plasti bolj, kot se je ugreznila spodnja ploskev te plasti ali talina se ne more strjevati tako, kot smo predpostavili, če naj bo »obešena« na stene stožca. Nujno potegne del strjene plasti na steni stožca navzdol, v sredino prereza bloka. Če gre za konično obliko taline, je mogoče celo izračunati, na kateri višini bi se to zgodilo, če bi šlo za krčenje v stožec. plastic material similar to a jelly that shrinks as a vvhole in one piece. We suppose that such a plastic "metal" is somehovv hung on its boundary to the vvalls that have been solidifi-ed previously. When it shrinks, it puliš down also some parts of the wall, that are stili plastic. So the V-segre-gates are formed. This idea is supported by the fact that vvith ali the sul-phur prints from Figs. 4 to 7 in the structure found about 0.5 m above the basis of the ingot, the crystals are de-formed similarly, as it can be seen in the region of the V-segregations above. In the čase of the ingot D vvhere we have to do vvith a cylinder vvith a rather large diameter above the core the appearance of V-segregations could not be avoided. From the photo of the etched surface it could be seen that in the region that is normally critical, in the čase of the ingot D the structure is extremely homogenous. It could be stated that in the upper part of the ingot an extremely good and homogenous structure is achieved and it is no more possible to say that there are any impurities penetrating from the top into the middle of the ingot. The only V-segregate in the middle of the ingot D can be explained in the follovving way. V-SEGREGATIONS Let us suppose the liquid metal is of the form of a cylinder and the vvalls of this cylinder are solid, made of the same material. Due to shrinkage the volume gets smaller and the upper level moves dovvn in the middle. Let us further suppose that the cylinder is cut uniformly to smaller equidistant layers according to Fig. 8. Due to solidification each layer shrinks for 4 % so that the upper flat circular surface moves dovvn and it attains the form of an inverted cone. In Fig. 8 the subsequent cylindrical layers are indicated by dotted lines. The full lines indi-cate the form of these layers after the solidification. Such picture vvould be obtained if the liquid stays att-ached to the vvalls of the cylinder. Another example: If instead of a cylinder vve have to do vvith a conical form of the liquid steel attached to solid vvalls of the same material. Let us again cut it into parallel layers of equal thickness and let us try to find out, hovv the flat ba-sic surfaces of these layers vvould be deformed at solidification, if the liquid stays attached to the conical vvalls. From Fig. 9 it can be seen that something nevv must happen if vve vvant to fulfil previous condition. According to Fig. 9 the upper surface of the 9th layer should move tovvards the center much deeper than the lovver surface. With the other words: The liquid metal can not be solidified in this way. It necessarily puliš a piece of the solid wall dovvn into the middle of the cross-section of the ingot. If the form of the liquid metal is conical, it is possible to evaluate also mathematically vvhere it happens, if the basic surfaces of the subsequent layers get the form of an inverted cone. Let the volume of the liquid steel of the form of a truncated cone vvith the radius Ro and vvith the angle a (Fig. 9) be reduced for t] in the process of solidification. The upper level is supposed to be lovvered so that it stays attached to the conical wall. In the middle it attains the form of an inverted cone. The critical height, vvhere Slika 9: Strjevanje taline konične oblike Fig. 9: Solidification of a conical liguid. Pri talini, ki ima obliko pokončnega prisekanega stožca, pri katerem je kot med stranskim robom osnega prereza in radijem a in polmer osnovne ploskve R<„ naj se volumen pri strjevanju zmanjša za ti. Gladina naj se ugrezne tako, da ostane na robu »privezana« na steno stožca, v sredini pa naj dobi obliko stožca, ki je z vrhom obrnjen navzdol. Kritična višina, pri kateri se v takem primeru prvič »strga« stena in ugrezne navzdol, se izračuna po formuli: Če se to ne bi zgodilo, bi se sicer morala zgornja ploskev 9. valjaste plasti po sliki 9 ugrezniti bolj kot spodnja. S slike 4 lahko ocenimo, da je približno tg a 10 in če je r| = 0.04, lahko ocenjujemo, da se bo pojavil prvi ugrez približno 1.5 metra nad osnovno ploskvijo ingota. Natančnejši izračuni pa kažejo, da se pri pogojih strjevanja valjasta plast taline v sredini ugrezne celo nekoliko bolj, kot če bi šlo za ugrezanje v obliki stožca. V primeru D opazujemo le eno večjo izcejo v obliki črke V v zgornjem delu prereza bloka, nekako tam, kjer se stožčasti del konča in začenja valjasti del. To je razumljivo, saj smo s stanjšano steno kokile dosegli, da se stožec nadaljuje v dosti širok valj. ZAKLJUČEK 1. Na potek izcej in na strukturo v prerezu jeklenega bloka, ki je bil ulit v kokilo, lahko vplivamo s primerno regulacijo hitrosti strjevanja. 2. Z rekonstruirano kokilo smo nakazali, kako je mogoče odpraviti sekundarni lunker, rahlo sredino in povečano koncentracijo nečistoč v sredini bloka. the boundary tears off and slips down, can be calculated from the follovving relationship: H.R.•*«•(,-f^) If it does not happen, the upper surface of the 9th lay-er according to Fig. 9 should be lovvered deeper than the bottom surface of the same layer. From Fig. 4 it can be estimated that tg a » 10 and if r| = 0.04 we can cal-culate the position of the first V-segregate. It vvould ap-pear approximately 1.5 m above the basis of the ingot. From more precise calculations it is evident that the lovvering of the middle of the surface of a layer vvould be even deeper than it is the čase vvith the inverted cone, if the real shrinkage is taken into account. In the čase of the ingot D only one V-segregate in the upper part of the ingot can be observed just in the plače vvhere the cone starts to continue into the cylin-der. The thinner wall in the upper part of the mould thus enables the formation of a cylinder vvith a large diameter. CONCLUSION 1. The course of segregates and the structure in the cross-section of a steel ingot čast into a mould can be modified by a proper regulation of the speed of solidification. 2. With the reconstructed mould it is shovvn how the secondary shrinkage hole can be suppressed together vvith the "soft middle" and vvith the increased concentra-tion of impurities in the middle of ingot. 3. Since vvith a proper form of the mould vvall the homogeneity of the upper part of the ingot can be con-siderably increased it could be expected that difficulties 3. Ker znamo s primerno obliko kokile znatno povečati homogenost v zgornjem delu bloka, pričakujemo, da bi s tem odpravili težave, ki se pojavljajo pri valjanju nekaterih kvalitet zaradi nehomogenosti v sredini (dvo-plastnost in podobno). 4. To trditev bo treba še praktično preveriti. Nova oblika kokile bo seveda nekoliko drugačna. Opisani poskusi bodo služili za izhodišče za delo v proizvodnji. 5. Odpirajo se možnosti za razvoj računalniškega krmiljenja in avtomatizacijo ohlajanja tudi ulitih blokov drugačnih oblik, kar je še posebno pomembno v livarstvu. appearing vvith the rolling of some qualities of steel slabs due to inhomogeneities in the middle could be avoided. 4. This statement must be also verified practically. The form of the reconstructed mould for practical pur-poses will be a little different. The results of the experi-ments discussed above will be starting point for the practical measures. 5. New possibilities in the development of the Computer regulation and automation of process of cooling of the čast steel ingots of different forms seem to be open-ed. It vvould be very important especially for the techno-logical development in foundries. LITERATURA/REFERENCES 1. P. N. Hansen: Numerical simulations of the solidification proces, 350—356, Solidification and Casting of Metals, Pro-ceedings of International conference on Solidification, Uni-versity Sheffield, 18—21 July 1977 2. F. VVeinberg, J. Lait, R. Pugh: Solidification of high carbon steel ingots, 334—339, Solidification and Casting of Metals, Proceedings of International conference on Solidification, University Sheffield, 18-21 July 1977 3. F. Oeters, K. Ruttiger, H.J.Selenz: VVarmeubergang beim Blockguss, Giessen und Erstarren von Stahl, Band I., 144—195 Informationstagung, Luxembourg 4. R. D. Pehlke, J. T. Berry, W. Erickson, C. H. Jacobs: Simulation of shaped casting solidification, 371—379, Solidification and Casting of Metals, Proceedings of International conference on Solidification, University Sheffield, 18—21 July 1977 5. P. R. Beeley: Keynote Address Solidification and aspects of čast metal quality, 319—324, Solidification and Casting of Metals, Proceedings of International conference on Solidification, University Sheffield, 18—21 July 1977 6. F. Oeters, K. Sardemann: Untersuchungen zum zeitlichen Verlauf der Erstarrung in der Randzone erstarrenden Eis-ens, Arch. Eisenhuttenwes. 45, (1974), 8, August, 517—524 7. K. Schvverdtfeger: Anvvendung der Methode des VVarmebi-lanzintegrales zur Berechnung der Erstarrungsgeschvvin-digkeit von Eisen-Kohlenstoff-Legierungen, Arch. Eisenhut-tenvves. 44 (1973) 6, Juni, 411-418 8. W. Schvvarz, R. Jeschar: Thermohydraulisches Analogiemo-dell zur Simulation der Blockerstarrung, Arch. Eisenhut-tenvves. 44 (1973), 6, Juni, 419—425 9. V. K. Chuang, K. Schvverdtfeger: Experimentalle und the-oretische Untersuchung der Erstarrung einer Eisen-Kohlen-stoff-Legierungen mit 0.6% C, Arch. Eisenhutenvves. 44 (1973), 5, Mai, 341-347 10 J Szekely, N. J. Themelis: Rate Phenomena in Process Me-tallurgy John Wiley & Sons Inc., New York 1971 11. T. Takahashi, K. Ichikavva, M. Kudou: Effect of fluid flow on macrosegregation in steel ingots 331—333, Solidification and Casting of Metals, Proceedings of International conference on Solidification, University Sheffield, 18—21 July 1977 12. J. Froeber, F. Oeters: On the mechanical behaviour of steel during solidification, Arch. Eisenhuttenvves. 51 (1980) 2, Februar, 43—49 13. R. Jha, T. Mukerjee: Shrinkage at peritectic temperature — its influence on cracking of steel ingots, Transactions of the Indian Institute of Metals, Vol. 29, 1, Feb. 1976, 30-35 14. T. Kolenko, B. Brudar, F.Mlakar, V. Tucič, S. Paulin; Vpliv oblike in forme na strjevanje ulitih valjev, Poročilo VTOZD Montanistika, December 1983 15. R. A. Entvvistle, J. E. Gruzleski, P. M. Thomas: Development of porosity in aluminium-base alloys, 345—349, Solidification and Casting of Metals, Proceedings of International conference on Solidification, University sheffield, 18—21 July 1977 16. A. Razinger: Mehanizem porazdelitve svinca v jeklu in njegov vpliv na strukturne in fizikalne lastnosti v odvisnosti od prisotnih elementov, Magistrsko delo, Univerza v Ljubljani, FNT, Oddelek za montanistiko, odsek za metalurgijo, 1973 17. A. J. Pokorny: De Ferri Metallographia III, 1967, Solidification of Steel, IRSID, Editor Berger — Levrault, Pariš 18. P. Oberhoffer: Das technische Eisen, S. 311, Julius Spring-erVerlag, Berlin 1936 19. Ju. Ja. Skok, G. A. Lubenec, F. I. Nečeporenko, V. M. Doro-feev, Z. L. Kozlova: Sniženie zonalnoj himičeskoj neodno-rodnosti slitkov putem modificirovanija stali, stal, 1986, 2, 19—22 20. F. Beneš, L. Beračkova, M. Kepka. L. Novak: Nestejnorod-nosti v ingotech o velkych hmotnosech, Hutnicke listy, 1986, č.2, 87-92 21. F. Esser, H. Brennecke: Rechnersimulation der Blockerstarrung in einer Kokille unter besonderer Berucksichti-gung des VVarmekontakts Block/Kokille, Neue Hutte, 24, 1979, 12,455-459 Bele kromove litine za valje, legirane z molibdenom White Chromium Čast Irons for Rolls, Alloyed vvith Molybdenum D. Kmetic*, F. Mlakar**, V. Tucič**, J. Žvokelj*, UDK: 669.15'26-194:669.14.018.255 F. Vodopivec*, M. Jakupovič*, B. Ralič* ASM/SLA: M28, N8b, TSk, 5, Cr, W23k Kromove bele litine, legirane z molibdenom in še nekaterimi drugimi elementi, se zaradi dobre obrabne obstojnosti, trdote in zadovoljivih mehanskih lastnosti vedno več uporabljajo za d vos lojno lite valje. Litine imajo tudi dobro korozijsko obstojnost. Delo obravnava mikrostrukturne značilnosti zlitin v litem stanju in po toplotni obdelavi. Narejena sta izotermna transformacijska diagrama za destabilizacijo avstenita in destabiliziran avstenit in kontinuirni transformacijski diagram za destabiliziran avstenit. UVOD V valjarnah je poleg ustrezne kvalitete valjanih proizvodov zelo pomembna ekonomičnost proizvodnje. Določena jekla se vroče valjajo v nizkih temperaturnih C v °/o SI. 1: Kemična sestava belih kromovih litin in litin za druge vrste valjev v faznem diagramu (KV — kovani valji, AD — adamitni valji, IND — indefinitni valji, KGR — nodularni valji, Cr — bele kromove litine, 8) Fig. 1 Chemical composition of vvhite chromium čast irons, and čast irons for other types of roils in the phase diagram (KV — forged rolls, AD — adamite rolls, IND — indefinite chill rolls, KGR — spheroidal-graphite rolls, Cr — vvhite chromium čast irons, 8) * SŽ — Metalurški inštitut Ljubljana ** SŽ — Železarna Štore * Institute of Metallurgy, Ljubljana ** Store Ironworks Chromium vvhite čast irons alloyed vvith molybdenum and some other elements are more and more applied for compound čast rolls due to good vvear resistance, hardness, and satisfactory mechanical properties. The čast irons have also good corrosion properties. Paper treats the microstructural characteristics of čast irons as čast, and after the heat treatment. Isother-mal transformation diagrams for the destabilization of austenite, and for the destabilized austenite vvere con-structed next to the continuous transformation diagram for the destabilized austenite. INTRODUCTION In rolling plants, the economy of manufacturing is very important next to the suitable quality of rolled pro-ducts. Some steel is hot rolled in low-temperature re-gions vvith high partial reductions, and low permissible dimensional tolerances. Narovver and narovver toler-ances are demanded also for the cold rolled strips. These are the reasons that rolls of vvhite čast iron vvith high chromium cpntent, and alloyed vvith Mo, Ni (Cu), V, Ti, and W are more and more used in hot and cold rolling plants. The rolls are čast by a compound centrifugal casting. Data on chemical composition of rolls of vvhite chromium čast irons are in references given in wide in-tervals. Data on manufacturing rolls, and on their heat treatment are scarce. The phase diagram in Fig. 1 pres- Ogljik v utežnih procentih Weight percentage of carbon SI. 2: Fazni diagram Fe-Cr-C za 17 % Cr Fig. 2 Fe-Cr-C phase diagram at 17 % Cr SI. 3: Likvidus površine in lega zlitin glede na razmerje Cr/C v faznem diagramu Fe-Cr-C po R. S. Jacksonu (1) Fig. 3 Liquidus surfaces and the position of alloys related to the Cr/C ratio in the Fe-Cr-C phase diagram, according to R. S. Jackson (1) področjih z velikimi parcialnimi redukcijami, pri čemer se zahtevajo ozke dimenzijske tolerance. Vedno bolj ozke tolerance se zahtevajo tudi pri hladno valjanih trakovih. To so razlogi, da se v vročih in hladnih valjarnah vedno bolj uporabljajo valji iz bele litine z visoko vsebnostjo Cr, legirane še z Mo, Ni (Cu), V, Ti in W. Valji se izdelujejo po postopku dvoslojnega centrifugalnega litja. Literaturni podatki o kemični sestavi valjev iz bele kromove litine so podani v širokih mejah. Podatki o izdelavi valjev in toplotni obdelavi so zelo skopi. V faznem diagramu na sliki 1 je za primerjavo navedeno področje kemične sestave valjev iz bele kromove litine in drugih vrst valjev. Zlitine Fe-Cr-C so že dolgo poznane in so v literaturi opisane številne raziskave. Fazni diagram Fe-Cr-C za 17 % Cr je prikazan na sliki 2 (3). Za razvoj teh zlitin so poleg začetnih raziskav F. Osmonda, ki je v mikrostruk-turi omenjenih litin že leta 1892 opazil kompleksne karbide, najpomembnejše raziskave R. S. Jacksona, ki je v faznem diagramu Fe-Cr-C opredelil likvidus površine (si. 3) in sistematične raziskave vpliva Mo na mikro-strukturne značilnosti, ki sta jih naredila F. Maratray in R. Usseglio-Nanot (1,2, 4). Mikrostruktura belih kromovih litin sestoji iz primarnih in evtektičnih karbidov in avstenitne matice, oziroma njenih transformacijskih produktov (sekundarni karbidi, perlit, bainit, martenzit). Za mikrostrukturne značilnosti je zelo pomembno razmerje Cr/C in vsebnost legiranih elementov, predvsem Mo, Mn, Ni (Cu) in W. EKSPERIMENTALNO DELO Na osnovi literaturnih podatkov, ki smo jih imeli na voljo, smo v železarni Štore izdelali preizkusne taline z različno vsebnostjo legirnih elementov, in sicer z 2,5 do 3,8 % C, 11,3 do 19,4 % Cr, 0,39 do 0,66 % Mo, 0,59 do 1,37 % Si, 0,68 do 0,93 % Mn, 0,56 do 0,78 % Ni, 0,023 do 0,11 % Ti in z 0,06 do 0,11 % V. Štiri zlitine smo legi-rali z 0,80 do 0,93 % W. Zlitine legirane z W so trše in se uporabljajo za valje za hladno valjanje trakov. Vsebnost P mora biti pod 0,08 % in S pod 0,05 %. Preizkusne zlitine imajo razmerje Cr/C od 3,62 do 7,76. Vzorce, preizkusne valjčke, premera 100 in višine 150 mm, smo ulili tako, da je bila polovica valjčka ulita v kokilo in polovica v pesek. Tako smo dobili na enem vzorcu dve različni hitrosti strjevanja. Pogoji litja bistveno vplivajo na izoblikovanje mikrostrukture in s tem na mehanske lastnosti litine. Zato ents the regions of chemical compositions of rolls of vvhite chromium čast iron, and of some other types of rolls (8), Fe-Cr-C alloys are already for a long tirne known, and numerous investigations are cited in references. The Fe-Cr-C phase diagram for 17 % Cr is shovvn in Fig. 2 (3). For development of these alloys, the most essential are the investigations by R. S. Jackson who determined the liquidus surfaces in the Fe-Cr-C phase diagram (Fig. 3), and the systematic investigations on the influence of Mo on the microstructural characteristics done by F. Mara-tray, and R. Usseglio — Nanot, beside the initial investigations by F. Osmond who already in 1892 observed complex carbides in the microstructure of the men-tioned čast irons (1, 2, 4). Microstructure of vvhite chromium čast irons con-sists of primary and eutectic carbides, and austenitic matrix, or of its transformation products (secondary carbides, pearlite, bainite, martensite). Essential for the microstructural characteristics are the Cr/C ratio and the content of alloying elements, mainly Mo, Mn, Ni (Cu), and W. EXPERIMENTAL WORK