KOVINE ZLITINE TEHNOLOGIJE METALS ALLOYS TECHNOLOGIES i m 0) oi I O ■ h- Ui ■ -J SELECTED PAPERS PRESENTED AT THE SLOVENIAN - HUNGARIAN - CROATIAN - AUSTRIAN 6 JOINT VACUUM CONFERENCE AND 3 MEETING OF THE SLOVENIAN AND CROATIAN VACUUMOLOGISTS BLED 4-7 APRIL 1995 GUEST EDITORS: F. VODOPIVEC AND M. JENKO PART II KOVINE LETNIK STEV. STR. LJUBLJANA JUL.-SEPT. ZLITINE 29 3-4 361-432 TEHNOLOGIJE VOLUME NO. P. SLOVENIJA 1995 Prosimo avtorje, da pri pripravi rokopisa za objavo članka dosled-noupoštevajo naslednja navodila: - Članek mora biti izvirno delo, ki ni bilo v dani obliki še nikjer objavljeno. Deli članka so lahko že bili podani kot referat. - Avtor naj odda članek oz. besedilo napisano na računalnik z urejevalniki besedil: - VVORDSTAR, verzija 4, 5, 6, 7 za DOS - VVORD za DOS ali WINDOWS. Če avtor besedila ne more dostaviti v prej naštetih oblikah, naj pošlje besedilo urejeno v ASCII formatu. Prosimo avtorje, da pošljejo disketo z oznako datoteke in računalniškim izpisom te datoteke na papirju. Formule so lahko v datoteki samo naznačene. na izpisu pa ročno izpisane. Celoten rokopis članka obsega: - naslov članka (v slovenskem in angleškem jeziku). - podatke o avtorju. - povzetek (v slovenskem in angleškem jeziku). - ključne besede (v slovenskem in angleškem jeziku), - besedilo članka, - preglednice, tabele. - slike (risbe ali fotografije), - podpise k slikam (v slovenskem in angleškem jeziku), - pregled literature. Članek naj bi bil čim krajši in naj ne bi presegal 5-7 tiskanih strani, pregledni članek 12 strani, prispevek s posvetovanj pa 3-5 tiskanih strani. Obvezna je raba merskih enot, ki jih določa zakon o merskih enotah in merilih, tj. enot mednarodnega sistema SI. Enačbe se označujejo ob desni strani besedila s tekočo številko v okroglih oklepajih. Preglednice (tabele) je treba napisati na posebnih listih in ne med besedilom. V preglednicah naj se - kjer je le mogoče - ne uporabljajo izpisana imena veličin, ampak ustrezni simboli. Slike (risbe ali fotografije) morajo biti priložene posebej in ne vstavljene (ali nalepljene) med besedilom. Risbe naj bodo izdelane praviloma povečane v merilu 2:1. Za vse slike po fotografskih posnetkih je potrebno priložiti izvirne fotografije, ki so ostre, kontrastne in primerno velike. Vsi podpisi k slikam (v slovenskem in angleškem jeziku) naj bodo zbrani na posebnem listu in ne med besedilom. V pregledu literature naj bo vsak vir oštevilčen s tekočo številko v oglatih oklepajih (ki jih uporabljamo tudi med besedilom, kadar se želimo sklicevati na določeni literarni vir). Vsak vir mora biti opremljen s podatki, ki omogočajo bralcu, da ga lahko poišče: kn|ige: - avtor, naslov knjige, ime založbe in kraj ter leto izdaje (po potrebi tudi določene strani): H. Ibach and H. Luth, Solid State Physics, Springer, Berlin 1991. p. 245 članki: - avtor, naslov članka, ime revije in kraj izhajanja, letnik, leto, številka ter strani: H. J. Grabke, Kovine zlitine tehnologije. 27, (1993) 1-2, 9 Avtorji naj rokopisu članka priložijo povzetek v omejenem obsegu do 10 vrstic v slovenskem in angleškem jeziku. Rokopisu morajo biti dodani tudi podatki o avtorju: - ime in priimek, akademski naslov in poklic, ime delovne organizacije v kateri dela. naslov stanovanja, telefonska številka. E-mail in številka fax-a. Uredništvo KZT - odloča o sprejemu članka za objavo, - poskrbi za strokovne ocene in morebitne predloge za krajšanje ali izpopolnitev, - poskrbi za jezikovne korekture. Authors are kindly requested to prepare the manuscripts according to the follovving instructions: - The paper must be original, unpublished and properly prepared for printing. - Manuscripts should be typed with double spacing and wide margins on numbered pages and should be submitted on flop-py disk in form of: - VVORDSTAR. version 4. 5. 6. 