Based on the data in references, being available, test melts vvith various contents of alloying elements, i. e. vvith 2.5 to 3.8 % C, 11.3 to 19.4 % Cr, 0.39 to 0.66 % Mo, 0.59 to 1.37 % Si, 0.68 to 0.93 % Mn, 0.56 to 0.78 % Ni, 0.023 to 0.11 % Ti, and 0.06 to 0.11 % V vvere made in the Štore Ironvvorks. Alloys vvith added W are harder and they are used for rolls for cold rolling of strips. Phospho-rus content must be belovv 0.08 %, and sulphur belovv 0.05 %. The test melts had the Cr/C ratio between 3.62 and 7.76. The samples as testing cylinders vvith diameter 100 mm and 150 mm high vvere čast so that one half of the cylinder vvas čast into mould, another one into sand. Thus two various solidification rates vvere obtained on the same specimen. Casting conditions have essential influence on the formation of microstructure, and thus on the mechanical properties. Therefore melting points and solidification in-tervals vvere determined for some alloys. Microstructural characteristics of as čast alloys, and after the heat treatment vvere determined by investigations vvith optical microscope, scanning electron micro-scope (SEM), and electron microanalyzer. To reveal the microstructural characteristics various etching agents (nital, Villela's, ferric chloride, alkaline picrate, Muraka-mi's, and 4% sodium hydroxide saturated vvith pota- smo za nekatere zlitine določili temperature tališča in intervale strjevanja. Mikrostrukturne značilnosti zlitin v litem stanju in po toplotni obdelavi smo opredelili s preiskavami z optičnim mikroskopom, v raster elektronskem mikroskopu (SEM) in v elektronskem mikroanalizatorju. Za odkrivanje mikrostrukturnih značilnosti smo uporabili različna jedkala (nital, Villela, feriklorid, alkalijski pikrat, Murakami in 4 % natrijev hidroksid, nasičen s kalijevim permanganatom). Sekundarne karbide in faze, nastale pri transformaciji avstenita, smo lahko dobro opredelili v SEM. V elektronskem mikroanalizatorju smo določili sestavo primarnih in evtektičnih karbidov in koncentracije nekaterih legirnih elementov v matici. Za eno od zlitin z najustreznejšo kemično sestavo in mikrostrukturo smo naredili izotermna transformacijska diagrama za nedestabilizirano in destabilizirano av-stenitno matico in kontinuirni transformacijski diagram za destabiliziran avstenit. Od mehanskih lastnosti smo merili le trdoto zlitin in posameznih mikrostrukturnih faz. V literaturi smo zasledili raziskave, ki obravnavajo upogibno trdnost in žilavost teh zlitin (13). Za valje je poznavanje teh parametrov zelo pomembno, vendar smo zaradi težavne priprave mehanskih preizkušancev te preiskave odložili na kasnejši čas. REZULTATI PREISKAV Tališča in interval strjevanja zlitin Žilavost in obrabna obstojnost litine je tem boljša, čim bolj drobni so evtektični karbidi in čim enakomer-neje so porazdeljeni po matici. (19) Zato mora potekati strjevanje belih kromovih litin hitro. Pri previsokem pregretju in počasnem strjevanju lahko nastanejo poleg grobih evtektičnih klarbidov še veliki primarni karbidi. V talilnem mikroskopu smo določili tališča in intervale taljenja nekaterih zlitin, izbranih tako, da smo pokrili ves interval razmerij Cr/C (tabela 1). Zaradi reka-lescence je razlika med talilnim in strjevalnim intervalom majhna. Zlitine so močno izcejane in se rezultati paralelk in vrednosti, izmerjene večkrat na istem vzorcu, med seboj precej razlikujejo. Tabela 1: Tališča in intervali taljenja Zlitina %C % Cr Cr/C Nastanek kapljic °C Začetek talj. Staljeno Interval talj. "C °C •c 1 2,49 19,31 7,76 1200 1245 1345 100 2 2,63 19,43 7,39 1170 1220 1350 130 (legira-noz W) 3 2,72 14,90 5,48 1200 1250 1305 55 4 2,76 19,21 6,96 1210 1250 1325 75 5 3,20 17,95 5,61 1210 1250 1285 35 7 3,31 11,97 3,62 1190 1225 1300 75 8 3,48 16,23 4,66 1170 1215 1250 35 (legira-noz W) Zlitine imajo tališča med 1350 in 1250 °C. Čim bolj se sestava zlitine približuje evtektični sestavi, ožji je interval strjevanja. Na sliki 2 se vidi, da ima zlitina s 17 % Cr evtektično sestavo pri 3,4 % Č. Na tališče in in- ssium permanganate) were applied. Secondary car-bides, and phases formed during the transformation of austenite were well determined by SEM. Electron micro-analyzer helped us to determine the composition of pri-mary and eutectic carbides, and the concentrations of some alloying elements in the matrix. For one of the alloys, vvith the most suitable chemical composition and the microstructure, the isothermal transformation diagrams for undestabilized and des-tabilized austenitic matrix, and the continuous transformation diagram for destabilized austenite were con-structed. Of mechanical properties only hardness of alloys and of single microstructural phases vvas measured. In refer-ences, investigations treating the bending strength, and the toughness of these alloys vvere found (13). Though the knovvledge of these properties is very important for the behaviour of rolls, these investigations vvere post-poned for later due to difficult preparation of testing specimens. RESULTS OF INVESTIGATIONS Melting Points and Soliditication Interval of Alloys Toughness and vvear resistance of the alloy are the better the smaller are eutectic carbides, and the more uniformly they are distributed in the matrix (19). There-fore the solidification of vvhite chromium čast irons mUs.t be fast. At a too high superheating and low solidification rate big primary carbides next to coarse eutectic carbides can be formed. Melting points and solidification intervals of some al-loys vvere determined by fusion microscope. The alloys vvere chosen in such a way that he vvhole interval of the Cr/C ratios vvas covered (Table 1). Due to recalescence the difference betvveen the melting and the solidification interval is small. The alloys exhibit intensive segregating, thus the results of parallel tests, and the values measured more times on the same sample differ a great deal. Table 1 Melting Points and Solidification Intervals Alloy %C % Cr Cr/C Formation of Begin. of Melted Melting interval "C drops 'C melting "C •c 1 2,49 19,31 7,76 1200 1245 1345 100 2 2,63 19,43 7,39 1170 1220 1350 130 (alloyed with W) 3 2,72 14,90 5,48 1200 1250 1305 55 4 2,76 19,21 6,96 1210 1250 1325 75 5 3,20 17,95 5,61 1210 1250 1285 35 7 3,31 11,97 3,62 1190 1225 1300 75 8 3,48 16,23 4,66 1170 1215 1250 35 (alloyed vvith W) The alloys have the melting points betvveen 1350 and 1250° C. The closer is the alloy composition to the eutectic composition the narrovver is the solidification interval. It is evident from the Fig. 2 that the alloy vvith 17 % Cr has eutectic composition at 3.4 % C. The melting points and the solidification intervals are mainly influenced by the carbon content, to a lesser extent by the Cr/C ratio, terval strjevanja vpliva predvsem vsebnost ogljika, manj pa razmerje Cr/C in koncentracije ostalih legirnih elementov. Od vsebnosti ogljika, ki sicer znižuje temperaturo tališča, in razmerja Cr/C je odvisen delež karbidne faze v mikrostrukturi, kar tudi vpliva na tališče in interval strjevanja zlitin. Iz faznih diagramov Fe-Cr-C se vidi, da se z naraščajočo vsebnostjo Cr evtektična točka pomika v levo in k višjim temperaturam. Mikrostruktura zlitin v litem stanju Mikrostruktura zlitin je odvisna od kemične sestave, razmerja Cr/C in pogojev strjevanja. Vse zlitine smo le-girali z Mo, zato imajo v mikrostrukturi poleg primarnih in evtektičnih karbidov M7C3 tudi karbide Mo2C. Mikrostruktura evtektika je odvisna od deleža avste-nitne faze, ki nastaja med procesom strjevanja. Če nastane med strjevanjem veliko avstenita in je majhen delež preostale taline, ki se strdi kot evtektik, imajo evtek-tični karbidi tendenco, da segregirajo vzdolž kristalnih mej avstenitnih zrn. Take mikrostrukture, ki je značilna za zlitine z do 20 % karbidne faze, pri naših zlitinah, ki imajo od 25 do 35 % karbidne faze, nismo opazili. V nekaterih zlitinah smo opazili v evtektiku bolj ali manj lamelarno izoblikovane karbide, ki rastejo iz sredine meddendritskih prostorov (si. 4) Pri drugih zlitinah, pri katerih je avstenitne faze zelo malo in ta praktično ni omejevala strjevanja evtektika, imajo karbidi popolnoma lamelarno obliko (si. 5). Čeprav so veliki primarni karbidi heksagonalne oblike značilni za litine z nad 35 % karbidne faze, smo te opazili tudi pri nekaterih naših zlitinah, in to predvsem na sredini preizkusnih valjčkov, kjer so bili za njihov nastanek ustreznejši pogoji (si. 6.). Deleže karbidne faze v mikrostrukturi smo za nekatere zlitind izračunali po enačbi (1): % K = 12,33 (%C) + 0,55 (% Cr) - 15,2 Izračunane vrednosti se dobro ujemajo z vrednostmi, ki smo jih dobili z meritvami po linearni intercepcij-ski metodi v optičnem mikroskopu (tabela 2). Vsebnosti Mo in W sta majhni in ne vplivata bistveno na delež karbidne faze. Tabela 2: Delež karbidne faze (% K) v mikrostrukturi Zlitina Cr/C %C % Cr % K izračunan % K izmerje 7 3,62 3,31 11,97 32,2 30 6 4,80 3,21 15,42 32,9 29 3 5,48 2,72 14,90 26,5 26 5 5,61 3,20 17,95 34,1 35 2 7,39 2,62 19,43 27,9 28 1 7,76 2,49 19,31 26,1 25 Mikrostruktura matice je odvisna od razmerja Cr/C, vsebnosti Mo in pogojev ohlajevanja. Matica ima v litem stanju avstenitno mikrostrukturo, oz. je med ohlajanjem potekla delna ali popolna transformacija avstenita v perlit. Pri litju v kokilo potekata strjevanje in ohlajanje hitreje, kot pri litju v pesek, in perlitna transformacija je zavrta. Pri zlitinah brez Mo lahko pričakujemo popolnoma avstenitno matico pri razmerju Cr/C večjem od 7,2 (1,2). Z legiranjem z Mo se razmerje Cr/C, pri katerem dobimo popolnoma avstenitno matico, pomika proti nižjim vrednostim. To pomeni, da ima lahko litina pri isti vsebnosti Cr več C in zato v mikrostrukturi večji delež karbidne faze in avstenitno matico. and the concentration of the other alloying elements. The carbon content vvhich namely reduces the melting point, and the Cr/C ratio determine the amount of car-bide phase in the microstructure vvhich has also influence on the melting point and the solidification interval of alloys. The Fe-Cr-C phase diagrams shovv that the in-creasing Cr content shifts the eutectic point tovvards the left and to higher temperatures. Microstructure of As Čast Alloys Microstructure of alloys depends on the chemical composition, the Cr/C ratio, and the conditions of solidification. Ali the alloys were alloyed vvith Mo, thus the microstructure contains also Mo2C carbides next to the primary and eutectic M7C3 carbides. Microstructure of eutectic depends on the amount of austenitic phase vvhich is formed during the solidification. If a high amount of austenite is formed during the solidification, and the portion of the remaining melt vvhich solidifies eutectically is small, the eutectic carbides exhi-bit tendency to segregate along the boundaries of austenitic grains. Such a microstructure being charac-teristic for the alloys vvith up to 20 % of carbide phase was not observed in our alloys which contained 25 to 35 % of carbide phase. In some alloys more or less lamellar carbides were observed vvhich grovv from the centre of interdendritic spaces (Fig. 4). In other alloys vvith a very lovv amount of austenitic phase vvhich did not hinder the solidification of eutectic, the carbides exhibited fully lamellar shape (Fig. 5). Though big primary carbides of hexagonal shape are characteristic for the čast irons vvith over 35 % of carbide phase, they vvere obseved also in some of our alloys, but mainly in the centre of the testing cylinders vvhere the conditions for their formation vvere the most suitable (Fig. 6). The portions of carbide phase in the microstructure vvas for some alloys evaluated by the equation (1): % K = 12.33 (% C) +0.55 (% Cr) - 15.2 The obtained values are in a good agreement vvith the values vvhich vvere obtained by the measurements in optical microscope by the intercept method (Table 2). Contents of Mo and W are lovv and they do not influence essentially the portion of carbide phase. SI. 4: Eutektični karbidi rastejo iz sredine meddendritskih prostorov, lito stanje (3,21 % C, 15,42 % Cr, 0,53 % Mo, Cr/C 4,80). Pov. 100 x Fig. 4 Eutectic carbides grovv from the centre of interdendritic spaces, as čast (3.21 % C, 15.42% Cr, 0.53% Mo, Cr/C 4.80). Magn. 100x Table 2 Portion of Carbide Phase (% K) in the Microstructure Zlitina Cr/C %C % Cr Karbidi primarni eutek. Matica perlit austenit 7 3.62 3.31 11.97 0.58 38.3Cr 32.9 Cr 6.7 Cr 0.55 Mo 0.45 Mo 0.26 Mo _ 0.8 Mn 0.7 Mn 0.6 Mn — 6 4.80 3.21 15.42 0.53 42.4Cr 37.0Cr 7.7 Cr 7.0 Cr 0.44Mo 0.46 Mo 0.26 Mo 0.34 Mo 3 5.48 2.72 14.90 0.56 44.8Cr 39.9 Cr — 10.5Cr 0.97 Mo 0.92 Mo - 0.13 Mo 0.9 Mn 0.8 Mn _ 0.7 Mn 5 5.61 3.20 17.95 0.63 48.4Cr 43.8 Cr — IO,2Cr 0.42 Mo 0.42 Mo — 0.23 Mo 0.95 Mn 0.9 Mn — 0.8 Mn 2 7.39 2.62 19.43 0.52 50.7 Cr 49.0 Cr _ 10.8Cr 0.83 0.49 Mo 0.46 Mo _ 0.3 Mo W 0.59 W 0.61 W 0.44 W 0.9 Mn 0.9 Mn _ 0.7 Mn 1 7.76 2.49 19.31 0.54 51.2Cr 49.0 Cr _ 10.5 Cr 0.41 Mo 0.44 Mo — 0.31 Mo Alloy Cr/C %C % Cr % K calculated measured 7 3.62 3.31 11.97 32.2 30 6 4.80 3.21 15.42 32.9 29 3 5.48 2.72 14.90 26.5 26 5 5.61 3.20 17.95 34.1 35 2 7.39 2.62 19.43 27.9 28 1 7.76 2.49 19.31 26.1 25 ■MM MHHBP* "»PSMfi^tess«. "'SfiSdK. HB SI. 5: Lamelami eutektični karbidi, lito stanje (3,20 % C, 17,95 % Cr, 0,63 % Mo, Cr/C 5,61). Pov. 100 x Fig. 5 Lamellar eutectic carbides, as čast (3.20% C, 17.95% Cr, 0.63% Mo, Cr/C 5.61). Magn. 100 x Preiskovane zlitine so legirane z Mo in preizkusni valjčki imajo na presekih, ulitih v kokilo pri razmerjih Cr/C nad 5,5, popolnoma avstenitno matico (si. 5). Le v večji oddaljenosti od površine smo pri nekaterih vzorcih opazili v mikrostrukturi manjša perlitna zrna. S padajočo vrednostjo razmerja Cr/C narašča v matici delež perlitne faze. Avstenitno perlitna mikrostruktura matice je prikazana na sliki 4. Perlitno matico lahko reavstenitiziramo in tako zagotovimo, da ima litina po destabilizaciji in transformaciji s stališča mehanskih lastnosti ustreznejšo mikrostrukturo matice (martenzit). Menimo, da z ogrevanjem avstenitno perlitnih litin 50 °C pod solidus temperaturo dobimo avstenitno mikrostrukturo matice. Za natančnejše pogoje reavstenitizacije so v literaturi podani diagrami (1, 2). Vsekakor pa je ugodneje, da z razmerjem Cr/C, legiranjem z Mo in pogoji strjevanja že v litem stanju zagotovimo litini avstenitno mikrostrukturo matice (9). Kemična sestava karbidov in matice Koncentracije Cr, Mo, Mn in W v primarnih in ev-tektičnih karbidih in v matici, izmerjene v elektronskem mikroanalizatorju, so podane v tabeli 3. Meritve smo naredili na vzorcih ulitih v kokilo. Tabela 3: Vsebnosti Cr, Mo, Mn in W v karbidih in matici Microstructure of matrix depends on the Cr/C ratio, amount of Mo, and conditions of solidification. Matrix as čast has austenitic microstructure, or a partial or com-plete transformation of austenite into pearlite occured during the solidification. Solidification and cooling are faster in casting into moulds than in casting into sand, and pearlitic transformation is retarded. In alloys vvithout Mo fully austenitic matrix can be expected at the Cr/C ratios higher than 7.2 (1, 2). Alloying vvith Mo shifts the Cr/C ratio at vvhich fully austenitic matrix is obtained tovvards lovver values. This means that čast iron can con-tain at the same Cr more C, and thus a greater portion of carbide phase can be in the microstructure. Investigated alloys were alloyed vvith Mo, and the testing cylinders exhibit on the cross sections čast into mould a fully austenitic matrix if Cr/C ratio vvas over 5.5 (Fig. 5). Only at a greater distance from the surface smaller pearlitic grains vvere observed in the microstructure of some samples. The reduced Cr/C ratio causes an increased amount of pearlitic phase in the matrix. The austenitic-pearlitic microstructure of the matrix is shovvn in Fig. 4. Pearlitic matrix can be reaustenitized, and thus it is ensured that the čast iron has a more suitable microstructure of matrix (martensite) from the vievvpoint of mechanical properties after the destabilization and the transformation. It is supposed that heating austenitic-pearlitic čast iron at 50° C belovv the solidus temperature gives austenitic microstructure of the matrix. More de-tailed conditions of reaustenitization are given in graphs in references (1,2). Anyhow, it is more favourable to en-sure the austenitic microstructure of the matrix in the as čast alloy by the Cr/C ratio, alloying vvith Mo, and the conditions of solidification. Chemical Composition of Carbides and of Matrix Concentrations of Cr, Mo, Mn, and W in the primary and the eutectic carbides, and in the matrix, measured by the electron microanalyzer are presented in Table 3. Table 3 Contents of Cr, Mo, Mn, and W in Carbides and in the Matrix V diagramu na sliki 7 je prikazana odvisnost med razmerjem Cr/C v zlitinah in razmerjem Fe/Cr v karbidih Alloy Cr/C % C % Cr M Carbides Primary Eutec. Matrix Pearlite Austen. 7 3.62 3.31 11.97 0.58 38.3 Cr 32.9 Cr 6.7 Cr _ 0.55 Mo 0.45 Mo 0.26 Mo — 0.8 Mn 0.7 Mn 0.6 Mn — 6 4.80 321 15.42 0.53 42.4 Cr 37.0Cr 7.7 Cr 7.0 Cr 0.44 Mo 0.46 Mo 0.26 Mo 0.34 Mo 3 5.48 2.72 14.90 0.56 44 8Cr 39.9 Cr — 10.5 Cr 0.97 Mo 0.92 Mo — 0.13MO 0.9 Mn 0.8 Mn 0.7 Mn 5 5.61 3.20 17.95 0.63 48 4Cr 43.8Cr — 10.2 Cr 0.42 Mo 0.42 Mo — 0.23 Mo 0.95 M n 0.9 Mn — 08 Mn 2 7.39 262 19.43 0.52 50 7Cr 49.0 Cr — 10.8 Cr 0.83 0.49 Mo 0.46 Mo — 03 Mo W 0.59 W 0.61 W — 0.44W 0.9 Mn 0.9 Mn — 0.7 Mn 1 7.76 2.49 19.31 0.54 51.2Cr 49.0Cr — 10.5Cr 0.41 Mo 0.44 Mo — 0.31 Mo SI. 6: Primarni karbidi heksagonalne oblike. Pov. 100x Fig. 6 Primary carbides of hexagonal shape. Magn. 100x M7C3. Z razmerjem Fe/Cr je podana sestava karbidov M7C3, ki se sicer lahko spreminja od (Cr2Fe5) C3 do (Cr5Fe2)C3. Primarni in evtektični karbidi imajo v naših zlitinah sestavo od malo nad stehiometričnim razmerjem (Cr3Fe4) Cr3 do (Cr4Fe3) C3. Krivulja za evtektične karbide leži nad krivuljo za primarne karbide, ker imajo evtektični karbidi pri istih vrednostih Cr/C manjšo vsebnost Cr. Razlika v vsebnosti Cr med primarnimi in evtektičnimi karbidi je največja pri najnižjem razmerju Cr/C. Z naraščajočo vsebnostjo tega razmerja proti 8 se vsebnost Cr v primarnih in evtektičnih karbidih približuje isti vrednosti. Poleg Cr, ki v kristalni mreži karbidov nadomešča atome Fe, smo v karbidih izmerili tudi določene koncentracije Mo, Mn in W. Vsebnost Mo v zlitinah je majhna, zato je v mikrostrukturi malo karbidov Mo2C. Ti karbidi so drobni, vendar smo jih lahko določili v elektronskem mikroanalizatorju, kot tudi karbide W v zlitinah, legiranih s tem elementom. Vsebnost Cr v matici narašča z vrednostjo razmerja Cr/C. Meritve koncentracij Cr in Mo v matici so pokazale, da je ta zelo nehomogena (5). Odstopanja od povprečnih vrednosti so pri Cr v mejah ±20%. Bistveno večje je izcejanje Mo, in sicer večinoma v mejah ± 50 %. V nekaterih primerih pa smo v izcejah izmerili tudi do 2 % Mo. Podobno smo ugotovili, da tudi karbidi nimajo homogene sestave. Pri večjih, predvsem primarnih karbidih, je koncentracija Cr največja v sredini in se zmanjšuje proti robu karbidnega zrna. V karbidih smo merili koncentracije Cr, Mo in W tudi na vzorcih, žarjenih 2, 4 in 8 ur na temperaturi 1050 °C. Pri tej temperaturi poteka izločanje sekundarnih karbidov in s tem destabilizacija avstenitne matice. Izmerjene razlike v koncentraciji omenjenih elementov med litim in žarjenim stanjem niso sistematične in so odstopanja v mejah merilnih napak. Diagram izotermne destabilizacije avstenita Sistematične preiskave destabilizacije avstenita, izo-termna transformacijska diagrama za nedestabiliziran in destabiliziran avstenit (TTT) in kontinuirni transformacijski diagram za destabiliziran avstenit (CTT) smo naredili za zlitino 5 z naslednjo kemično sestavo: 3,2% C, 1,22% Si, 0,86% Mn, 0,035 % S, 0,030% P, 17,95% Cr, 0,63% Mo, 0,69% Ni, 0,08% Ti in 0,095% V. Zlitina ima razmerje Cr/C 5,61 in v litem stanju avstenitno matico (si. 5). Izotermna transformacijska diagrama smo naredili na osnovi metalografskih Measurements vvere made on the samples čast into mould. Plot in Fig. 7 gives the relationship betvveen the Cr/C ratio in alloys and the Fe/Cr ratio in M7C3 carbides. The Fe/Cr ratio defines the composition of M7C3 carbides vvhich varies betvveen (Cr2Fe5)C3 and (Cr5Fe2)C3. Primary and eutectic carbides in our alloys have the composition betvveen the composition vvhich is slightly above the stoichiometric one of (Cr3Fe4)C3, and the composition of (Cr4Fe3)C3. The curve for eutectic carbides is above the curve for primary carbides since eutectic carbides at equal Cr/C ratios have smaller contents of Cr. The differ-ence in Cr content betvveen the primary and the eutectic carbides vvas the highest at the lovvest Cr/C ratio. If this ratio goes tovvards 8, the Cr content in primary and eutectic carbides approaches to the same value. Beside Cr vvhich in crystal lattice of carbides substi-tutes Fe atoms, certain concentrations of Mo, Mn, and W vvere found in carbides. Mo content in alloys is small therefore the microstructure contains small amount of Mo2C. These carbides are fine but they were determined by the electron microanalizer, as well as the tungsten carbides in the alloys alloyed vvith that element. Cr content in matrix is increased vvith the increased Cr/C ratio. Measurements of Cr and Mo concentrations in the matrix shovved that matrix is very unhomogeneous (5). Deviations from the mean values are for Cr in the li-mits ±20 %. Essentially greater are segregations of Mo, mainly in limits ±50%. In some cases in segregations, even up to 2 % Mo vvas found. Similarly, it vvas found that also carbides do not have a homogeneous composition. In bigger, mainly primary carbides the concentration of Cr is the greatest in the centre and it is reduced tovvards the edge of the carbide grain. In carbides, the concentrations of Cr, Mo, and W vvere measured also in the samples annealed 2, 4 and 8 hours at 1050° C. At this temperatures secondary carbides are precipitated and thus the austenitic matrix is (Cr3Fe4)C3 (Cr4Fe3)C3 -(Cr2Fe5)C3 "234 5 6 7 8 Razmerje - Ratio : Cr/C SI. 7: Odvisnost med razmerjem Cr/C in razmerjem Fe/Cr v karbidih M7C3 Fig. 7 Relation betvveen the Cr/C ratio and Fe/Cr ratio in M7C3 carbides -(Cr5Fe2)C3 eutektični - eutectic karbidi carbides primarni karbidi -primary carbides preiskav. Mikrostrukturne spremembe smo opredelili v optičnem mikroskopu in v SEM. V nekaterih primerih, ko je bilo težko določiti mikrostrukturne komponente, smo si pomagali še z meritvami mikrotrdot in selektivnim jedkanjem (15, 16). Za izdelavo diagrama izotermne destabilizacije av-stenita smo vzorce izotermno žarili različno dolgo časa v temperaturnem področju med 500 in 1150° C. Razpad avstenita poteka v dveh temperaturnih področjih, ki se v ozkem področju prekrivata (si. 8). Za toplotno obdelavo zlitin je pomembna destabilizacija avstenita z izločanjem sekundarnih karbidov (Ks), ki poteka v višjem temperaturnem področju. Izločanje sekundarnih karbidov je najhitrejše med 940 in 990 °C. Nad temperaturo Ac3 poteka transformacija y = y + Ks in pod to temperaturo y = y + a + K,. Nad temperaturo Ac3 se iz avstenita izločajo karbidi M7C3. V temperaturnem področju med Ac3 in Ac, pa se iz avstenita izločajo tudi karbidi M23C6 (i0). Izločanje sekundarnih karbidov je za nadaljnjo toplotno obdelavo bistvenega pomena. Brez predhodne destabilizacije, pri kateri se zaradi izločanja sekundarnih karbidov v avstenitu zmanjša vsebnost Cr in C, transformacija avstenita v martenzit, kot tudi v bainit, niti ni mogoča. Perlitna transformacija pa poteka v destabiliziranem avstenitu počasneje. V nižjem temperaturnem področju razpada poteka transformacija avstenitne matice v perlit. Transformacija poteka najhitreje med 670 in 710 °C. Sekundarni karbidi se začnejo izločati iz avstenita po določeni inkubacijski dobi, in to ob kristalnih mejah med avstenitnimi zrni in na meji avstenitnih zrn z evtek-tičnimi karbidi. Proti sredini avstenitnih zrn poteka izločanje hitreje po določenih kristalografskih ravninah (si. 9, 10, 11). V začetni fazi izločanja so karbidi drobni, destabilized. The measured differences of concentra-tions of the mentioned elements betvveen the čast and annealed state are not systematic, and the deviations are in the limits of measuring errors. Diagram of Isothermal Destabilization of Austenite Systematic investigations of the destabilization of austenite, isothermal transformation diagrams for un-destabilized and destabilized austenite (TTT), and the continuous transformation diagram for destabilized austenite (CTT) were constructed for the alloy 5 with the follovving composition: 3.2% C, 1.22% Si, 0.86% Mn, 0.035 % S, 0.030 % P, 17.95 % Cr, 0.63 % Mo, 0.69 % Ni, 0.08 % Ti, and 0.095 % V. The Cr/C ratio of the alloy was 5.61, and the as čast alloy exhibits austenitic matrix (Fig. 5). The isothermal transformation diagrams were constructed from data of metallographic investigations. Microstructural variations were determined in optical microscope and by SEM. In some cases when the microstructural components were not easy to be determined, measurements of microhardnesses, and selec-tive etching were applied (15, 16). To construct the diagram of isothermal destabilization of austenite, the samples were isothermally annealed for various times in the temperature interval 500 to 1150° C. Decomposition of austenite occurs in two temperature intervals vvhich overlap in a narrow region (Fig. 8). The destabilization of austenite vvith precipita-tion of secondary carbides (Ks) occuring in the higher temperature interval is important for the heat treatment of alloys. Precipitation of secondary carbides is the fas-test betvveen 940 and 990° C. Above Ac3 transformation y = Y + Ks takes plače, and belovv that point y = y + + a + Ks. Above Ac3 M7C3 carbides are precipitated from austenite. In the temperature interval betvveen Ac3 and Ac, also M23C(5 carbides are precipitated from austenite (10). Precipitation of secondary carbides is essential for fur-ther heat treatment. VVithout the predestabilization vvhen due to the precipitation of secondary carbides the con-centrations of Cr and C are reduced, the transformation of austenite into martensite as well as into bainite is not possible. Pearlitic transformation is slovver in the destabilized austenite. In the lovver temperature interval of the decomposition, the transformation of austenitic matrix into pearlite takes plače. It is the fastest betvveen 670 and 710° C. SI. 9: Izločanje sekundarnih karbidov v avstenitu, 80 s žarjeno na 1050°C. Pov. 500x Fig. 9 Precipitation of secondary carbides in austenite, annealed 80 s at 1050° C. Magn. 500 x Kemična sestava: Chemical composition Sekunde 1_1 1 1 nun_1 1 ..........111111 Seconds 1 10 100 1000 CQS *" Minute 1_I 1 1 111111_l_ Time Minutes 1 10 Ure Hours SI. 8: Diagram izotermne destabilizacije avstenita (A — austenit, F — ferit, P — perlit, Ks — sekundarni karbidi) Fig. 8 Diagram of isothermal destabilization of austenite (A — austenite, F — ferrite, P — pearlite, Ks — secondary carbides) 800 SI. 10: Morfologija izločanja sekundarnih karbidov iz avstenita, 40 s žarjeno na 1050° C (matica je iz avstenita in martenzita) Fig. 10 Morphology of precipitation of secondary carbides from austenite, annealed 40 s at 1050" C (Matrix is of austenite and mart-ensite) s časom izotermnega žarjenja pa rastejo. Največji vpliv na ,rast sekundarnih karbidov ima temperatura in nad 1050 °C je njihova rast že zelo hitra. Inkubacijski čas za potek premene v perlitnem področju je daljši. Morfologija izločanja cementita je podobna kot pri izločanju sekundarnih karbidov, le izločanje cementita po prednostnih kristalografskih ravninah je manj izrazito. Sam potek transformacije je hitrejši kot proces destabilizacije. Na nekaterih mestih se vidi, da je transformacija potekla hitro po celem zrnu avstenita (si. 12). Oblika cementitnih lamel in medlame-larna razdalja v perlitu sta odvisni od temperature transformacije. Pri 750 °C je cementit grob in globula-ren, le na sredini večjih zrn je nakazana lamelama oblika. S padajočo temperaturo transformacije ima cementit vedno bolj lamelarno obliko (si. 13, 14). Najmanjšo SI. 12: Perlitna transformacija, 5 min žarjeno na 690° C. Pov. 500 x Fig. 12 Pearlitic transformation, annealed 5 min. at 690° C, Magn. 500 x SI. 11: Morfologija izločanja sekundarnih karbidov iz avstenita, 5 min žarjeno na 950° C (matica je iz avstenita in martenzita) Fig. 11 Morphology of precipitation of secondary carbides from austenite, annealed 5 min. at 950° C (Matrix is of austenite and martensite) Secondary carbides start to precipitate from austenite after a certain induction period, and this occurs on the grain boundaries betvveen the austenite grains, and on the boundaries of austenite grains vvith eutectic carbides. Tovvards the centre of austenite grains the precipitation is faster on certain crystallographic planeš (Figs. 9,10, and 11). In the initial phase of precipitation, the carbides are fine, but they grovv vvith the time of isothermal annealing. The greatest influence on the grovvth of secondary carbides has the temperature, and above 1050° C their grovvth is already very fast. Induction period for the transformation in the pearlitic region is longer. Morphology of cementite precipitation is similar to that of the secondary carbides, only precipitation of cementite on the preferred crystalogra-phic planeš is less pronounced. Transformation itself is faster than the process of destabilization. On same spots it is evident that the transformation vvas fast through the vvhole austenite grain (Fig. 12). The shape of cementite lamellae and the interlamellar spacing in per-lite depend on the transformation temperature. At 750° C cementite is coarse and globular, only in the centre of bigger grains lamellar formation is indicated. With de-creasing temperature of transformation the shape of cementite is becoming more lamellar (Figs. 13 and 14). The smallest interlamellar spacing in pearlite is found in the alloys at the transformation temperature around 650° C. At lovver temperatures pearlitic transformation is slovver, and lamellae can be observed only after longer annealing times. The curves of the initial precipitation of secondary carbides, and of pearlitic transformation were metallo-graphically exactly determined by optical microscopy and by SEM. Bigger problem vvas to determined the time vvhen both processes are completed. Precipitation of carbides and the pearlitic transformation move from the grain boundaries into the interior of the grains. The time SI. 13: Morfologija perlitne transformacije, 10 min žarjeno na 750°C Fig. 13 Morphology of pearlitic transformation, annealed 10 min. at 750" C SI. 14: Morfologija perlitne transformacije, 120 min žarjeno na 650° C Fig. 14 Morphology of pearlitic transformation, annealed 120 min. at 650° C medlamelarno razdaljo v perlitu ima zlitina v temperaturnem področju transformacije okoli 650 "C. Pri nižjih temperaturah poteka perlitna transformacija počasneje in lamele opazimo le pri daljših časih žarjenja. Krivulji začetka izločanja sekundarnih karbidov in perlitne transformacije smo lahko metalografsko točno določili z optičnim mikroskopom in v SEM. Večji problem je določiti čas, v katerem sta oba procesa končana. Izločanje karbidov in perlitna transformacija potekata s kristalnih mej v notranjost zrn. Čas, v katerem je proces končan, je zato odvisen od velikosti kristalnih zrn. V nekaterih primerih tudi sicer težko točno opredelimo konec procesa izločanja sekundarnih karbidov, ker v martenzitni osnovi težko ločimo karbidna zrna. Z meritvami trdote si prav tako težko pomagamo, saj se trdota, ko je izločenih že več kot 80 % sekundarnih karbidov, ali se perlitna transformacija približuje koncu, bistveno ne spremeni in so odstopanja v mejah merilnih napak. Na potek destabilizacije in perlitne transformacije pa vpliva tudi izcejanje legirnih elementov. Iz teh razlogov sta krivulji, ki označujeta konec obeh procesov, opredeljeni le približno. of the process termination thus depends on the size of crystal grains. In some cases the exact determination of the end of the precipitation of secondary carbides is dif-ficult since carbide grains can hardly be distinguished in the martensitic matrix. Measurements of hardness can also not help since the hardness changes very little when more than 80 % of secondary carbides are precipi-tated or the pearlitic transformation approaches to its end, and the deviations are in the limits of measuring er-rors. The destabilization process and the pearlitic transformation are influenced also by the segregations of al-loying elements. Therefore the curves determining the completion of both processes are approximate. In some samples being destabilized belovv 900° C also retained austenite was observed in the microstructure. Due to fast cooling (microstructure vvas stabilized by quenching) beside the stable austenite also residual austenite is present in the microstructure of the matrix in partial destabilization. The both austenites differ in the content of alloying elements (Cr, Mo, and C) (Fig. 15). The Ac3 point vvas determined dilatometrically. SI. 15: Delno destabiliziran avstenit (850° C, 2 min). V sredini zrn je stabilni avstenit. Ob kristalnih mejah, kjer so se izločili sekundami karbidi, je med martenzitnimi iglami zaostali avstenit. Pov. 500 x Fig. 15 Partially destabilized austenite (850° C, 2 min.). In the centre of grains there is stable austenite. On grain boundaries vvhere secondary carbides are precipitated there is residual austenite betvveen the martensitic needles. Magn. 500 x Pri nekaterih vzorcih, destabiliziranih pri temperaturah pod 900 °C, smo opazili v nikrostrukturi tudi zaostali avstenit. Zaradi hitrega ohlajanja (mikrostruktura smo stabilizirali z gašenjem) je v mikrostrukturi matice pri delni destabilizaciji prisoten poleg stabilnega avstenita še zaostali avstenit. Avstenita se razlikujeta po vsebnosti legirnih elementov Cr, Mo in C (si. 15). Temperaturo Ac3 smo določili z dilatometrom. Izotermni transformacijski diagram za destabiliziran avstenit. Temperatura destabilizacije 970 °C je istočasno izhodna temperatura nadaljnje toplotne obdelave. Omenili smo že, da višje razmerje Cr/C in legiranje z Mo zavirata transformacijo avstenita v perlit. Pri de-stabiliziranem avstenitu moramo upoštevati še vpliv sekundarnih karbidov. V primerjavi s TTT diagramom za nedestabiliziran avstenit je pri destabiliziranem avstenitu področje nastajanja perlita pomaknjeno močno v desno, v temperaturah pa ni nobene razlike. Več je v avstenitu izločenih sekundarnih karbidov, daljša je inkubacijska doba. Za praktično uporabo diagrama je seveda pomembna le popolna destabilizacija avstenita (si. 16). Določena razlika je v morfologiji nastajanja perlita. V nedestabiliziranem avstenitu poteka transformacija predvsem s kristalnih mej proti sredini avstenitnih zrn. Pri destabiliziranem avstenitu poteče premena hitro po celem, oziroma delu avstenitnega zrna. S časom žarje-nja narašča število transformiranih kristalnih zrn. (si. 17). Sekundarni karbidi delujejo kot kali in v destabiliziranem avstenitu poteka kontinuirna transformacija (si. 18). Tudi pri teh pogojih transformacije je iz že omenjenih razlogov nemogoče točno opredeliti konec premene. Destabilizacija: 970°C , 60min -Oestabilization 800,-1 ; MIHI-1 I ' "Hll- 1 ....... I 1 llllll-rn^p 1 10 Sekunde Seconds Čas —• Time 10* 103 10 10= ' i 'i.....i........"