7 for DOS. - VVORD for DOS or VVINDOVVS. - ASCII text vvithout formulae, in vvhich čase formulae should be clearly written by hand in the printed copy. Preparation of Manuscript: - the paper title (in English and Slovenian Language)* - author(s) name(s) and affiliation(s) - the text of the Abstract (in English and Slovenian Language)' - key vvords (in English and Slovenian Language)' - the text of the paper (in English and Slovenian Language)* - tables (in English Language) - figures (dravvings or photographs) - captions to figures (in English and Slovenian Language)" - captions to tables (in English) - acknovvledgement - references * The Editorial Board will provide for the translation in Slovenian Language for foreign authors. The length of published papers should not exceed 5-7 journal pages. of revievv papers 12 journal pages and of contributed papers 3-5 journal pages. The international system units (SI) should be used. Equations should be numbered sequentially on the right-hand side in round brackets. Tables should be typed on separate sheets at the end of manuscript. They should have a descriptive caption explaining dis-played data. Figures (dravvings or photographs) should be numbered and their captions listed together at the end of the manuscript. The dravvings for the line figures should be tvvice the size than in the print Figures have to be original, sharp and well contrasted. enclosed separately to the text. References must be typed in a separate reference section at the end of the manuscript. vvith items refereed too in the text by numerals in square brackets. References must be presented as follovvs: - books: author(s). title. the publisher, location. year. page num-bers H. Ibach and H. Luth. Solid State Physics. Springer. Berlin 1991, p. 245 - articles: author(s). a journal name. volume. a year. issue num-ber, page H. J. Grabke. Kovine zlitine thenologije, 27. (1993). 1-2. 9 The abstract (both in English and in Slovenian Language) should not exceed 200 vvords. The title page should contain each author(s) full names. affiliation vvith full address, E-mail number. telephone and fax number if available. The Editor - will decide if the paper is accepted for publication. - will take care of th.e refereeing process. - language corrections. The manuscripts of papers accepted for publication are not re-turned. Rokopisi člankov ostanejo v arhivu uredništva Kovine zlitine tehnologije. KOVINE ZLITINE TEHNOLOGIJE tt 2 2 9 2 8 0 METALS ALLOYS TECHNOLOGIES KOVINE ZLITINE TEHNOLOGIJE Izdajatelj (Published for): Inštitut za kovinske materiale in tehnologije Ljubljana Soizdajatelji (Associated Publishers): SŽ ŽJ ACRONI Jesenice. IMPOL Slovenska Bistrica, Kemijski inštitut Ljubljana. Koncem Slovenske Železarne. Metal Ravne. Talum Kidričevo Izdajanje KOVINE ZLITINE TEHNOLOGIJE sofinancira: Ministrstvo za znanost in tehnologijo Republike Slovenije (Journal METALS ALLOYS TECHNOLOGIES is financially supported by Ministrstvo za znanost in tehnologijo Republike Slovenije) Glavni in odgovorni urednik (Editor-in-chief): prof. Franc Vodopivec, Inštitut za kovinske materiale in tehnologije Ljubljana. 61000 Ljubljana, Lepi pot 11, Slovenija Urednik (Editor): mag. Aleš Lagoja Tehnični urednik (Technical Editor): Jana Jamar Lektorji (Linguistic Advisers): dr. Jože Gasperič in Jana Jamar (slovenski jezik), prof. dr. Andrej Paulin (angleški jezik) Uredniški odbor (Editorial Board): doc. dr. Monika Jenko. prof. Jakob Lamut, prof. Vasilij Prešeren, prof. Jože Vižintin, prof. Stane Pejovnik, dipl. ing. Sudradjat Dai. Jana Jamar Mednarodni pridruženi člani uredniškega odbora (International Advisory Board): prof. Hans Jurgen Grabke. Max-Planck-lnstitut tur Eisenforschung, Dusseldorf, Deutschland prof. Thomas Bell. Faculty of Engineering School of Metallurgy and Materials. The University of Birmingham, Birmingham. UK prof. Jožef Zrnik. Technicka Univerzita. Hutnicka fakulteta. Košice. Slovakia prof. Ilija Mamuzic, Sveučilište u Zagrebu, Hrvatska prof. V. Lupine. Istituto per la Tecnologia dei Materiali Metallici non Tradizionali. Milano. Italia prof. Gunther Petzov. Max-Planck-lnstitut fur Metallforschung. Stuttgart, Deutschland prof. Hans-Eckart Oechsner. Universitat Darmstadt. Deutschland Izdajateljski svet (Editorial Advisory Board): prof. Marin Gabrovšek. prof. Blaženko Koroušič, prof. Ladislav Kosec. prof. Alojz Križman, prof. Tatjana Malavašič. dr. Tomaž Kosmač, prof. Leopold Vehovar. prof. Anton Smolej. dr. Boris Ule, doc. dr. Tomaž Kolenko. dr. Jelena