_i i i iiiiii 1 10 100 1000 Minute i i i...... i Minutes 1 10 Ure Hours Isothermal Transtormation Diagram tor Destabilized Austenite The destabilization temperature of 970° C is simul-taneously the starting temperature for further heat-treat-ment processes. It was already mentioned that higher Cr/C ratio and alloying vvith Mo retard the transtormation of austenite into pearlite. In destabilized austenite also the influence of secondary carbides must be taken in account. Compared to the TTT diagram for understabilized austenite the region of formation of pearlite in destabilized austenite is shifted significantly to the right while there are no differences related to the temperatures. Amount of secondary carbides precipitated in austenite is greater, longer is also the induction period. Only the complete destabilization of austenite (Fig. 16) is certain-ly important for practical application of the diagram. There is a certain difference in the morphology of pearlite formation. In undestabilized austenite the transtormation goes mainly from crystal boundaries tovvards the centre of austenite grains. In destabilized austenite the transtormation is fast over the vvhole or over a part of austenite grain. Longer annealing tirne increases the number of transformed crystal grains (Fig. 17). Secon-dary carbides are nuclei, and continuous transtormation takes plače in destabilized austenite (Fig. 18). Also in these conditions of transtormation it is not possible to determine exactly the termination of the transtormation due to the reasons already mentioned. Formation of cementite is influenced also by secon-dary carbides beside the temperature of isothermal transtormation. They are bigger in the bigger austenite grains, but their density is lovver. Thus the grovvth of cementite in the bigger grains is less hindered than in the smaller ones. Bainitic transtormation is possible only after the destabilization of austenite and at Cr/C ratios smaller than 5.2. By alloying vvith Mo, bainite can be obtained also at higher concentrations of Cr (1, 6). Induction period for the bainitic transtormation is long. Significant portion of bainitic phase in the microstructure is obtained only after longer times of isothermal annealing. A completely bainitic matrix can be expected only after a very long times of isothermal annealing when the alloy has a suit-able Cr/C ratio and is alloyed vvith Mo. SI. 16: Izotermni transformacijski diagram za destabiliziran avstenit (A — avstenit, K, — sekundarni karbidi, P — perlit, B — bai- nit) Fig. 16 Isothermal transtormation diagram for destabilized austenite (A — austenite, K — secondary carbides, P — pearlite, B — bainite) SI. 17: Perlitna transformacija destabilizirane avstenitne matice, 60 min žarjeno na 650° C. Pov. 500 x Fig. 17 Pearlitic transtormation of destabilized austenitic matrix, an-nealed 60 min. at 650° C. Magn. 500 x SI. 18: Meja med perlitnim zrnom (temnejše) in netransformiranim avstenitom (svetlejše) v destabilizirani matici Fig. 18 Boundary betvveen the pearlite grain (darker) and not trans-formed austenite (brighter) in the destabilized matrix Na izoblikovanje cementita vplivajo poleg temperature izotermne transformacije še sekundarni karbidi. Ti so v večjih avstenitnih zrnih večji, njihova gostota pa je manjša. Zato je v večjih zrnih rast cementita manj ovirana, kot v manjših. Bainitna transformacija je možna le po destabilizaciji avstenita in razmerjih Cr/C manjših od 5,2. Z legi-ranjem zlitin z Mo lahko dobimo bainit tudi pri višjih koncentracijah Cr (1,6). Inkubacijska doba za potek bainitne premene je dolga. Pomemben delež bainitne faze v mikrostrukturi dobimo le pri daljših časih izo-termnega žarjenja. Popolnoma bainitno matico pa lahko pričakujemo po zelo dolgih časih žarjenja zlitin z ustreznim razmerjem Cr/C in legiranih z Mo. Bainitno in martenzitno mikrostrukturo lahko ločimo le v SEM (17). Bainit stabilizira avstenit, zato je v mikrostrukturi matice še precej netransformiranega avstenita (si. 19). To so potrdile tudi meritve mikrotrdot. Začetek martenzitne transformacije smo določili z dilatometrom. Ms temperatura za popolnoma destabili-zirano zlitino (970 °C, 60 min) je 180 °C. Potek martenzitne transformacije ni odvisen le od delne ali popolne destabilizacije, temveč tudi od temperature destabiliza-cijskega žarjenja. Ms temperatura se znižuje z naraščajočo temperaturo destabilizacije (1, 6, 7). Pri kaljenju delno destabiliziranega avstenita je v matici poleg marten-zita tudi avstenit. Kontinuirni transformacijski diagram za destabiliziran avstenit Diagram, prikazan na sliki 20 velja za popolnoma destabilizirano zlitino (970 °C, 30 min). Preizkuse smo naredili tudi pri drugih temperaturah destabilizacije in na delno destabiliziranih vzorcih, da smo dobili čim več podatkov o vplivu različnih pogojev destabilizacije na mikrostrukturne značilnosti pri icontinuirnem ohlajanju. SI. 19: Morfologija bainitne transformacije v destabiliziranem avste-nitu (300° C, 4 ure). Mikrostruktura matice je iz avstenita, bai-nita, martenzita in sekundarnih karbidov Fig. 19 Morphology of bainitic transformation in destabilized austenite (300° C, 4 hours). Micostructure of matrix is of austenite, bai-nite, martensite, and secondary carbides Bainitic and martensitic microstructure can be distin-guished only by SEM (17). Bainite stabilizes austenite therefore a good deal of not transformed austenite can be found in the microstructure of the matrix (Fig. 19). This was confirmed also by the microhardness measure-ments. The beginning of the martensitic transformation was determined by dilatometer. Ms point for a completely destabilized alloy (970° C, 60 min.) is at 180° C. Course of martensitic transformation does not depend only on the partial or complete destabilization but also on the temperature of the destablization annealing. Ms temperature is lovvered vvith the increasing temperature of destabilization (1, 6, 7). After hardening the partially destabilized austenite, also austenite next to the martensite is found in the matrix. Continuous Transformation Diagram for Destabilized Austenite Diagram is presented in Fig. 20 and it is valid for a completely destabilized alloy (970° C, 30 min.). Experi-ments vvere made also at other temperatures of destabilization, and vvith partially destabilized samples in order to obtain the most possible data about the influences of various conditions of destabilization on the microstruc-tural characteristics in continuous cooling. In slovv cooling at 300° C/h, the transformation oc-curs at first in pearlitic stage, then also in the bainitic stage. The extent of transformation is in both stages ap-proximately equal. Martensitic transformation is namely observed, but amount of formed martensitic phase is small. At shorter times of destabilization the extent of transformation in pearlitic stage is smaller and in the bainitic one greater, but at temperatures around 200° C also martensitic transformation is observed. Destabilizacija 970°C, 30min - Destabilization 10 Sekunde Seconds Čas-► Time 1 Minute Minutes ' 600 0 1 500 8. 1400 300 100......... 1000 1 ......io Ure Hours SI. 20: Kontinuirni transformacijski diagram za destabiliziran avstenit (A — avstenit, K, — sekundarni karbidi, P — perlit, B — bainit) Fig. 20 Continuous transformation diagram for destabilized austenite (A — austenite, Ks — secondary carbides, P — pearlite, B — bainite) Pri počasnem ohlajevanju 300 °C/h pride do premene najprej v perlitni, nato pa v bainitni stopnji. Obseg transformacije je v obeh stopnjah približno enak. Mar-tenzitna premena se sicer opazi, martenzitne faze pa je nastalo malo. Pri krajših časih destabilizacije je obseg tranformacije v perlitni stopnji manjši, v bainitni pa večji, s tem da se pri temperaturah okrog 200 °C opazi tudi martenzitna premena. Pri večjih hitrostih ohlajanja (merjeno v sekundah ohlajanja med 800 in 500° C, oznaka t8/5) dobimo pri hitrosti t8/5 = 165 s že popolnoma martenzitno premeno. V začetku martenzitne premene smo opazili anomalijo (črtkana krivulja Ms). Podobni rezultati iz literature to anomalijo omenjajo, vendar brez ustrezne razlage. Ugotavljamo pa, daje anomalija pri krajših časih destabilizacije bolj izrazita. Naša preizkušanja so bila izvedena v omejenem obsegu, vendar se vidi, da ima čas destabilizacije, ki vpliva na obseg izločanja sekundarnih karbidov, bistveno vlogo na transformacijo avstenita pri kasnejšem ohlajevanju. Večja stopnja destabilizacije avstenita pospešuje obseg transformacije v perlitni stopnji. Anomalija pri martenzitni premeni, ki je večja pri manjši destabilizaciji avstenita, je verjetno povezana z nehomogenostjo avstenita. Ta je vsekakor večja pri nepopolni destabilizaciji. Premenske točke, ki so vrisane na diagramu, smo določili na podlagi dilatometrskih krivulj pri kontinuir-nem ogrevanju s hitrostjo 300 °C/h. Mikrostruktura matice je odvisna od hitrosti ohlajanja. Pri delni destabilizaciji je v sredini kristalnih zrn še avstenit. At higher cooling rates (measured in seconds for cooling from 800 to 500° C, marked by t8/5) a complete martensitic transformation is obtained at the rate t8/5 = 165 s. In the beginning of the martensitic transformation an anomaly vvas observed (dashed curve Ms). Similar results in references mention this anomaly but vvithout any explanation. It vvas found that the anomaly is more pronounced at shorter times of destabilization. Our testing vvas limited but it is evident that the tirne of destabilization vvhich influences the extent of secon-dary-carbide precipitation has an essential role in the transformation of austenite at further cooling. Higher stage of austenite destabilization accelerates the extent of transformation in the pearlitic stage. The anomaly in the martensitic transformation vvhich is greater at small-er destabilization of austenite is probably connected vvith the unhomogeneity of austenite. This is anyhow greater at uncomplete destabilization. Transformation points being plotted into the diagram vvere determined from dilatometric curves in continuous heating at the rate 300° C/h. The microstructure of the matrix depends on the cooling rate. In partial destabilization stili austenite is found in the centre of crystal grains. Residual austenite can be obtained at harsher condi-tions of hardening and at higher contents of Mn and Ni. Morphology of pearlite depends on the degree of destabilization as it vvas already explained at the isother-mal conditions of the transformation. Hardness of Alloys Corresponding applicable properties of vvhite chrom-ium čast irons depend on the amount of eutectic carbides in the microstructure, and on the hardness of ma-trix (6, 7). The best vvear resistance and the hardness possess the alloys vvith martensitic matrix. Alloys vvith martensitic-pearlitic matrix are softer, their hardness depends on the portion of pearlitic phase (11, 12). The alloy for vvhich the transformation diagrams vvere constructed exhibited as čast on the cross section čast in mould the average hardness 650 HV, and on the cross section čast into sand 635 HV respectively. Destabilized samples vvith martensitic matrix had hardnesses betvveen 760 and 800 HV. The highest hardness of alloys vvith 15 to 18 % Cr is obtained by hardening from the temperature betvveen 940 and 970° C. For the alloys vvith higher contents of Cr (over 20 %) the temperatures of hardening are higher, up to 1010° C. Martensite is harder if more carbon is dis-solved in austenite. Solubility of carbon increases vvith the increased temperature of hardening (18). By hardening from higher temperatures, also residual austenite can be obtained in the microstructure, especially if al-loys contain over 1 % Mn. Also Ni (Cu) acts like Mn vvhich on the other hand improves the through-hardena-bility. In our alloys the residual austenite vvas not detect-ed by dilatometric investigations. The residual austenite vvas obtained in the matrix only under harsher conditions of hardening from higher temperatures. The samples vvere hardened in air from the destabilization temperature of 970° C. Martensite is stable and the hardness starts to drop in tempering above 400° C (Fig. 21). The hardness of alloys vvith sufficient portion of residual austenite increases in tempering betvveen 450 and 550° C due to the decomposition of residual austenite into martensite. The transformation is connected vvith volume changes, and it is undesired be-cause of internal stresses. Microstructural characteris-tics of tempered čast iron are shovvn in Figs. 22 to 26. SI. 21: Vpliv temperature popuščanja na trdoto destabilizirane litine z različno mikrostrukturo matice (M — martenzit, B — bainit, P — perlit, Az — zaostali avstenit) Fi9' 21 intiuence of temperature of tempering on the hardness of de-stabilized čast iron vvith various microstructures of matrix (M — martensite, B — bainite, P — pearlite, Az — residual austenite) SI. 24: Mikrostruktura kaljene litine 2 uri popuščane na 600" C. Pov. 500 x Fig. 24 Microstructure of hardened čast iron, tempered 2 hours at 600° C. Magn. 500 x Zaostali avstenit lahko dobimo pri ostrejših pogojih kaljenja in pri višjih vsebnostih Mn in Ni. Morfologija perlita je odvisna od stopnje destabilizacije, kot smo to že razložili pri izotermnih pogojih transformacije. Trdota zlitin Ustrezne uporabne lastnosti belih kromovih litin so odvisne od deleža evtektičnih karbidov v mikrostrukturi in trdote matice (6, 7). Najboljšo obrabno obstojnost in trdoto imajo zlitine z martenzitno matico. Zlitine z martenzitno perlitno matico so mehkejše, njihova trdota pa je odvisna od deleža perlitne faze. (11, 12). Zlitina, za katero so narejeni transformacijski diagrami, ima v litem stanju na preseku ulitem v kokilo, povprečno trdoto 650 HV in na preseku ulitem v pesek 635 HV. Destabilizirani vzorci z martenzitno matico imajo trdoto med 760 in 800 HV. Največjo trdoto zlitin s 15 do 18% Cr dobimo pri kaljenju s temperature med 940 in 970 °C. Za zlitine z višjo vsebnostjo Cr (nad 20 %) so temperature kaljenja višje, do 1010 °C. Martenzit je trši, čim več C je raztopljenega v avstenitu. Topnost C narašča z naraščajočo temperaturo kaljenja (18). Pri kaljenju z višjih temperatur pa lahko dobimo v mikrostrukturi še zaostali avstenit, zlasti še, če vsebujejo zlitine nad 1 % Mn. Podobno kot Mn učinkuje tudi Ni (Cu), ki sicer izboljša prekalji-vost. V naših zlitinah zaostalega avstenita z dilatometr-skimi preiskavami nismo zasledili. Zaostali avstenit smo dobili v matici le pri ostrejših pogojih kaljenja z višjih temperatur. Vzorce smo kalili na zraku s temperature destabilizacije 970 °C. Martenzit je stabilen in trdota prične padati pri popuščanju nad 400° C (si. 21). Trdota zlitin z zadostnim deležem zaostalega avstenita pri popuščanju med 450 in 550 °C naraste zaradi razpada zaostalega avstenita v martenzit. Premena je povezana z volumskimi spremembami in je zaradi notranjih napetosti nezaže-ljena. Mikrostrukturne značilnosti popuščene litine so prikazane na slikah 22, 23, 24, 25 in 26. Fig. 23 Microstructure of hardened čast iron, tempered 2 hours at 400° C. Magn. 500 x SI. 22: Mikrostruktura destabilizirane litine kaljene na zraku. Matica je iz sekundarnih karbidov in martenzita. Pov. 500 x Fig. 22 Microstructure of destabilized čast iron hardened in air. Matrix is of seconday carbides and martensite. Mag. 500 x Si. 23: Mikrostruktura kaljene litine 2 uri popuščane na 400° C. Pov. 500 x 200 300 400 500 600 700 800 Temperatura popuščanja v °C Tempering temperature . °C I oM XP+B»M(20%P •P*M(907.P; □ P + M+Az- SI. 25: SEM posnetek na zraku kaljene litine. Mikrostruktura matice je iz sekundarnih karbidov in martenzita. Fig. 25 SEM picture of air-hardened čast iron. Microstructure of matrix is of secondary carbides and martensite. S kontinuirnim ohlajanjem smo pripravili vzorce z mešano perlitno-bainitno-martenzitno mikrostrukturo. Potek trdote v odvisnosti od temperature popuščanja je podoben kot pri vzorcih z martenzitno matico, le izhodna trdota je nižja. Trdota vzorcev s perlitno matico se pri popuščanju ne spremeni. Trdote ostalih preiskovanih zlitin z avstenitno matico v litem stanju so podobne, tiste z avstenitno-perlitno ali popolnoma perlitno matico pa so mehkejše in je njihova trdota med 480 in 550 HV. Temu ustrezne so tudi trdote po toplotni obdelavi. Trdota je pri zlitinah z avstenitno matico v litem stanju enaka po celem preseku preizkusnih valjčkov. Mikrotrdote posameznih mikrostrukturnih faz so podane v tabeli 4. Pri karbidih M7C3 moramo upoštevati, da je njihova trdota odvisna od kristalografske smeri, v kateri jo merimo (6). Tabela 4: Mikrotrdote mikrostrukturnih faz HV Primarni in evtektički karbidi M7C3 900-1300 Martenzit, sekundarni karbidi 650—700 Avstenit 400-520 Perlit, sekundarni karbidi 360—420 SI. 26: SEM posnetek kaljene litine 2 uri popuščane na 300° C Fig. 26 SEM picture of hardened čast iron, tempered 2 hours at 300° C By continuous cooling the samples vvith mixed pearl-itic-bainitic-martensitic microstructure vvere prepared. The variation of hardness depending on the tempering temperature is similar to that in the samples vvith mart-ensitic matrix, only the initial hardness is lovver. Hardness of samples vvith pearlitic matrix does not change in tempering. Hardnesses of the other investigated as čast alloys vvith austenitic matrix are similar to that of alloys vvith austenitic-pearlitic matrix, but the alloys vvith fully pearlitic matrix are softer and their hardnesses varied betvveen 480 and 550 HV. Similar relationship remains also after the heat treatment. Hardness of the as čast alloys vvith austenitic matrix is practically equal through the vvhole cross section of testing cylinders. Microhardnesses of single microstructural phases are given in Table 4. It is necessary to take into account that the hardness of M7C3 carbides depends on the crys-tallographic direction in vvhich it is measured (6). Table 4 Microhardnesses of Microstructural Phases Phase HV M7C3 primary and eutectic carbides 900 .. . 1300 Martensite, secondary carbides 650 .. . 700 Austenite 400 .. . 520 Pearlite, secondary carbides 360 .. . 420 ZAKLJUČEK Opisane so nekatere mikrostrukturne značilnosti in pogoji toplotne obdelave belih kromovih litin, legiranih z Mo in z Ni, Si, Mn, V, Ti, in W, namenjenih za centrifugalno dvoslojno lite valje. Iz preiskav se vidi, da so mehanske lastnosti odvisne od mikrostrukturnih značilnosti in s tem od kemične sestave, pogojev strjevanja in ohlajanja ter toplotne obdelave. Z ustreznim razmerjem Cr/C, legiranjem z Mo in hitrim strjevanjem dobimo v litem stanju drobne evtek- CONCLUSIONS Some microstructural characteristics and the condi-tions for heat treatment of vvhite chromium čast irons al-loyed vvith Mo, and vvith Ni, Si, Mn, V, Ti, and W vvhich are intended for centrifugal compound casting of rolls, are presented in the paper. The investigations show that mechanical properties depend on the microstructural characteristics, and thus on the chemical composition, conditions of solidification and cooling, and on the heat treatment. tične karbide in avstenitno matico, ki je s stališča nadaljnje toplotne obdelave najustreznejša. Sicer pa je matica lahko tudi avstenitno-perlitna ali perlitna. Delež karbidne faze je odvisen predvsem od vsebnosti C. Primarni heksagonalni karbidi, ki nastajajo pri počasnem ohlajanju ali pri deležu karbidne faze, večjem od 35 %, so nezaželjeni, ker bistveno poslabšajo žilavost litin. Po kemični sestavi ustrezajo karbidi ste-hiometričnemu razmerju od (Cr3Fe4) C3 do (Cr4Fe3) C3. Z naraščajočo vrednostjo razmerja Cr/C narašča vsebnost Cr v karbidih. Karbidi M7C3 z več Cr so trši. V ev-tektičnih karbidih je pri istem razmerju Cr/C manj Cr kot v primarnih karbidih. Trdota litin je odvisna od deleža evtektičnih karbidov in trdote mikrostrukture, ki jo dobimo po toplotni obdelavi. Osnova za toplotno obdelavo litin sta TTT diagrama za destabilizacijo avstenita in za destabiliziran avstenit in CTT diagram za destabiliziran avstenit. Morfologija izločanja sekundarnih karbidov (destabilizacija avstenita) in perlitne transformacije nedesta-biliziranega avstenita je podobna. Oba procesa potekata s kristalnih mej proti sredini kristalnih zrn, in to hitreje po določenih kristalografskih ravninah. V destabi-liziranem avstenitu poteka kontinuirna perlitna transformacija. Izločanje sekundarnih karbidov, ki poteka najhitreje med 940 in 990° C, je bistvenega pomena za nadaljnjo toplotno obdelavo litin. Le v destabilizirani matici je mogoča martenzitna in tudi bainitna transformacija. Bainitna transformacija poteka počasi in je za prakso manj pomembna. Najtrše so litine z martenzitno matico. Z ustrezno toplotno obdelavo pa lahko dobimo zlitine z martenzitno-perlitno ali perlitno matico, ki so mehkejše. Pri popuščanju litin z martenzitno ali martenzitno-perlitno matico prične trdota padati pri temperaturah popuščanja nad 400 °C. Suitable Cr/C ratio, alloying vvith Mo, and fast solidification enable the formation of fine eutectic carbides in the austenitic matrix in čast state vvhich is the most de-sired from the vievvpoint of further heat treatment. Matrix can also be austenitic-pearlitic or only pearlitic. Portion of carbide phase depends mainly on the car-bon content. Primary hexagonal carbides formed during slovv cooling or in alloys containing more than 35 % of carbide phase are undesired since they essentially re-duce the toughness of alloys. According to the chemical compositions the carbides correspond to stoichiometric ratios from (Cr3Fe4)C3 to (Cr4Fe3)C3. The increasing value of the Cr/C ratio causes the increased content of Cr in carbides. M7C3 vvith higher Cr content are harder. Eutectic carbides at the same Cr/C ratio contain less Cr than the primary ones. Hardness of čast irons depend on the amount of eutectic carbides and on the hardness of the microstructure obtained by the heat treatment. Basis for the heat treatment of čast irons are the TTT diagrams for the de-stabilization of austenite, and for destabilized austenite, beside the CTT diagram for the destabilized austenite. Morphologies of precipitation of secondary carbides (destabilization of austenite), and of pearlitic transformation of undestabilized austenite are similar. Both pro-cesses start on the crystal boundaries and proceed tovv-ards the centre of grains, and the process is faster along certain crystallographic planeš. In destabilized austenite continuous pearlitic transformation takes plače. Precipitation of secondary carbides vvhich is the fas-test betvveen 940 and 990° C is essential for further heat treatment of čast irons. Only in destabilized matrix the martensitic and also bainitic transformation is possible. Bainitic transformation is slovv and it is less important for practical applications. The hardest are the alloys vvith martensitic matrix. By a suitable heat treatment the al-loys vvith martensitic-pearlitic or pearlitic matrix can be obtained, but they are softer. In tempering the čast irons vvith martensitic or martensitic-pearlitic matrix, their hardness begins to decrease at the tempering tempera-tures above 400° C. LITERATURA/REFERENCES 1. F. Maratray, R. Usseglio-Nanot: Factors Affecting the Structure of Chromium and Chromium-Molybdenum White Irons, Climax Molybdenum, 1979 2. F. Maratray, R. Usseglio-Nanot: Atlas, Transformation Char-acteristics of Chromium and Chromium-Molybdenum VVhite Irons, Climax Molybdenum, 1979 3. Metals Handbook: Metallography, Structures and Phase Diagrams, vol. 8, American Society for Metals, 1973 4. R. S. Jackson: Journal of the Iron and Steel Institute, 1970, febr., 163—167 5. J. D. B. DeMello, M. Durand-Charre, S. Hamar-Thibault: Me-tallurgical Transactions, 1983, sept., 1793—1801 6. W. Fairhurst, K. Rohrig: Foundry Trade Journal, 1974 May 685—698 7. J. Drabina, A. Mazur: Giesserei Praxis, 1981, 6, 108—112 8. Centrifugally Čast High-chromium Rolls, Transactions ISIJ, 1986, pp. 168 9. T. Minemura, A. Inoue, T. Masumoto: Transactions ISIJ, vol. 21, 1981, pp. 649-655 10. J. T. H. Pearce, D. W. L. Elvvell: Journal of Materials Science Letters, 5, 1986, pp. 1063—1064 11. L. H. Priče: Metal Progress, 1983, aug., 21—27 12. R. W. Durman: Journal of Mechanical VVorking Technology, 8, 1983, 217-223 13. I. Katavič: Ljevarstvo (Foundry), 28,1981, 2, 3—6 14. J. Honda: Transactions ISIJ, vol. 24, 1984, 85—100 15. D. Kmetič et alt.: Development of čast irons for rolls vvith high content of chromium (Razvoj valjčnih litin z visoko vsebnostjo kroma) — Part I. Report of The Institute of Me-tallurgy, Ljubljana, 1983 16. D. Kmetič et alt.: Development of čast irons for rolls vvith high content of chromium (Razvoj valjčnih litin z visoko vsebnostjo kroma) — Part II. Report of The Institute of Me-tallurgy, Ljubljana, 1984 17. D. Kmetič et alt.: Development of čast irons for rolls vvith high content of chromium (Razvoj valjčnih litin z visoko vsebnostjo kroma) — Part III. Report of The Institute of Me-tallurgy, Ljubljana, 1985 18. F. Henke: Giesserei-Praxis, 1975, 23—24, 377—407 19. F. Mlakar, V. Tucič, B. Mlač: Optimal parameters in manu-facturing indefinite chill rolls (Optimalni parametri izdelave indefinite chill valjev). Report of The Institute of Metallurgy, Ljubljana, 1980 Osnovni koncept numerične simulacije radialnega kovanja Basic Concepts of Numerical Simulation of a Radial Forging Process T. Rodič*, D. R. J. Owen** UDK: 621.73.045:519.6 ASM/SLA: F22, Q24,1-66, U4g, U4k 1. UVOD S spoznanji splošnih principov fizikalne metalurgije1 postajajo preoblikovalni procesi v vročem pomembnejši. Preoblikovanje v vročem že dolgo ni več samo spreminjanje oblike preoblikovanca, temveč termomehanska obdelava materiala, ki naj privede do ugodnih strukturnih sprememb. Pri upoštevanju medsebojnih odvisnosti med strukturo, lastnostmi in obnašanjem materiala med plastičnim preoblikovanjem so deformacija, hitrost deformacije in temperatura tiste fizikalne večine, ki imajo odločilni vpliv. Nadzorovana porazdelitev teh termo-mehanskih parametrov med preoblikovanjem je potrebna pri optimiranju preoblikovalne operacije. Preoblikovalni procesi v vročem so zahtevni za eksperimentalna opazovanja zaradi visokih temperatur in preoblikovalnih hitrosti. Od tod tudi potreba po matematičnih in numeričnih modelih, ki pripomorejo k boljšemu razumevanju eksperimentalnih rezultatov ali celo delno nadomeščajo draga preizkušanja. Tri klasične metode za analizo preoblikovalnih procesov so bile pogosto uporabljane v preteklosti2: — metoda elementarne plastomehanike — metoda drsnih linij — metoda zgornje in spodnje meje. Analiza preoblikovalnih procesov je zahtevna in mnogo poenostavitvenih predpostavk je bilo vpeljanih v klasičnih metodah, da bi se izognili matematičnim težavam, kar je seveda zmanjševalo njihovo uporabnost. Napredek numeričnih metod v zadnjem času, posebej metode končnih elementov3 (MKE) in vzporedno zmanjševanje cen računalniških obdelav, ponuja možnost za realnejše simulacije preoblikovalnih procesov. 2. MATEMATIČNI MODEL Za analizo porazdelitev napetosti, deformacij in temperature, ki se spreminjajo znotraj deformacijske cone med preoblikovanjem, je nujna uporaba numeričnih metod. Razvoj MKE na področju plastomehanike4 in prenosa toplote ponuja zadovoljivo orodje za računalniško simulacijo preoblikovalnih procesov v vročem. Simulacijo preoblikovalnega procesa v vročem lahko idealiziramo5, kot je to prikazano na sliki 1. 1. INTRODUCTION Since the general principles of physical metallurgy were recognised1, the hot vvorking processes are no longer only concerned with shape changes but also con-sider the thermomechanical treatment vvhich contributes to beneficial structural changes vvithin the material. In considering the interaction betvveen the structure, properties and performance of the material under plastic deformation, the strain, strain rate and temperature are quantities vvhich have a fundamental influence and a controlled variation of these thermomechanical parame-ters is essential for optimising the forming operation. The nature of hot vvorking processes makes experimen-tal observations difficult, due to the high temperatures and speeds involved. Therefore mathematical and numerical models have a role to play in either improving the interpretation of experimental results or even replac-ing, in part, an expensive testing programme. Three classical methods for analysing metal forming problems have been widely used in the past2: — elementary plasticity — the slip line method — the upper and lovver bound method. Metal forming processes are complex and many sim-plifying assumptions have been introduced to these classical methods in order to avoid mathematical diffi-culties. This, hovvever, limits their applicability. Recent developments of numerical methods, in particular the Fi-nite Element Method3 (FEM), and a parallel reduction in unit computing costs offer an opportunity for a more realistic simulation of vvorking processes. 2. MATHEMATICAL MODEL The complex stress-strain and temperature distribu-tions vvhich vary across the deformed region during the deformation process require the use of numerical methods. Developments in the FEM in the field of plastome-chanics4 and heat transfer offer a satisfactory tool for computer simulation of hot vvorking processes. The numerical simulation of the hot vvorking process can be idealised5 as shovvn in Fig. 1. * FNT — Odsek za metalurgijo, Univerza E. Kardelja, Ljubljana ** Dept. of Civil Engineering, University of VVales, Swansea, U. K. Slika 1: Simulacija preoblikovalnega procesa. Vhodni podatki predstavljajo lastnosti in začetno stanje preoblikovanca, kot so: oblika, porazdelitev temperature, sestava in mikrostruktura materiala. Matematična simulacija preoblikovalnega procesa: Pri MKE razdelimo preoblikovanec na manjša območja, imenovana elementi. Togost vsakega elementa je določena z njegovo geometrijo in lastnostmi materiala, ki jih elementu pripišemo. Oba vpliva obravnavamo ločeno in zato relativno enostavno vgrajujemo različne materialne modele. Model preoblikovanca dobimo s sestavljanjem togostnih matrik elementov. Takšen pristop omogoča analizo različnih geometrijsko zahtevnih preoblikovalnih procesov. Z robnimi pogoji simuliramo različne pogoje, v katerih poteka proces. Izhodni rezultati. Po obdelavi rezultatov numerične analize določimo optimalne tehnološke pogoje. Nazoren grafični prikaz rezultatov je pomemben sestavni del analize z MKE. 3. MATERIALNI MODEL Pri modeliranju obnašanja materiala med plastično deformacijo je potrebno poznavanje ustrezne napetosti tečenja. V splošnem je ta odvisna od sestave in mikro-strukture materiala in od hitrosti deformacije, temperature ter deformacijskega stanja, povzročenega s plastično deformacijo. Napetost tečenja določimo z nateznim, tlačnim ali torzijskim preizkusom6. Pogosto napetosti tečenja niso dosegljive za specifične kombinacije termome- Slika 2: Ploskve napetosti tečenja v prostoru termomehanskih parametrov. Fig. 2: Flow stress surfaces in the space of the thermomechanical par-ameters. Fig. 1: Simulation of the vvorking process. Input data represent the material properties and in-itial state of the vvorkpiece; such as shape, temperature distribution, composition and microstructure of the material. Mathematical simulation of the vvorking process: In FEM the workpiece is divided into small regions termed elements. The stiffness of each element is determined by its geometry and material properties. Both effects are considered separately and therefore it is relatively sim-ple to incorporate different material models. The vvorkpiece model is obtained by combining the stiffness con-tribution of each element. Thus any complex shape of the vvorkpiece model can be analysed using the FEM. Boundary conditions are applied to simulate different conditions under vvhich the process is to operate. Output results: A decision on the most suitable set of operating conditions can be made by postprocessing the numerical results. Graphical representations play an integral part in the interpretation of the numerical results of a FEM analysis. 3. MATERIAL MODELS In the modelling of material behaviour during metal-vvorking processes a knovvledge of the appropriate flovv stress for the material is essental. In general this will de-pend on the composition and microstructure of the material and on the strain rate, temperature and deforma-tion modes imposed by the vvorking process. The flovv 0 0d=0-0p H \op O. > U (u > w 77777777777777 Slika 3: Osnovni enodimenzionalni elasto-viskoplastični reološki model. Fig.3: Basic one-dimensional elasto-viscoplastic rheological model. hanskih parametrov in jih ocenimo s pomočjo poznanih podatkov. Znane so različne interpolacijske enačbe, kot na primer Hajdukova7 ali Sellars-Tegartova enačba8. Na podlagi teh enačb lahko določimo potencial napetosti tečenja v prostoru deformacije, hitrosti deformacije in temperature5 (SI. 2). Ploskev A-B-C-D je določena z (p, 3F 3{# (c) and unstable and leading to failure if QCcp(C). Če pa je Kth=> [H]r = [H]" (6) At present, it is knovvn that almost ali the kinds of im-perfections in the crystal lattice of metals, both disloca-tions310, microvoids11,12, grain boundaries13, metal-car-bide interfaces14, and the surfaces of non-metalic inclu-sions15-16 can act as trapping sites. A mathematical model of hydrogen traps was deve-loped by Foster and McNabb17. According to the model, trapped hydrogen is locally in equilibrium vvith the inter-stitial hydrogen. The interaction betvveen hydrogen and the traps is a thermally activated process with an activa-tion energy constituted of the binding trap energy, and the activation energy of diffusion of hydrogen in an ideal iron lattice being 12 kJ/mole hydrogen18. The driving force for the interstitial diffusion of the dis-solved hydrogen in metals is the gradient of chemical potential. This gradient is influenced by the differences in hydrogen concentrations and by the effects of elastic-stress fields. The thermodynamic effect of the elastic-stress fields is caused by the reversible dilatation of the crystal lattice of metals, and the positive volume change accompanyed by the insertion of hydrogen interstitials into the areas of positive strain, vvhile compressively strained regions are impoverished vvith hydrogen. Thus a locally independent chemical potential of hydrogen in an inhomogeneous elastic-stress field is obtained through an inhomogeneous redistribution of hydrogen. The thermodynamic analysis of this process, under supposition that hydrogen is a completely mobile com-ponent, vvas established bi Li, Oriani and Darken19 who developed the follovving equation: Rh-^ + RT ln[H] — ohVH (1) The first two terms determine the chemical potential of hydrogen depending on its concentration, VH is the phenomenological partial atomic volume of hydrogen in iron, vvhile oh is the hydrostatic component of the stress tensor describing the triaxial state of stresses. Suppos-ing that the chemical potential of hydrogen in equilibrium is locally independent, the concentration of hydrogen in the region of the maximal hydrostatic component of stress tensor is given at a distance r from the notch root by (1): [H]r = [H]exp(ahVH/RT), (2) vvith [H] as an average concentration of hydrogen in the specimen. For a plane strain state, for a mode of loading I, and for the plane crack propagation from the notch root, the follovving equation is given20: CTh = 2 (1 4-v) K,/3 l/2jtr (3) From equations (2) and (3) vve obtain: [H]r = [H]exp2'1+;)^H (4) 3 Rl]/2nr The equation (4) connects the stress intensity factor K, vvith the concentration of hydrogen at the distance r from the notch root. The fact, that the delayed fracture does not occur if the applied stress intensity factor K, is lovver than the threshold stress intensity factor KTH, has been applied by Beachem21 as the criterion to determine the conditions under vvhich the delayed fracture takes plače: Kth Kth=> [H]r=[Hr Combining Eqs. (4) and (6) we obtain: Kth = 3 RT ]/27tr 2(1+v)VH In [H] (6) (7) The equation (7) is correct only vvhen the plastic zone at the notch tip is limited to less than the grain diameter d. Considering that also hydrogen on the grain bounda-ries at the crack tip is involved, vve can vvrite: 3RTJ^ |nW 2 (1 + v) VH l [H] ; d>Rh (8) vvith R, as the size of the strain plastic zone under mode of loading I. On the basis of equation (8) it is not possible to ex-plain the connection betvveen the yield strength effect and Kth, though it has been established by experiments. If taking into account the influence of a slip-line field at the notch root, according to Gerberich22, vve obtain for Kth: Kth = ~ in ctVH m [H] 2a (9) It is necessary to mention that disagreements are of-ten observed betvveen Eq. (9) and experimental results at yield strength belovv 1200 N mm-2. These disagreements vvere explained partly by the dependence of the [H]cr/[H] ratio on the yield point. Namely Farrell and Quarrell25 ascertained that larger concentrations of hy-drogen are needed to produce embrittlement in steel vvith lovver yield strength, and postulated the relationship [H]cr oc 1/oys. Kim and Loginovv26 suggested that the content of sol-uble hydrogen in steel vvas proportional to the yield strength, therefore [H]«ays. Finally, vve obtain: (10) [H]*_ ft [H] oys' vvith P as constant for a single type of steel. 3. EXPERIMENTS AND RESULTS 3.1 Selection of Steel and the Geometry of Speci-mens The aim of the investigation vvas to establish the influence of microstructure on the hydrogen induced sus-ceptibility to cracking of a high-strength steel vvith the composition: 0.40 % C, 0.31 % Si, 0.71 % Mn, 0.019 % P, 0.006 % S, 1.03 % Cr, 0.21 % Mo, 0.26 % Cu, 0.009 % Al and 0.010 % Sn. The steel vvas manufactured by the VOD process, thus the content of sulphur vvas low and the concentration of residual hydrogen did not exceed 0.05 ppm. Tensile tests vvere made on notched tensile speci-mens vvith the geometry shovvn in Fig. 1. For these specimens, the relationship betvveen the stress intensity factor K|, the geometry, and the axial force P is given by the eguation27: K| = 5^i (-1.27+1,72 D/d) (11) macije dosežene v samem korenu zareze, kjer se sicer pojavljajo prve mikrorazpoke. Slika 1 Geometrija cilindričnih nateznih preiskušancev z zarezo po obodu. Fig. 1 Geometry of cylindrical round notched tensile specimens. under condition that: 0,5 < d/D <0,8 The ratio p/D, vvith p as the notch root radius, vvas close to the value 0.02. Moran and Noriš28 found by the computer simulation of the tension test vvith cylindrical, peripherally notched specimens, that the maximum stresses at fracture occur at about two notch-root radii belovv the surface vvhen the p/D ratio is 0.01 to 0.02. On the other hand the maximum strain occurs at the notch root vvhere also the first microcracks appear. 