Vojvodič-Gvardjančič Naslov uredništva (Editorial Address): KOVINE ZLITINE TEHNOLOGIJE IMT Ljubljana Lepi pot 11 61000 Ljubljana. Slovenija Telefon:+386 61 125 11 61 Telefax: +386 61 213 780 Žiro račun: 50101-603-50316 IMT pri Agencija Ljubljana Cover: Experiment. first performed in Magdeburg in 1657. von Guericke demonstrated the effect of air pressure on an evacuated cavity formed by two tightly fitted copper hemispheres. Naslovnica: Magdeburški poskus z dvema polkroglama in evakuiranim vmesnim prostorom, ki ju ni moglo ločiti osem parov konj. Poskus je javno prikazal Otto von Guericke leta 1657. Oblikovanje ovitka: Ignac Kofol Stavek: Majda Kuraš Tisk: Gorenjski tisk. Kranj Po mnenju Ministrstva za znanost in tehnologijo Republike Slovenije št. 23-335-92 z dne 09. 06.1992 šteje KOVINE ZLITINE TEHNOLOGIJE med proizvode, za katere se plačuje 5-odstotn davek od prometa proizvodov Preface The 1995 Bled Joint Vacuum Conference continues the tradition of small family type meetings of Austrian, Hungarian, Croatian and Slovenian scientists in the field of vacuum physics, techniques and related topics. Former Austrian-Hungarian Joint Conferences were held in Gyor (1979), Brunn am Gebirge (1981) and in Debrecen (1985) vvhere for the first tirne Yugoslavia was also included, than Portorož (1988) and Vienna (1991). At the Fourth European Vacuum Conference (EVC-4. Uppsala 1994), the delegates from Austria, Croatia, Hungary and Slovenia expressed the vvillingness to continue regional cooperation and agreed to organize Slovenian-Hungarian-Croatian-Austrian Sixth Joint Vacuum Conference, vvith the Slovenian Vacuum Society and Institute of Metals and Technologies, Ljubljana as organizers. Scientists from Czech Republic, Slovak Republic and Poland joined the Conference. An impressive number of excellent contributions were presented on the Conference, instead of its small family type. Papers related to applied surface science, electronic materials, plasma science, surface science, thin films and vacuum science were already published in Journal FIZIKA, 1995, Number 2. Papers related to vacuum metallurgy are published in the present number of the Journal KOVINE ZLITINE TEHNOLOGIJE - METALS ALLOYS TECHNOLOGIES. Various subjects, such as electron beam vvelding, laser vvelding, vacuum brazing, surface treatment by plasma nitriding hard coatings as well as industrial experiences by vacuum manufacturing of stainless steels are included. The authors are affiliated mostly to academic research institutions. In spite of this, results of applied research as well as technological experience from various countries are presented in the papers, vvhich should meet the interest of a large audience. Editors: F. Vodopivec and M. Jenko Nagrajeni mladi raziskovalci za najboljši raziskovalni dosežek predstavljen na 45. Posvetovanju o metalurgiji in kovinskih gradivih in 2. Posvetovanju o materialih na 14. Slovenskem vakuumskem posvetovanju, Portorož, 3.-5. oktober 1994 Darja Steiner Petrovič, dipl. ing. metalurgije, se je na 45. Posvetovanju o metalurgiji in kovinskih gradivih, 2. Posvetovanju o materialih in 14. Slovenskem vakuumskem posvetovanju, predstavila z delom: Razogljičenje in rekristalizacija neorientirane elektro pločevine. Darja Steiner Petrovič se je rodila 18. oktobra 1969 v Ljubljani. Srednjo naravoslovno-matematično šolo je zaključila leta 1988 in nato vpisala študij metalurgije na Univerzi v Ljubljani, FNT -Montanistika, kjer je leta 1993 tudi diplomirala. Kot mlada raziskovalka je pod mentorstvom doc. dr. Monike Jenko zaposlena na Inštitutu za kovinske materiale in tehnologije v Ljubljani. V času podiplomskega študija se je izpopolnjevala tudi na Inštitutu za fiziko Zagreb, v laboratoriju za fiziko trdne snovi. Contents Original Scientific Papers Vacuum Metallurgy Future Aspects of Electron Beam VVelding and Surface Modification Friedel K. P.. W. Sielanko...................................................... 364-376 Operational Aspects of Experiences in Vacuum Technology by Production of High Quality Stainless and Alloyed Steels Koroušič B., A. Rozman, J. Triplat, J. Lamut ....................................... 377-384 Electron Beam VVelding of Chromium-Nickel Stainless Steel to Duralumin Grodzinski A., J. Senkara, M. Kozlovvski........................................... 385-390 Modification of Steel Surface vvith Nickel Alloy by an Electron Beam Kozlovvski M., J. Senkara ...................................................... 391-395 Influence of Fracture Toughness on Vacuum Hardened HSS Leskovšek l/., B. Ule, A. Rodič .................................................. 397-404 Characteristics of Cemented Carbide Particles/Structural Steel Vacuum Brazing Joint Šuštaršič B., I/. Leskovšek, A. Rodič ............................................. 405-412 Pulsed Plasma Nitriding of Stainless Steel Torkar M., V. Leskovšek, B. Rjazancev ...........................................413-416 Hydrogen and Temper Embrittlement of Medium Strength Steel Ule B.. V. Leskovšek .........................................................417-422 A New Concept of Quality Evaluation of High Energy Electron Beam Used in VVelding VVojcicki S.................................................................. 423-426 Laser Induced Reaction betvveen Fe Layer and CuNi30Mn1Fe Alloy Spruk S., B. Praček, A. Zalar, A. Rodič ........................................... 427-430 Applications of a New Electron Beam VVelding Technology in Vacuum Equipment Design Dupak J., P. Kapounek, M. Horaček.............................................. 431-432 INSTITUTE OF METALS AND TECHNOLOGIES p.o. 61000 LJUBLJANA. LEPI POT 11. POB 431 SLOVENIJA Telefon: 061/1251-161, Telefax: 061 213-780 VACUUM HEAT TREATMENT LABORATORY Vacuum Brazing Universally accepted as the most versatile method of joining metals. Vacuum Brazing is a precision metal joining technique suitable for many component configurations in a wide range of materials. ADVANTAGES • Flux free process yields clean, high integrity joints • Reproducible quality • Components of dissimilar geometrv or material type may be joined • Uniform heating & cooling rates minimise distortion • Fluxless brazing alloys ensure strong defect free joints • Bright surface that dispense with expensive post cleaning operations • Cost effective Over five years of Vacuum Brazing expertise at IMT has created an unrivalled reputation for excellence and quality. Our experience in value engineering will often lead to the use of Vacuum Brazing as a cost effective solution to modern technical problems in joining. INDUSTRIES • Aerospace • Mechanical • Electronics • Hydraulics • Pneumatics • Marine • Nuclear • Automotive QUALITY ASSURANCE Quality is fundamental to the IMT philosophy. The choice of process, ali processing operations and process control are continuously monitored by IMT Quality Control Department. The high level of quality resulting from this tightly organised activity is recognised by government authorities, industry and International companies. Future Aspects of Electron Beam VVelding and Surface Modification Perspektive varjenja in obdelave površin z elektronskim curkom Friedel K. P.,1 Institute of Vacuum Technology, Warsaw, and Institute of Electronic Technology, The VVroclavv Technical University W. Sielanko, Institute of Vacuum Technology, Warsaw Some aspects of future electron beam technology are given, especially for vvelding and modification of metal surfaces. One of the prime advantages of EB vvelding is the ability to make vveld deeper and narrovver than one can make using other methods of vvelding. The vveld bead may contain some defects typical for EB vvelding. The essential requirement for the future equipment development is to avoid or eliminate these defects. Advances in the equipment from this point of vievv are revievved. Key words: electron beam vvelding. vveld defects, equipment for vvelding Podane so nekatere perspektivne tehnologije varjenja z elektronskim curkom, posebno za varjenje in obdelavo površin kovinskih materialov. Najpomembnejša prednost varjenja z elektronskim curkom je možnost izdelave ožjih in globjih zvarov v primerjavi z drugimi varilnimi tehnikami. Nizi zvarov lahko vsebujejo napake, ki so tipične za postopek varjenja z elektronskim curkom. Pomembna