3.2 Thermal Treatment Thermal treatment of specimens consisted of a 30 mins. austenitisation at 850° C, quenching in vvater or in oil, and tempering. Fig. 2 shovvs the microstructure of the lath-formed martensite in the steel quenched in vvater. After tempering 2 hrs. at 480° C or 420°C yield strengths of 1185 and 1290 N mm-2 respectively vvere obtained. The hardness of oil quenched specimens vvas betvveen 52 and 53.7 HRC. This means, that the predominantly martensitic microstructure of oil quenched specimens (Fig. 3) stili contains up to 3 % of bainite. After tempering specimens quenched in oil 2 hrs. at 450° C, a yield strength of 1230 N mm-2 vvas obtained. 3.2 Toplotna obdelava Toplotna obdelava preizkušancev je obsegala 1/2-ur-no avstčnitizacijo pri 850°C s kaljenjem v vodi oziroma olju ter popuščanje. Slika 2 prikazuje mikrostrukturo letvastega martenzita, izoblikovanega pri kaljenju v vodi. S popuščanjem 2 uri pri 480 oziroma 420° C je bila dosežena meja plastičnosti 1185 oziroma 1290 Nmm-2. Slika 2 Avstenitizirano pri 850°C in kaljeno v vodi. Letvasti martenzit. Fig. 2 Austenitized at 850° C, and quenched in vvater. Lath-shaped martensite. Slika 3 Avstenitizirano pri 850°C in kaljeno v olju. Spodnji bainit v mart-enzitni osnovi. Fig. 3 Austenitized at 850°C, and quenched in oil. Lovver bainite in the martensitic matrix. 3.3 Hydrogen Charging After thermal treatment, the specimens vvere charged vvith hydrogen by etching 24 hrs. in a 0.1 N aqueous so-lution of hydrochloric acid. Chemical analysis of specimens immediately after the removal from the acid solution shovved a hydrogen con-centration of 2.9±0.1 ppm independently upon the yield Trdota v olju kaljenih preizkušancev je dosegala vrednosti med 52 in 53,7 HRc. Pomeni, da je v pretežno martenzitni mikrostrukturi tovrstnih preizkušancev po kaljenju v olju (si. 3) še tudi do 3 % bainita. S popuščanjem v olju kaljenih preizkušancev 2 uri pri 450° C je bila dosežena meja plastičnosti 1230 Nmm"l 3.3 Navodičenje Toplotni obdelavi je sledilo navodičenje preizkušancev z jedkanjem 24 ur v 0,1 N vodni raztopini solne kisline. Kemične analize vzorcev neposredno po odstranitvi iz kisline kažejo, da dobljena koncentracija vodika 2,9 ±0,1 ppm praktično ni odvisna od meje plastičnosti jekla. Z drugimi besedami: s 24-urnim jedkanjem še ni dosežena stacionarna koncentracija nasičenja jekla z vodikom. Zaključujemo, da del pasti še ni zaseden, saj bi v takšnem primeru bila koncentracija vodika različna v preizkušancih z različno mejo plastičnosti26'29. 24 ur po odstranitvi iz raztopine je koncentracija vodika v preizkušancih padla na 0,82 ±0,1 ppm ter 48 ur po odstranitvi na 0,58 ±0,08 ppm. Ob predpostavki, da je hitrost uhajanja vodika iz cilindričnih preizkušancev majhnega premera (D je približno 7 mm) premo sorazmerna razliki med trenutno ter residualno koncentracijo vodika v njih, dobimo za koncentracijo vodika v odvisnosti od časa po jedkanju (v urah) naslednji izraz: [H] = 0,55+ 2,35 exp( —0,09 t) (12) Polempiričen izraz (12) je uporaben za fenomenološki opis uhajanja vodika iz cilindričnih preizkušancev in ugotovljeno je bilo30, da v navodičenih preizkušancih ostaja še okrog 0,55 ppm vodika tudi dolgo časa po končanem jedkanju. 3.4 Določevanje kritičnega ter mejnega napetostnega intenzitetnega faktorja Razvit je bil eksperimentalni sklop za registriranje inkubacijskega časa, to je časa do porajanja prve mikro-razpoke ter za registriranje počasnega napredovanja mikrorazpok do hipnega loma statično obremenjenih preizkušancev z zarezo po obodu. Sestavljen je bil iz polovičnega Wheatstonovega mostička z variabilnim uporom, ki ga je predstavljal uporovni listič, nalepljen preko ustja zareze. Porajanje ter napredovanje mikrorazpok je bilo registrirano posredno z odpiranjem ustja zareze, kot sprememba upornosti aktivnega uporovnega strength of steel. In the other vvords, a 24 hrs. etching did not produce the saturation of steel with hydrogen. It was concluded that ali the traps were not filled since in such a čase the concentration of hydrogen in steel vvould be different in samples with different yield strengths26 Twenty-four hours after removal from the acid solution, the concentration of hydrogen in speci-mens dropped to 0.82 ±0.1 ppm and after 48 hours to 0.58±0.08 ppm. Supposing that the escape rate of hydrogen from the cylindric specimens vvith small diameter (D is approx. 7 mm) is proportional to the difference betvveen the actual and the residual hydrogen concentration, the follovv-ing equation can be derived for the variation of hydrogen concentration vvith tirne (hours) after the removal of samples from the acid solution: [H] = 0,55+ 2,35 exp (— 0,09 t) (12) This semiempirical equation is useful for the pheno-menological description of hydrogen losses from the cylindric specimens and as it has been established30 that the specimens charged vvith hydrogen stili contain residual hydrogen of about 0.55 ppm even a long tirne after the etching. 3.4 Determination of the Critical and the Threshold Stress lntensity Factor An experimental set-up was developed for the regis-tration of the incubation period i. e. of the tirne neces-sary for the nucleation of the first microcrack, as vvell as for the registration of the slovv propagation of micro-cracks to the instantaneous fracture of the round-notched specimens under static load. It consisted of a half-Wheatston bridge vvith a variable resistor represent-ed by a strain-gauge sticked across the notch opening. Nucleation of microcracks and their propagation were indirectly registered by the displacement of the notch opening as the change of the resistance of the active strain-gauge compared vvith the resistance of the reference strain-gauge. This experimental set-up (schemati-cally shovvn in Fig. (4)) permitted to detect the propagation steps of about 0.1 um. An almost similar set-up vvas used to measure the critical stress intensity factor i. e. fracture toughness of steel. Fig. 5 shovvs hovv the measurements vvere made on the "Instron" tensile machine vvith an accurate exten-someter mounted on the notch opening of the specimen for calibration of the strain-gauges. Slika 4 Eksperimentalni' sklop za zasledovanje porajanja ter napredovanja mikrorazpok (1 — natezni preiskušanec z uporovnim lističem, 2 — referečni uporovni listič, 3 — merilna enota z izvorom napetosti, galvanometrom ter ojačevalcem, 4 — registrator). Fig. 4 Experimental set-up for the detection of crack nucleation and propagation (1 — tensile specimen vvith strain-gauge, 2 — reference strain-gauge, 3 — measuring unit vvith povver source, galvanometer and amplifier, 4 — recorder). Slika 5 Merjenje lomne žilavosti. Ekstenzometer na preiskušancu služi kalibriranju uporovnih lističev. Fig. 5 Measurement of fracture toughness. Extensometer on the ten-sile specimen used for the calibration of the strain-gauges. lističa glede na upornost referenčnega uporovnega lističa. Eksperimentalni sklop (shematsko prikazan na sliki 4) je dovoljeval zaznavanje koraka propagacije okrog 0,1 ji,m. Skoraj podoben sklop opreme je bil uporabljen za merjenje kritičnega napetostnega intenzitetnega faktorja, t. j. lomne žilavosti jekla. Na sliki 5 je prikazana izvedba merjenja na trgalnem stroju »Instron« s preciznim ekstenzometrom, montiranim preko ustja zareze preizkušanca ter uporabljenim za kalibriranje uporovnih lističev. Z merjenjem časa do loma preizkušancev v odvisnosti od uporabljene obremenitve je bil eksperimentalno določen mejni napetostni intenzitetni faktor KXH, to je mejna statična obremenitev, pri kateri še ne pride do nukleacije mikrorazpok. Kritični napetostni-intenzitetni faktor Klc, t. j. lomna žilavost jekla je bila izmerjena na cilindričnih preizku-šancih z zarezo ter utrujenostno razpoko v korenu zareze. Tako kot za izračun mejnega, je bila tudi za izračun kritičnega napetostnega intenzitetnega faktorja uporabljena formula (11), globina utrujenostne propagacije mikrorazpoke pa je bila izmerjena z optičnim mikroskopom po vsakokratnem preizkusu. Za preverjanje rezultatov je bila lomna žilavost izračunana še s korelacijo Rolfe-Novak31 za takoimenovano upper shelf področje. Odvisnost med izmerjenimi faktorji (Kth, K,c) ter mejo plastičnosti jekla je prikazana S -C* "E J3C O >4— E -z. C i- 4-> c C c <11 S C J) £ 1/1 i> o 1— c to o u č S KJ 4-. L. b •D C C O X I t— ^ "O E a? O) Z -C t- Slika 6 Odvisnost med napetostnim intenzitetnim faktorjem (KTH, l|C) in mejo plastičnosti preiskovanega jekla. Fig. 6 Relationship betvveen the stress intensity factor (KTH, K,c) and the yield strength of the investigated steel. The threshold stress intensity factor KTH vvhich repre-sents the limiting value of static load belovv vvhich micro-cracks do not appear, was experimentally determined by measuring the tirne till fracture occurs related to the ap-plied load. The critical stress intensity factor K,0 i. e. the fracture toughness of steel was measured on round notched tensile specimens vvith fatigue crack at the notch tip. The Eq. (11) was applied to calculate both the threshold and the critical stress intensity factor. The vvidth of the fatigue crack vvas measured vvith an optical microscope after each experiment. To check the obtained results, the fracture toughness vvas also calculated by the Rolfe-Novak correlation31 for the upper shelf region. The relation betvveen the measured factors (Kth, K,c) and the yield strength of the investigated steel is shown in Fig. 6. The plot presents also the relation by eq. (9) calculated on the basis of measured values for steel vvith fully martensitic microstructure after quenching (points M). Considering the equa-tion (10), a vaule of 5770 N mm-2 vvas calculated for the constant p. The straight line for [H]cr/[H] = const. proves that linear interpolation is acceptable at yield strengths above 1200 N mm"2. The threshold stress intensity factor KTH for tempered martensitic-bainitic microstructure vvith only fevv per-cents of bainite after quenching (point M + B) has the same value as that for a fully martensitic tempered microstructure vvith the same yield strength (KTH = = 2100 N mm-3'2). Hydrogen has no noticeable influence, if any at ali, on the fracture toughness of the investigated steel. How-ever, at the same yield strength, the fracture toughness, M'. TH] - konstx N o 0,05 ppm hydrogen j}0,55 ppm hydrogen K™-avHlnmr 2ot [H]rr„ 5770 - IhT 800 900 1000 1100 1200 1300 1400 Meja plastičnosti, Yield strength cSys(Nmm2) na sliki 6. V diagramu je vrisana tudi odvisnost (9), izračunana na osnovi izmerjenih vrednosti za jeklo s povsem martenzitno mikrostrukturo po kaljenju (točki M). Upoštevaje izraz (10) ima konstanta p vrednost 5770 Nmm-2. Premica za [H]cr/[H] = konst. dokazuje, da je pri meji plastičnosti nad 1200 Nmm-2 dopustna linearna interpolacija. Mejni napetostni intenzitetni faktor KXH za popušče-no martenzitno-bainitno mikrostrukturo z le nekaj odstotki bainita po kaljenju (točka M + B) ima enako vrednost, kot je bila določena za jeklo z mikrostrukturo popuščenega martenzita ter enako mejo plastičnosti (Kth = 2100 Nmm_3/2). Vodik nima večjega, če ima sploh kakšen vpliv na lomno žilavost preiskovanega jekla. Pri enaki meji plastičnosti pa je lomna žilavost, v nasprotju z mejnim napetostnim intenzitetnim faktorjem, nekoliko odvisna od majhnih mikrostrukturnih variacij jekla. Tako ima jeklo s popuščeno martenzitno-bainitno mikrostrukturo ter mejo plastičnosti 1230 Nmm-2 lomno žilavost med 2310 in 2360 Nmm~3/2, kar je nekoliko več od z linearno interpolacijo določene lomne žilavosti za popuščeno martenzitno mikrostrukturo enake meje plastičnosti (K]c med 2250 in 2320 Nmm"3/2). 3.5 Mikromorfologija prelomov Mikrofraktografske preiskave prelomnih površin na-vodičenih cilindričnih nateznih preizkušancev z zarezo po obodu, so bile opravljene z vrstičnim elektronskim mikroskopom. Slika 7 kaže del prelomne površine nateznega preizkušanca z zarezo, z mikrostrukturo popuščenega martenzita ter mejo plastičnosti 1185 Nmm-2. Statična Slika 7 Z zapoznelim lomom nastala prelomna površina preiskušanca z mikrostrukturo popuščenega martenzita ter mejo plastičnosti 1185 N mm-2 posneta z vrstičnim elektronskim mikroskopom. Fig. 7 Scanning electron micrographs of the delayed fracture surfaces of specimen vvith the tempered martensitic microstructure and yield strength 1185 N mm-2. contrary to the threshold stress intensity factor, de-pends slightly on small microstructure variations of steel. For instance, steel vvith the martensitic-bainitic microstructure vvith yield strength 1230 N mrrr2 has fracture toughness betvveen 2310 and 2360 N mm~3/2, vvhich is slightly above the value of K,c being betvveen 2250 and 2320 N mm"3'2 found by linear interpolation for the tempered martensitic microstructure vvith the same yield strength. 3.5 Micromorphology of Fractures Microfractographic investigations of fracture surfaces of round notched tensile specimens charged vvith hy-drogen were performed in a scanning electron micro-scope. Fig. 7 shovvs a part of fracture surface of a round notched tensile specimen vvith fully martensitic tempered microstructure and vvith yield strength 1185 N mm-2. Static load i. e. applied stress intensity factor 2190 N mm-3'2 vvas close to the limiting value (KTH = 2180 N mm-3'2) vvhich caused the delayed fracture of the specimen after 181 hours. Right along the notch (above), an area of slovv crack propagation can be seen separated by an unsharp boundary from the fracture surface formed by an instantaneous failure (belovv). The area of slovv crack propagation, vvhich is com-pletely undefined at low magnification, is shovvn in Fig. 8 at a higher magnification. This area is predominantly ductile and irregularly shaped dimples are found next to the well defined dimples. Irregular dimples indicate that decohesion occurred at a very lovv plastic deformation. It is also possible that some details are a consequence of cleavage too. A similar micromorphology of the area of slovv crack propagation is observed on samples vvith the tempered Slika 8 Področje počasne propagacije s slike 7. Pretežno duktilna oblika preloma. Fig. 8 The area of slovv crack propagation from Fig. 7. Predominantly ductile fracture. Slika 9 Področje počasne propagacije na preiskušancu s popuščeno martenzitno mikrostrukturo ter mejo plastičnosti 1290 N mm"2. Duktilna oblika preloma s posameznimi cepilnimi ploskvami. Fig. 9 The area of slovv crack propagation in specimen vvith the tem-pered martensitic microstructure and yield strength 1290 N mm"2. Ductile fracture vvith single cleavage facets. obremenitev, namreč aplicirani napetostni intenzitetni faktor 2190 Nmm~3/2, je bila blizu mejni vrednosti (KTH = 2180 Nmm-3/2), kar je povzročilo zapozneli lom preizkušanca po 181 urah. Neposredno ob zarezi (zgoraj) je moč videti cono počasnega napredovanja mikrorazpok, ne ostro ločeno od prelomne površine, nastale s hipno porušitvijo (spodaj). Cona počasnega napredovanja mikrorazpok, ki je pri nizki povečavi povsem neopredeljiva, je pri večji povečavi prikazana na sliki 8. To področje je pretežno duk-tilno, poleg dobro definiranih jamic pa najdemo tudi jamice nepravilnih oblik. Nepravilne jamice kažejo,, da se je dekohezija izvršila z zelo malo plastične deformacije in prav mogoče je, da so nekateri detajli tudi proizvod cepljenja. Podobno mikromorfologijo preloma v coni počasnega napredovanja mikrorazpok zasledimo tudi na preizkušancih s popuščeno martenzitno-bainitno mikrostrukturo, medtem ko je na preizkušancih s popuščeno martenzitno mikrostrukturo ter mejo plastičnosti 1290 Nmm~2 že tudi občutnejši delež cepilnih ali kvazicepilnih ploskvic (slika 9). Področje naglo zlomljenega osrednjega dela preizkušanca s slike 7 je pri večji povečavi prikazano na sliki 10. Prevladujejo področja duktilnega tipa preloma, čeprav je opaziti tudi cepilne oziroma kvazicepilne ploskvice, a v manjšem obsegu. Podobna je mikromorfologija preloma osrednjega, naglo zlomljenega dela preizkušancev s popuščeno martenzitno-bainitno mikrostrukturo, kot tudi preizkušancev s popuščeno martenzitno mikrostrukturo ter mejo plastičnosti 1290 Nmm"2, čeprav je v slednjem primeru število kvazicepilnih ploskvic povečano. Pojavljanje cepilnega oziroma kvazicepilnega tipa preloma v coni počasnega napredovanja mikrorazpok je le sporadično, v nasprotju s prevladujočo duktilno obliko preloma, zato sklepamo, da je nukleacija mikro- Sllka 10 Področje naglega loma iz slike 7. Duktilno, cepilno in kvazi-ce-pilno. Fig. 10 Region of fast fracture from Fig. 7. Ductile, cleavage and quasi-cleavage. martensitic-bainitic microstructure, vvhile a noticeable amount of cleavage or quasi-cleavage facets appear in the samples vvith the martensitic microstructure and the yield strength 1290 N mm"2 (Fig. 9). The area of fast fracture, already shovvn in Fig. 7, is shovvn again in Fig. 10 at a higher magnification. Here, the ductile type of fracture prevails though cleavage or quasi-cleavage facets can also be observed but in small-er extent. A similar micromorphology of the areas of fast fracture is also observed on the samples vvith the tempered martensitic-bainitic microstructure as vvell as on the samples vvith the tempered martenistic microstructure and vvith yield strength 1290 N mm-2, though in the lat-ter čase the number of quasi-cleavage facets is larger. Cleavage or quasi-cleavage type of fracture in the area of slovv crack propagation is merely sporadic in comparison to the prevailing ductile type of the fracture, thus the conclusion can be made that the crack nucieation as vvell as the slovv crack propagation are mainly strain induced processes related to the decrease of fracture ductility i. e. decrease of the microplasticity in the area of stress induced segregation of hydrogen at the crack tip. 4 CONCLUSIONS An appropriate method vvas developed for the detec-tion of the nucieation and the propagation of micro-cracks at the notch tip of hydrogen charged static loaded cylindrical round notched tensile specimens. The limit of the detectability of microcrak propagation vvas about 0.1 (im. Measurements of the threshold stress intensity factor Kth of the hydrogen charged chromium-molybdenum Č.4732 steel vvith the tempered martensitic microstruc- razpok ter njih počasno napredovanje deformacijsko induciran proces povezan s poslabšanjem lomne duktil-nosti, t. j. poslabšanjem mikroplastičnosti v področju napetostno induciranega segregiranja vodika ob konici razpoke. 