zahteva za razvoj opreme v bodočnosti je odprava tovrstnih napak. Ključne besede: varjenje z elektronskim curkom, napake zvarov, oprema za varjenje 1. Introduction The industrial application of electron beam (EB) vvelding and surface modification has continually in-creased since the early 1960's in the number of ma-chines as vvell as in the number of applications. Throughout this period continuous development of equipment and production engineering has taken plače. It is anticipated that EB vvelding and surface modification will continue to develope extensively in future vvithin the two main domains: - enhancement of joint quality and reliability as vvell as process reproducibility, - thick-section vvelding and surface hardening of large vvorkpieces vvith complex shapes at atmos-pheric pressure. 2. Joint quality and reiiability The main reason for vvhich the EB vvelding vvill be able to satisfy the industrial user needs in the future can be summarised as improvement in joint quality and reliability. The joint quality and reliability de-pends directly upon: - fusion vveld and heat affected zone shape and di-mensions and microstructure, - number and type of defects. Prof. K. P. Friedel,1 Institute of Vacuum Technology, Warsaw, and Institute of Electronic Technology. The VVroclavv Technical University One of the prime advantages of EB vvelding is the ability to make vveld deeper and narrovver than one can make using other methods of vvelding. Unfortunately, the vveld of such a shape may contain some defects. The most of them is also encountered in other methods of vvelding but there are defects typ-ical only for EB vvelding. The macrodefects of vveld caused by EB tend to occur in deep vveld both in its central part (cracks, porosity, unmelted lumps etc.), and in vveld root (porosity, cold shut, spiking). In čase of very high speed of vvelding, humping and under-cut phenomena often occur. 3. Methods of defect prevention The essential requirement is to avoid or eliminate defects by applying of suitable methods. There are two main groups of such methods10: - material selection methods that rely on improvement in chemical composition either "a priori" or by alloying during the vvelding, - system methods vvhich are connected vvith the possibilities offered by the equipment itself. The last group of methods can be divided into further three groups: - methods of defect prevention connected vvith improvement in the original construction of EB vvelding machine (e.g. increasing the equipment resistance against electrical breakdovvn and electromagnetic interference, etc.). These methods are usually out of range for the equipment operator, - parametric methods are crucial in improving the joint quality. Certain parameters of an equipment or vvelding process can be set to make the perfor-mance less sensitive to causes of variation and imperfection, - optional methods that rely on application of special process technology requiring usually the installa-tion of additional optional devices of exactly defined destination. 3.1 Material selection methods The volume of melted metal during EB vvelding is small in comparison vvith other classical vvelding methods. Besides, the rapid heating and self-quenching in region of vvelding lead to highly unbal-anced changes of microstructure. Additionally the high microstructure diversity vvithin the zone of vvelding depends frequently on extremaly different vvelding conditions (e.g. vvelding of steel plate vvith a thick-ness of 5 mm, at vvelding speed of 30 - 50 mm/s, causes the fevv milisecond self-quenching in heat af-fected zone vvithin the range of 800 - 500 C, vvhrereas at thickness of 150 mm and vvelding speed of 1,6 mm/s the rate of self-quenching is equal to ap-proximately 50 s20. The knovvledge of microstructure and particularly of hardness dependence of heat af-fected zone on weld parameters is by no means less important than the knovvledge of bead shape and di-mensions dependence on these parameters. The purity of materials. The basic precondition of high quality joint is the initial purity of materials used. This purity refers not only to surface