4. ZAKLJUČKI V okviru opravljenega dela je bila razvita primerna metoda za preučevanje nastajanja ter napredovanja mi-krorazpok iz korena zareze na obodu navodičenih ter statično obremenjenih cilindričnih nateznih preizkušan-cev z zarezo. Najmanjši korak propagacije mikroraz-pok, ki ga je bilo moč zaslediti, je znašal okoli 0,1 ^m. Merjenja mejnega napetostnega intenzitetnega faktorja Kth navodičenega krom-molibdenskega jekla, vrste Č.4732, z mikrostrukturo popuščenega martenzita oziroma popuščeno martenzitno-bainitno mikrostrukturo enake meje plastičnosti, kažejo, da majhne mikro-strukturne variacije preiskovanega jekla ne vplivajo na k-th- Ti rezultati se ujemajo z ugotovitvami Nakasata in Terasakija32, ki potrjujeta, da mejni napetostni intenzi-tetni faktor KIscc pri enaki trdnosti jekla ni odvisen od mikrostrukturnih variacij visokotrdnega jekla. Če je ta ugotovitev splošna, potem je utemeljena hipoteza, po kateri je nukleacija mikrorazpok, ki povzroče zapozneli lom jekla, vedno omejena na martenzitne dele mikrostrukture (najprej dosežena [H]cr). Le na ta način namreč lahko razložimo, da majhni deleži bainita v pretežno martenzitni mikrostrukturi popuščenega visokotrdnega jekla nimajo vpliva na mejni napetostni intenzite-tni faktor. Zdi se, da se različna jekla pri enaki vsebnosti vodika ter enaki meji plastičnosti v pogledu nuklea-cije mikrorazpok obnašajo kot elastični kontitfuum. Merjenja kritičnega napetostnega intenzitetnega faktorja kažejo, da majhne koncentracije vodika v preiskovanem jeklu nimajo opaznejšega vpliva na lomno žila-vost jekla. Pač pa je pri enaki meji plastičnosti lomna žilavost (v nasprotju z mejnim napetostnim intenzite-tnim faktorjem) odvisna tudi od majhnih mikrostrukturnih variacij jekla, saj ima jeklo s popuščeno martenzitno-bainitno mikrostrukturo nekoliko višjo lomno žilavost, kot isto jeklo s popuščeno martenzitno mikrostrukturo enake meje plastičnosti. Rezultati teh preiskav se ujemajo s podatki, ki so jih objavili Ohtani, Terasaki in Kunitake3 , ki so povečano žilavost duplex mikrostrukture razlagali s koristno vlogo majhnih deležev bainita pri zmanjšanju delov posameznih avstenitnih zrn, ki se transformirajo v martenzit. Takšna mikrostruktura ima povečano odpornost proti napredovanju razpok, t. j. manjšo občutljivost k zapoznelemu lomu. Porajanje mikrorazpok v področju maksimalnih deformacij, kot tudi pretežno duktilna oblika preloma v coni počasnega napredovanja mikrorazpok navajata k sklepu, da je nukleacija mikrorazpok deformacijsko induciran proces, povezan s poslabšanjem lomne duktil-nosti med trajanjem obremenjevanja. ture or the tempered martensitic-bainitic microstructure and with the same yield strength shovved that small microstructure variations of the investigated steel had no influence on Kth. These results confirm the Nakasato's and Terasaki's statements32 according to vvhich the threshold stress in-tensity factor K|SCC at the same tensile strength does not depend on the microstructure variations of high strength steel. If this statement is general, the hypothesis sug-gesting that the microcrack nucleation leading to the de-layed fracture is always confined to martensitic areas of the microstructure ([H]cr is at first reached) has argument. It seems that this is the only way to explain the lack of influence of small portion of bainite in a predomi-nantly martensitic microstructure of the tempered high strength steel on the threshold stress intensity factor. As far as the crack nucleation is concerned, it seems that various steels vvith the same yield strength and the same hydrogen concentration behave as an elastic con-tinuum. The measurements of the critical stress intensity factor show that small concentrations of hydrogen in the investigated steel has no noticeable influence on the fracture toughness of steel. Hovvever, at the same yield strength the fracture toughness (contrary to the threshold stress intensity factor) depends also on small microstructure variations of steel, since steel vvith the tempered martensitic-bainitic microstructure has slightly higher fracture toughness than the same steel vvith the tempered martensitic microstructure and vvith the same yield strength. The results of this investigation agree vvith the data published by Ohtani, Terasaki and Kunitake33 who explained the higher toughness of the duplex microstructure by the beneficial effect of small quantity of bainite vvhich reduces the size of single austenite-grain parts in vvhich the martensitic transformation takes plače. Such a microstructure has a better resistance to crack propagation i. e. it is less sensitive to the delayed fracture. The nucleation of the crack in the region of maximal strain as well as the predominantly ductile type of fracture in that area suggest the conclusion that the nucleation of microcracks is a strain induced process related to the decrease of fracture ductility during the loading. LITERATURA/REFERENCES 1. G. L. Hanna, A. R. Troiano in E. A. Steigerwald: Transacti-ons of the ASM, 57, 1964, 658-671 2. A. R. Troiano: Transactions of the ASM, 52, 1960, 54 3. L. S. Darken in R. P. Smith: Corrosion 5,1949,1 4. R. A. Oriani: Acta Metali. 18, 1970,147 5. V. Sakamoto in J. Eguchi: Proc. Japan Congress on Materials Research 19,1976, 91 6. A. P. Miodovvnik: Stress corrosion cracking and hydrogen embrittlement of-lron base Aloys, NACE, Huston, 1977 7. A. Zielinski, B. Lunarska in M. Smialovvski: Acta Metali. 25, 1977, 551 8. A. J. Kumnick in H. H. Johnson: Acta Metali. 28, 1980, 33 9. A. M. Adair in R. E. Hook: Acta Metali. 10,1962, 741 10. A. J. Kumnick in H. H. Johnson: Metallurgical Transactions 5A, 1974, 1199 11. G. M. Evans in E. C. Rollason: Japan Iron Steel Inst., 1969, 1484 12. D. M. Allen-Booth in J. Hevvitt: Acta Metali. 22, 1974,'171 13. H. Hargi, V. Hayashi in L. L. Shreir: Corrosion Science 11, 1971,25 14. G. W. Hong in J. Y. Lee: J. Mat. Sci. 18, 1983, 271 15. T. Asaoka, G. Lapasset in M. Aucouturier: Corrosion 34, 1978, 39 16. G. M. Pressouyre in I. M. Bernstein: Metallurgical Transactions 9A, 1978, 1571 17. A. McNabb in P. K. Foster: Trans. Am. Inst. Min. Engrs. 227, 1963, 618 18. W. Tyson: Canadian Metallurgical Quarterly, Vol. 18, 1979, 1-11 19. J. C. M. Li, L. S. Darken in R. A. Oriani: Zeitschrift fur Physt-kalische Chemie Neue Folge 9,1966, 271 20. H. L. Ewalds in R. J. H. Wanhill: Fracture mechanics, Edvvard Arnold Ltd., 1985 21. povzeto po A. G. Guy: Essentials of Materials Science, McGraw-Hill, 1976 22. W. W. Gerberich: Effect of hydrogen on high-strength and martensitic steels, Hydrogen in Metals — Proceedings of an international conference, 23,—27. September 1973, Seven Springs Conference Center, Champion, Pa. USA, Library of Congress Catalog Card Number: 73— 86455 23. R. A. Oriani: Ber. der Buns. Gesell. 76, 1972, 848 24. C. D. Beachem: Metallurgical Transactions 3, 1972, 437 25. K. Farrell in A. G. Ouarrell: J. Iron Steel Inst., 202, 1964, 1002 26. C. D. Kim in A. W. Loginovv: Corrosion 24,1968, 313 27. K. Heckel: Einfiihrung in die technische Anvvendung der Bruchmechanik, Carl Hanser Verlag, Munchen 1970 28. B. Moran in D. M. Noriš: Metallurgical Transactions A, 1978, 1685 29. L. S. Darken in R. P. Smith: Corrosion 5, 1949, 60 30. B. Ule: Zapozneli lom zaradi vodika v jeklu, Magistrsko delo, Univerza E. Kardelja v Ljubljani, 1987 31. S. T. Rolfe in S. R. Novak: Slow-bend K,c testing of medium-strength high-toughness steels, STP 463, 1970, 124—159 kot tudi R. B. Scarlin in M. Shakeshaft: Metals Technology, Jan. 1981, 1—9 32. F. Nakasato in E. Terasaki: Transactions ISIJ, 15, 1975, 290—291 33. H. Ohtani, F. Terasaki in T. Kunitake: Transactions ISIJ, 12, 1972,185 Doktorska in magistrska dela v letu 1986 Ph. D. and M. Sc. Theses at the Department of Geology, Mining and Metallurgy, Section Metallurgy, in Year 1986 DOKTORSKO DELO Mirko Dobršek: Konstrukcija trokomponentnih sistemov Pd-Au-Zn, Pd-Cu-Zn, Au-Cu-Zn (Mentor: I. Kosovinc, 18/12/1986) Zgoraj omenjeni sistemi so osnova sistema Pd-Au-Ag-Cu-Zn. Paladij zaradi svoje majhne gostote, korozijske obstojnosti, katalizatorskih sposobnosti ter nizke cene vedno uspešneje nadomešča drago zlato in platino. Avtor je izdelal neizotermne preseke teh ternernih sistemov, določil fazna področja z rentgensko fazno analizo, ki jo je dopolnil še z metalografsko analizo. V vseh treh sistemih je ugotovil široko enofazno področje ternerne trdne raztopine, ki se širi iz obrobnih sistemov popolne topnosti v ternerni prostor. Delo ima praktično uporabno vrednost predvsem na področju dentalnih zlitin in omogoča nadaljnje raziskave večkomponentnih sistemov v ožjih, tehnično-ekonomsko zanimivih koncentracijskih območjih. 82 strani 29 cit. PH. D. THESIS Mirko Dobršek: Construction of Ternary Pd-Au-Zn, Pd-Cu-Zn, and Au-Cu-Zn Phase Diagrams (Supervisor: I. Kosovinc, 18/12/1986) The upper mentioned systems are the basis of the Pd-Au-Ag-Cu-Zn system. Paladium due to its lovv densi-ty, corrosion resistance, catalytic properties, and lovv priče successfully substitutes more expensive gold and platinum. Nonisothermal cross sections in these ternary diagrams vvere constructed. The phase regions vvere de-termined by the X-ray phase analysis vvhich was comple-mented also vvith the metallographic analysis. In ali the three systems a wide one-phase region of the ternary solid solution was found vvhich extends from the edge binary systems of complete solubility into the'-ternary space. The project has practical applicability maihly in the field of dental alloys, and it enables further investiga-tions of the multi-component systems in the narrovver, technically and economically interesting regions of com-positions. 82 pages 29 ref. MAGISTRSKA DELA Mihael Tolar: Kohezivne cone v plavžu (Mentor: J. Lamut, 4/3/1986) Plavž je agregat, v katerem potekajo fizikalno-ke-mične reakcije v odvisnosti od razporeditve plinskih tokov ter temperaturnega polja. Lega in oblika kohezivne cone v delovnem prostoru plavža sta odločilni pri razporeditvi plinskih tokov iz spodnjega v zgornji del peči. Avtor je v svojem delu obravnaval termostabilnost mineralnega vsipa in obnašanje koksa po višini delovnega prostora peči, prehod silicija in žvepla iz vročega koksa v grodelj, oblikovanje mehčalno-talilne cone v peči ter kohezivne cone, skupaj z ukrepi, ki vplivajo na njihovo obliko in lego. Kohezivne cone so kompaktne nataljene plasti mineralnega vsipa, ki so vrinjene med plasti koksa in jih med seboj ni možno več ločiti. Predstavljajo prehod med začetkom mehčanja in dokončno stalitvijo vsipa. Običajno imajo te cone obliko V ali W. Obnašanje vsipa pri pogrezanju je avtor zasledoval z jemanjem vzorcev na 7 ravneh v plavžu. Ugotavljal je stopnjo redukcije in metalizacije, delež FeO ter nastajanje žlindre. Hod preizkovanega jeseniškega plavža je bil periferen. 110 strani 173 cit. M. SC. THESES Mihael Tolar: Cohesive Zones in the Blast Furnace (Supervisor: J. Lamut, 4/3/1986) Course of physico-chemical reactions in the blast furnace depends on the distribution of gas flovvs and the temperature field. Position and shape of the cohesive zones in the operation area of the blast furnace are es-sential for the distribution of the gas flovvs from the lovv-er into upper section of the furnace. Further, termostability of the burden, behaviour of coke along the height of the furnace region of operation, transfer of silicon and sulphur from hot coke into pig iron, formation of the softening-melting zone in the furnace and the cohesive zones are treated together vvith the measures vvhich can influence the shape and the position of the cohesive zones. These zones are compact partially melted layers of the burden vvhich intrude into the coke layers, and they cannot be separated anymore. They represent the transition betvveen the initial soften-ing and the final melting of the burden. Usually these zones have the shape of letter V or W. Behaviour of the burden during descending was follovved by sampling on 7 levels in the furnace. Degrees of reduction and metalli-zation, portion of FeO and formation of slag vvere ana-lyzed. Running of the blast furnace in the Jesenice Iron-vvorks vvas peripheral. 110 pages 173 ref. Miraš Djurovič: Vpliv induktivnega mešanja taline na kinetiko reakcij med talino in žlindro pri izdelavi jekla za kroglične ležaje Č.4146. (Mentor: J. Lamut, 14/5/1986) Moderni postopki izdelave elektrojekla vključujejo tudi obdelavo taline izven peči. Kroglični ležaji zahtevajo zelo kakovostno jeklo. V delu je bil . poudarek na študiju vpliva induktivnega mešanja. Raziskave so bile v jeklarni Železarne Boris Kidrič v Nikšiču. Avtor je najprej obdelal osnovne značilnosti ležajnega jekla Č.4146, osnove reakcij med talino in žlindro ter izven-pečno obdelavo tekočega jekla s poudarkom na intenzi-ftkaciji reakcij. Zasledoval je gibanje žvepla v jeklu med obdelavo v napravi ASEA-SKF, vpliv sestave sintetične žlindre, pomen delovne temperature, sestavo dobljene žlindre in možnost reoksidacije. Na osnovi rezultatov raziskav je avtor predložil tehnologijo za izdelavo ležajnega jekla v obločni peči skupaj z rafinacijo v ASEA-SKF napravi in litjem ingotov. 73 strani 9 cit. Miraš Djurovič: Influence of Inductive Stirring of Melt on the Kinetics of the Melt/Slag Reactions in Manufacturing Č.4146 Ball-Bearing Steel (Supervisor: J. Lamut, 14/5/1986) Modern processes in manufacturing steel in electri-cal furnaces include also the treatment of melt outside the furnace. Bali bearings demand a high-grade steel. Influence of the induction stirring was analyzed. Investigations were made in the Boris Kidrič Ironvvorks in Nikšič. Initially, the basic characteristics of the ball-bearing Č.4146 steel are presented, together vvith the basic reactions betvveen melt and slag, and the out-of-furnace treatment of molten steel vvith a special emphasis on the intensification of reactions. Variation of the sulphur content in steel during the treatment in the ASEA-SKF equipment vvas follovved, together vvith the influence of the composition of synthetic slag, and the importance of operating temperature, composition of the obtained slag, and the possiblity of the reoxidation. Based on the results of the investigation an improved technology for manufacturing ball-bearing steel in electric are furnace together vvith the refining in the ASEA-SKF set-up, and casting the ingots was proposed. 73 pages 9 ref. Henrik Kaker: Kvantitativna energijsko disperzijska analiza Nimonic 80 A v rastrskem elektronskem mikroskopu (Mentor: V. Marinkovič, 8/7/1986) Avtor je na nikljevi zlitini Nimonic 80 A obravnaval postopke obdelave zbranega spektra z energijskim spektrometrom, metode kvantitativne mikroanalize in računske postopke za popravke v koncentracijah analiziranih elementov vsled razlik v atomskem številu med vzorcem in standardom (etalonom), absorpcije in sekundarne fluorescence rentgenskega sevanja. Naredil je tudi primerjavo med metodo s standardi in brez njih ter dobljene rezultate primerjal z rezultati kemijske analize. Analiziral je napake, ki vplivajo na točnost kvantitativne mikroanalize z energijskim spektrometrom. Ugotovil je, da je relativna napaka pri metodi brez standardov pod 1 % pri koncentracijah nad 20 mas. %, naraste pa na okoli 16 % pri koncentracijah pod 1 mas. %. Glavna prednost metode je njena hitrost. Zelo primerna je, če nimamo na razpolago ustreznega standarda in če točnost pri nizkih koncentracijah ni odločilna. Metoda s standardi je pri majhnih koncentracijah sicer bolj natančna, a je bistveno počasnejša. Mikroanaliza faz v zlitini Nimonic 80 A je pokazala prisotnost Ti karbonitri-dov in Cr karbidov na kristalnih mejah ter enakomerno porazdeljeno fazo y' v osnovi. 74 strani 43 cit. Henrik Kaker: Ouantitative Energy Dispersion Analysis of Nimonic 80 A by the Scanning Electron Microscope (Supervisor: V. Marinkovič, 8/7/1986) The author has chosen the Nimonic 80 A nickel alloy to present the methods for treating the spectrum obtained by the energy spectrometer, the methods of quantitative microanalysis, and the mathematical methods to correct the obtained concentrations of analyzed elements due to the differencies in atomic numbers betvveen the sample and the standard, in the absorption, and the secondary fluorescence of X-radiation. A com-parison vvas made betvveen the method vvhere stand-ards vvere applied and the method vvithout applying standards. The both obtained results vvere compared to the results of chemical analysis. Further, the analysis of errors influencing the accuracy of quantitative microan-alysis by the energy spectrometer vvas made too. It vvas found, that the relative error vvith the method vvhere standards vvere not applied vvas belovv 1 % for concentrations above 20 mass % of elements, but this error is increased to about 16% if concentrations are reduced belovv 1 mass %. The basic advantage of this method is that it is fast. It is suitable when a corresponding standard is not available and vvhen the accuracy at lovver concentrations is not essential. The method applying standards is more ccurate for lovver concentrations but it is essentially slovver. The microanalysis of phases in the Nimonic 80 A alloy revealed the presence of Ti carboni-trides and Cr carbides on the grain boundaries, and the uniformly dispersed y' phase in the matrix. 74 pages 43 ref. o Železarski zbornik, 21, 1987, 1—4 1. KRONOLOŠKO KAZALO Todorovič Gojko, J. Lamut, B. Dobovišek, J. Kramer, J, Zapu-šek, B. Sedlar: Naogljičenje železa med redukcijo in taljenjem plavžnega vsipa........ŽZB 21 (1987) 1,1—5 Arh Joža, V. Prešeren: Dosežki Železarne Jesenice na področju sekundarne obdelave jekla . . ŽZB 21 (1987) 1, 7—17 Vodopivec Franc, M. Gabrovšek: Meja plastičnosti konstrukcijskih jekel, fizikalno metalurške osnove ŽZB 21 (1987) 1,19—28 Vodopivec Franc, F. Marinšek, F. Grešovnik: Rekristalizacija in rast zrn pri žarjenju hladno valjanega jekla z 0,03 % C, 1,8 % Si, 0,3 %Mn in 0,3% Al......ŽZB 21 (1987) 1, 29—37 Kaker Henrik: Mikroanaliza faz v zlitini Nimonic 80 A s kombinacijo REM-EDS .......ŽZB 21 (1987) 1,39-43 Ažman Alojz, D. Slkošek, A. Šteblaj, J. Triplat, J. Arh: Tehnična novica: Novi konstrukcijski mikrolegirani jekli Niomol 390 in Niomol 490 ..........ŽZB 21 (1987) 1,45—52 Leskovšek Vojteh: Tehnična novica: Predstavitev enokomorne vakuumske peči IPSEN VTC 324-R s homogenim plinskim hlajenjem pod visokim tlakom .... ŽZB 21 (1987) 1, 53—57 Arh Jože: Tehnična novica: Druga evropska konferenca o elek- tro jeklarstvu........ŽZB 21 (1987) 1, 59—63 Paulin Andrej, J. Lamut, D. Dretnik: Reaktivnost koksa in njen vpliv na delo plavža......ŽZB 21 (1987) 2, 65—71 Gnamuš Janko, G. Rlhar: Reparaturno varjenje orodnih jekel ... ŽZB 21 (1987) 2, 73-76 Grešovnik Ferdo: Računanje temperaturnih napetosti v elastičnem področju .......ŽZB 21 (1987)2, 77—83 Brudar Božidar: Ohlajanje jeklenega valja na vozičku . . . . . . ŽZB 21 (1987) 2, 85—92 Vodopivec Franc, O. Kurner, A. Lagoja, F. Grešovnik, A. Rodič, S. Senčič.F. Vizjak:Tehničnanovica:Orazvojnihmožnostihjekelin nekaterih posebnih zlitin ter postopkov za njihovo izdelavo, predelavo, ulivanje in plemenitenje .... ŽZB 21 (1987) 2, 93—98 Paulin Andrej: Mehanizem zgorevanja koksa . . . . . . ŽZB 21 (1987) 3, 105-112 Vodopivec Franc, F. Grešovnik, F. Marinšek, M. Kmetič, O. Kurner: O teksturi valjanja, razogljičenja in rekristalizacije v jeklu z 0,03 C, 1,8 Si in 0,3 Al . . ŽZB 21 (1987) 3,113—118 Uranc Franc: Vpliv trenja, poti, drsne hitrosti in pritiska na obrabo.........ŽZB 21 (1987) 3, 119—125 Risteski B. Ice: O periodičnosti kristalizacije kovin . . . . . . ŽZB 21 (1987)3, 127-130 Čurčlja Dušan: Vpliv hitrosti valjanja na proces hladnega valja- nja z mazivi .......ŽZB 21 (1987) 3, 131 —136 Brudar Božidar: Strjevanje jekla v kokili . . . ... ŽZB 21 (1987) 4, 137-150 Kmetič Dimitrij, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M. Jakupovič, B. Ralič: Bele kromove litine legirane z molibdenom za valje .......ŽZB 21 (1987) 4, 151 — 165 Rodlč Tomaž, D. R. J. Owen: Osnovni koncept numerične simulacije radialnega kovanja . . ŽZB 21 (1987) 4, 167—174 Kosec Ladislav, F. Kosel: Nastanek in rast utrujenostne razpoke v korozijskem mediju . . . ŽZB 21 (1987) 4, 175—182 UleBoris, F. Vodopivec, J. Žvokelj, M. Grašič.L. Kosec:Zapozneli lom jekla z visoko trdnostjo . . . ŽZB 21 (1987) 4, 183—192 2. AVTORSKO KAZALO Arh Jože, V. Prešeren: Dosežki Železarne Jsenice na področju sekundarne obdelave jekla . . . ŽZB 21 (1987) 1, 7—17 Arh Jože: Tehnična novica: Druga evropska konferenca o elek- tro jeklarstvu........ŽZB 21 (1987) 1, 59—63 Ažman Alojz, D. Sikošek, A. Šteblaj, J. Triplat, J. Arh: Tehnična novica: Novi konstrukcijski mikrolegirani jekli Niomol 390 in Niomol 490 ..........ŽZB 21 (1987) 1,45-52 Brudar Božidar: Ohlajanje jeklenega valja na vozičku . . . . . . ZZB 21 (1987)2,85—92 Brudar Božidar: Strjevanje jekla v kokili . . . ... ŽZB 21 (1987)4, 137-150 Curčija Dušan: Vpliv hitrosti valjanja na proces hladnega valja- nja z mazivi .......ŽZB 21 (1987) 3, 131-136 Gnamuš Janko, G. Rihtar: Reperaturno varjenje orodnih jekel ŽZB 21 (1987) 2, 73-76 Grešovnik Ferdo: Računanje temperaturnih napetosti v elastičnem področju .......ŽZB 21 (1987) 2, 77—83 Kaker Henrik: Mikroanaliza faz v zlitini Nimonic 80 A s kombinacijo REM-EDS .......ŽZB 21 (1987) 1, 39-43 Kmetič Dimitrij, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M. Jakupovič, B. Ralič: Bele kromove litine legirane z molibdenom za valje .......ŽZB 21 (1987) 4, 151 — 165 Kosec Ladislav, F. Kosel: Nastanek in rast utrujenostne razpoke v korozijskem mediju . . . ŽZB 21 (1987) 4, 175—182 Leskovšek Vojteh: Tehnična novica: Predstavitev enokomorne vakuumske peči IPSEN VTC 324-R s homogenim plinskim hlajenjem pod visokim tlakom .... ŽZB 21 (1987) 1, 53—57 Paulin Andrej, J. Lamut, D. Dretnik: Reaktivnost koksa in njen vpliv na delo plavža......ŽZB 21 (1987) 2, 65—71 Paulin Andrej: Mehanizem zgorevanja koksa . . . . . . ŽZB 21 (1987) 3, 105-112 Risteski B. Ice: O periodičnosti kristalizacije kovin . . . . . . ŽZB 21 (1987)3, 127-130 Rodič Tomaž, D. R.J. Owen: Osnovni koncept numerične simulacije radialnega kovanja . . ŽZB 21 (1987) 4, 167—174 Todorovič Gojko, J. Lamut, B. Dobovišek, J. Kramer, J. Zapu-šek, B. Sedlar: Naogljičenje železa med redukcijo in taljenjem plavžnega vsipa........ŽZB 21 (1987) 4,1 -5 Ule Boris, F. Vodopivec, J. Žvokelj, M. Grašič.L. Kosec:Zapozneli lom jeklaz visoko trdnostjo . . . ŽZB 21 (1987) 4, 183—192 Uranc Franc: Vpliv trenja, poti, drsne hitrosti in pritiska na obrabo.........ŽZB 21 (1987) 3, 119-125 Vodopivec Franc, M. Gabrovšek: Meja plastičnosti konstrukcijskih jekel, fizikalno metalurške osnove . . ŽZB 21 (1987) 1,19—28 Vodopivec Franc, F. Marinšek, F. Grešovnik: Rekristalizacija in rast zrn pri žarjenju hladno valjanega jekla z 0,03 % C, 1,8 % Si, 0,3 % Mn in 0,3 % Al......ŽZB 21 (1987) 1, 29—37 Vodopivec Franc, O. Kurner, A. Lagoja, F. Grešovnik, A. Rodič, S. Senčič.F. Vizjak:Tehničnanovica:Orazvojnih možnostih jekel in nekaterih posebnih zlitin ter postopkov za njihovo izdelavo, predelavo, ulivanje in plemenitenje .... ŽZB 21 (1987) 2, 93—98 Vodopivec Franc, F. Grešovnik, F. Marinšek, M. Kmetič, O. Kurner: O teksturi valjanja, razogljičenja in rekristalizacije v jeklu z 0,03 C, 1,8 Si in 0,3 Al . . ŽZB 21 (1987) 3, 113-118 3. KAZALO PO STROKAH — UDK 53 FIZIKA 531 Splošna mehanika Grešovnik Ferdo: Računanje temperaturnih napetosti v elastičnem področju .......ŽZB 21 (1987)2,77-83 536 Nauk o toploti. Termodinamika Brudar Božidar: Ohlajenje jeklenega valja na vozičku ŽZB 21 (1987) 2, 85-92 62 — INŽENIRSTVO, TEHNIKA 620.17 Preskušanje mehanskih lastnosti Uranc Franc: Vpliv trenja, poti, drsne hitrosti in pritiska na obrabo.........ŽZB 21 (1987) 3, 119-125 620.18 Metalografija Kaker Henrik: Mikroanaliza faz v zlitini Nimonic 80 A s kombinacijo REM-EDS .......ŽZB 21 (1987) 1, 39-43 620.19 Napake v materialu Kosec Ladislav, F. Kosel: Nastanek in rast utrujenostne razpoke v korozijskem mediju . . . ŽZB 21 (1987) 4, 175-182 621.73 Kovanje Rodič Tomaž, D.R.J. Ovven: Osnovni koncept numerične simulacije radialnega kovanja . . . ŽZB 21 (1987) 4, 167-174 621.74.047 Kontinuirno litje Risteski B. Ice: O periodičnosti kristalizacije kovin . . . ŽZB 21 (1987)3, 127-130 621.77 Valjanje, stiskanje, vlečenje Čurčiia Dušan: Vpliv hitrosti valjanja na proces hladnega valjanja z mazivi .......ŽZB 21 (1987) 3, 131 -136 Vodopivec Franc, F. Grešovnik, F. Marinšek, M. Kmetič, O. Kur-ner- O teksturi valjanja, razogljičenja in rekristalizacije v jeklu z 0,03 C, 1,8 Si in 0,3 Al .... ŽZB 21 (1987) 3, 113-118 621.78 Toplotna obdelava kovin Leskovšek Vojteh: Tehnična novica: Predstavitev enokomorne vakuumske peči IPSEN VTC 324-R s homogenim plinskim hlajenjem pod visokim tlakom .... ŽZB 21 (1987) 1,53-57 621.791 Varjenje in podobni postopki Gnamuš Janko, G. Rihtar: Reperaturno varjenje orodnih jekel . . . ŽZB 21 (1987)2,73-76 66 KEMIJSKA TEHNIKA, KEMIČNE IN SORODNE INDUSTRIJE 669 — Metalurgija Vodopivec Franc, O. Kurner, A. Lagoja, F. Grešovnik, A. Rodič, S. Senčič, F. Vizjak: Tehnična novica: O razvojnih možnostih jekel in nekaterih posebnih zlitin ter postopkov za njihovo izdelavo, predelavo, ulivanje in plemenitenje ŽZB 21 (1987) 2, 93—98 669.01 Splošna in teoretična metalurgija Paulin Andrej: Mehanizem zgorevanja koksa . . . . . . ŽZB 21 (1987)3, 105-112 Vodopivec Franc, F. Marinšek, F. Grešovnik: Rekristalizacija in rast zrn pri žarjenju hladno valjanega jekla z 0,03 % C, 1,8 % Si, 0,3 % Mn in 0,3 % Al......ŽZB 21 (1987) 1, 29-37 669.14 Zlitine železa z ogljikom. Jeklo nasploh Ule Boris F. Vodopivec, J. Žvokelj, M. Grašič: Zapozneli lom jekla z visoko trdnostjo .... ŽZB 21 (1987)4, 183-193 Vodopivec Franc, M. Gabrovšek: Meja plastičnosti konstrukcijskih jekel, fizikalno metalurške osnove . . ŽZB 21 (1987) 1,19—25 Ažman Alojz, D. Sikošek, A. Šteblaj, J. Triplat, J. Arh: Tehnična novica: Novi konstrukcijski mikrolegirani jekli Niomol 390 in Nio-mol 490 ..........ŽZB 21 (1987) 1,45-52 669.15 Zlitine železa z drugimi elementi, razen ogljikom. Legi-rana jekla. Ferozlitine Kmetič Dimitrij, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M Jakupovič, B. Ralič: Bele kromove litine legirane z molibdenom za valje .......ŽZB 21 (1987)4, 151-165 669.16 Proizvodnja grodlja Todorovič Gojko, J. Lamut, B. Dobovišek, J. Kramer, J. Zapu-šek, B. Sedlar: Naogljičenje železa med redukcijo in taljenjem plavžnega vsipa........ŽZB 21 (1987) 1,1-5 669.18 Proizvodnja jekla Brudar Božidar: Strjevanje jekla v kokili . . . ... ŽZB 21 (1987) 4, 137-150 Arh Jože, V. Prešeren: Dosežki Železarne Jesenice na področju sekundarne obdelave jekla . . . ŽZB 21 (1987) 1, 7—17 Arh Jože: Tehnična novica: Druga svetovna konferenca o elektro jeklarstvu ........ŽZB 21 (1987) 1,59-63 669.4 Svinec. Svinčeve zlitine Paulin Andrej, J. Lamut, D. Dretnik: Reaktivnost koksa in njen vpliv na delo plavža......ŽZB 21 (1987) 2, 65-71 VSEBINA UDK: 669.18:669. i 12.223:620.192.43 ASM/SLA: N21, M28h. E25n, D9p, 9-69 Metalurgija — jeklarstvo — strjevanje jekla — izceje — sekundarni lunker B Brudar Strjevanje jekla v kokili Železarski zbornik 21 (1987) 4 s 137—150 Na osnovi simulacij ohlajanja bloka v kokili po poenostavljenem modelu smo prišli do nekaterih novih ugotovitev glede poteka strjevanja. Predvsem smo želeli povezati potek izračunanih izoterm s potekom izcej. Da bi se prepričali o pravilnosti svojih predpostavk. smo rekonstruirali kokilo OK 650. Vlili smo 4 poskusne bloke avtomatnega jekla Č 3990, jih prerezali po dolgem, naredili Baumannov odtis in fotografirali jedkano ploskev. Iz primerjave med posameznimi slikami sklepamo, da je mogoče s primerno stanjšano steno kokile vplivati tudi na porazdelitev izcej V v glavi bloka. Pričakujemo, da bi na ta način lahko odpravili probleme, ki nastopajo pri valjanju nekaterih kvalitet, kadar je za to vzrok nehomogenost v sredi bloka. Avtorski izvleček UDK: 620.193.01 ASM/SLA: Rlh, Rle, R2j, Q26p Metalurgija — Temperaturna utrujenost — Korozija L. Kosec, F. Kosel Nastanek in rast utrujenostne razpoke v korozijskem mediju Železarski zbornik 21 (1987) 4 s 175-182 V korozijskeh medijih se površine kovin prekrijejo s korozijskimi produkti. To je posebej izrazito pri oksidaciji kovin na povišanih ali visokih temperaturah. Mehanska nestabilnost teh plasti je vzrok, da skozi nastale razpoke prihaja korozijski medij v stik s kovino, zaradi česar rastejo korozijski produkti v obliki klinov. Menjajoče ali stalne natezne napetosti povzroče trajno nestabilnost korozijskih produktov in pospešeno napredovanje klinov v globino kovine. V sistemih s korozijskimi klini nastanejo značilna napetost-no-deformacijska stanja, ki vplivajo na oblikovanje sistema. Motnje, ki preprečujejo hiter dotok korozijskega medija do kovine na vrhu klinov, zavirajo njihovo rast. Posebej uspešne v tem so kompozitne plasti (kovina-oksid), ki so mehansko zelo stabilne. Skupine oksidnih klinov rastejo v enakih pogojih počasneje kol če je v sistemu en sam klin enake velikosti. V kemično nehomogenih materialih (s kristalnimi izcejami) so praviloma negativne izceje mesto, v katerem nastanejo in rastejo oksidni klini. Avtorski izvleček UDK: 669.15"26-194:669.14.018.255 ASM/SLA: M28, N8b, TSk, 5, Cr, W23k Metalurgija — bele kromove litine — mikrostruktura — 111 in CTT diagrami D. Kmetič, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec, M. Jakupovič, B. Ralič Bele kromove litine legirane z molibdenom za valje Železarski zbornik 21 (1987) 4 s 151-165 Delo obravnava mikrostrukturne značilnosti belih kromovih litin legiranih z molibdenom v litem in toplotno obdelanem stanju. Na izoblikovanje eutektičnih karbidov, delež karbidne faze in mikrostrukturo matice v litem stanju vplivajo poleg pogojev strjevanja in ohlajanja, razmerje Cr/C in vsebnost ogljika in molibdena v zlitini. Podani so izotermni transformacijski diagrami za destabilizacijo austenita in za destabiliziran austenit in kontinuirni transformacijski diagram za destabiliziran austenit. Avtorski izvleček UDK: 669.14.018.2:539.56:620.192.3 ASM/SLA: Q26s, SGBa, ST, 2-60, EGn, 3-66 Metalurgija — fizika kovin B. Ule, F. Vodopivec, J. Žvokelj, M. Grašič in L. Kosec Zapozneli lom jekla z visoko trdnostjo Železarski zbornik 21 (1987) 4 s 183-192 Prispevek obravnava teoretično analizo napetostno inducirane-ga segregiranja vodika, ki pri jeklu z visoko trdnostjo povzroči zapozneli lom. Na osnovi merjenja kritičnega in mejnega napetostnega intenzitetnega faktorja ter ob pomoči mikrofraktografskih preiskav je bilo ugotovljeno, da ima popuščena martenzitna mikrostruktura enak Kth kot popuščena martenzitno-bainitna mikrostruktura, enake meje plastičnosti, da pa ima slednja boljšo lomno žilavost. Avtorski izvleček UDK: 621.73.045:519.6 ASM/SLA: F22, Q24, 1-66, U4g, U4k Metalurgija, kovanje, matematični model T. Rodič, D. R. J. Ovven: Osnovni koncept numerične simulacije radialnega kovanja Železarski zbornik 21 (1987) 4 s 167-174 MKEje pogosto uporabljena metoda za analizo preoblikovalnih procesov. V primerjavi s klasičnimi metodami ima naslednje prednosti: z MKE lahko rešujemo primere z zahtevno geometrijo preo-blikovanca; problem lahko rešujemo z različnimi materialnimi modeli; možno je obravnavanje nestacionarnih napetostno-deforma-cijskih in temperaturnih polj. V prihodnje pričakujemo vgrajevanje novih spoznanj s področja fizikalne metalurgije v numerične modele. Prvi koraki v tej smeri so bili že storjeni."-18 Avtorski izvleček CONTENTS UDK: 620.193.01 ASM/SLA: Rlh, Rle, R2j, Q26p Metallurgy — Thermal Fatigue — Corrosion L. Kosec, F. Kosel Occurrence and Crowth of Fatigue Cracks in Corrosion Environment Železarski zbornik 21 (1987) 4 P 175-182 In corrosion environment metal surfaces become covered by corrosion products. This is a typical feature of metal oxidation at temperature rise and high temperatures. The mechanical instability of these layers is the reason for the occurrence of cracks through which the oxidant gets in contact vvith the metal, making the corrosion products grovv in the form of vvedges. Changing or constant tensile stresses cause a permanent instability of corrosion products and intensify the grovvth of vvedges deep into the metal. In corrosion vvedge systems typical stress-strain states occur, affecting their morphology. Barriers preventing a rapid access of the oxidant to the metal at the vvedge tip, retard the grovvth. Especially successfrul barriers are composite metal/oxide layers vvhich are mechanically very stable. In equal conditions groups of oxide vvedges grovv more slowly than if there is only one vvedge of the same size. In chemical-ly nonhomogeneous materials vvith crystal segregations oxide vvedges vvould as a rule start grovving in the areas of negative segregations. Authors Abstract UDK: 669.18:669.112.223:620.192.43 ASM/SLA: N21, M28h, E25n, D9p, 9-69 Metallurgy — steelmaking — solidification of steel — segregations — secondary pipe B. Brudar Solidification of Steel in a Mould Železarski zbornik 21 (1987) 4 P 137-150 On the basis of simulations of solidification of an ingot in the mould using the simplified model nevv facts about the process of solidification vvere obtained. We vvished to correlate the calculated is-otherms vvith the course of segregations. In order to check the val-idity of our assumptions we reconstructed the mould OK 650. There vvere 4 test ingots of the free cutting steel Č 3990 čast. They-vvere cul longitudinally and the sulphur prints and the photos of the etched cross-sections vvere made. From the comparison among different figures it could be concluded that it vvas possible to influence the distribution of segregates by a proper thinning of the mould vvall. This holds especially for the V-segregates in the upper part of the ingot. We expect that in this way some problems occuring vvith the rolling of some qualities could be suppressed vvhen it is sup-posed that the unhomogeneities in the middle of the ingot are the reason for them. Author's Abstract UDK: 669.14.018.2:539.56:620.192.3 ASM/SLA: Q26s, SGBa, ST, 2-60, EGn, 3-66 Metallurgy — materials Science B. Ule, F. Vodopivec, J. Žvokelj, M. Grašič in L. Kosec Delayed Fracture of High-strength steel Železarski zbornik 21 (1987) 4 P 183-192 The paper presents theoretical analysis of stress induced hydrog-en segregation, vvhich produces delayed fracture of high-strength steels. On the basis of measurements of the critical and the threshold stress intensity factor and by means of microfractographic ex-aminations, it vvas established that the tempered martensitic microstructure has the same KTH as the tempered martensitic-bainitic microstructure vvith the same yield strenth, the latter having a bet-ter fracture toughness. Author's Abstract UDK: 669.15'26-194:669.14.018.255 ASM/SLA: M28. N8b, TSk, 5, Cr. W23k Metallurgy — White Chromium Čast Irons — Microstructure — TTT and CTT Diagrams D. Kmetič, F. Mlakar, V. Tucič, J. Žvokelj, F. Vodopivec. M. Jaku-povič, B. Ralič White Chromium Čast Irons for Rolls, Alloyed vvith Molybdenum Železarski zbornik 21 (1987) 4 P 151-165 The paper treats the microstructural characteristics of the vvhite chromium čast irons aloyed vvith molybdenum, as čast and as heat treated. Formation of eutectic carbides, portion of carbide phase. and microstructure of matrix in the čast state are influenced by the Cr/C ratio, and carbon and molybdenum contents in the alloy be-side the conditions of solidification and cooling. Isothermal transformation diagrams for destabilization of austenite, and for destabilized austenite, next to the continuous transformation diagram for destabilized austenite are presented. Author's Abstract UDK: 621.73.045:519.6 ASM/SLA: F22. Q24. 1-66, U4g, U4k Metallurgy, forging. mathematical model T. Rodič, D. R. J. Ovven: Basic Concepts of Numerical Simulation of a Radial Forging Process Železarski zbornik 21 (1987) 4 P 167—174 The FEM is novv widely used for the analysis of metal forming processes. In comparison vvith classical methods the FEM has cer-tain advantages: various complex shapes can be considered, it en-ables implementation of diferent material models and treatment of transient stress-strain and temperature fields. In the future the in-clusion of the metallurgical development in the numerical modell-ing of hot working processes is expected. The first steps tovvard this goal have already been made17 Authors Abstract