contamination, but after ali to internal impurities, particularly to the contents of gases as vvell as metallic and nonmetal-lic alloying constituents. Some impurities like sulfur, phosphorus, arsenic, tin and antimony are very un-desirable24. Intergranular cracking arises from the grain-boundary segregation of impurity elements42. The deeper the vveld depth, the higher should be re-quirements concerning material purity (e.g. at vveld depth > 100 mm the contents of oxygen and nitrogen should not exceed 30 and 80 ppm respectively26. Deep vvelding vvith filler materials. Appropriately selected filler alloy could diminish the risk of hot cracks during the vvelding of austenitic steels (some alloying components facilitate the development of ferritic microstructure37. Insertion of nickel foil into the gap may enable to join the nodular iron34. The filler vvire is usually fed directly to the vapour chan-nel (Fig. 1) either ahead of EB, on the front edge of the channel entrance orifice5 23, or behind the EB vvith simultaneous dislocation of vvire melting zone alongside the channel17. Instead of the filler foil the overlay layer may be used24-26. For example, onto the surface of structural steel, contaminated vvith non-metallic elements, the layer of high quality steel can be overlaid (Fig. 2). This way it is possible to ob-tain the intermediate layer of filler material of very good adherence to core material and of very good weldability as vvell. / / / ž EB \ - V.,--- __ .—✓> Figure 1: Deep vvelding vvith filler vvire fed ahead of EB (a) or behind EB. 1 - vvire feeder Figure 2: The overlaying of filler material onto the surface of vvorkpiece. 1 - electron gun, 2 - vvire feeder, 3 - overlayed layer New materials for EB vvelding. There is an urgent need fot not only the entirely nevv materials but also for improved materials destined for EB vvelding. These improvements consist usually in lovvering the concentration of alloying components undesirable from the joint quality standpoint. It vvas found that the decrease in Al content in steel favour an acicular ferrite microstructure in EB vveld and this has enable the development of nevv type of steel for EB vvelding25. The reduction of phosphorus content in steel dimin-ishes the intergranular fracture surface area fraction and improves the toughness42. Nevv steels designed for EB vvelding are usualy low-carbon steels of higher toughness ovving to such alloying components like Cr, Mn, Mo and Ni20 21 44. Besides the properties of many materials are closely related to various manu-facturing operations and particulary the thermal treatment. The proper thermal treatment can improve the properties of EB vveld7. 3.2 Parametric methods It is well known that product or process quality must be properly designed. The Japanese have, through the efforts of Genichi Taguchi14, built quality methods into the engineering process. This method-ology comprises three-phase program involved in the engineering optimization of a product or process: system design, parameter design, tolerance design. The operator of EB vvelding machine is able to un-dertake the optimization procedure only by selecting the appropriate process parameter levels. The parameter design phase is crucial also for the improving to the repeatability of a process. It permits to deter-mine the parameter levels such that the process is optimal and present a minimal sensitivity to causes of variation. The fundamental for parametric method of im-provement in joint quality is the knovvledge of dependence of weld cross-section shape and defec-tiveness degree on vvelding conditions. It is almost impossible to find this dependence theoretically by creating the model of EB-material interactions. Therefore the key stage of such optimization procedure is the properly designed experiment. The most commonly applied design of experiments is the frac-tional factorial experiment. The main limitation of this method is that no interaction among the parameters studied can be observed. To avoid this disadvantage it is necessary to pertorm the full factorial experi-ment. Fig. 3 shovvs the example10 of the relationship be-tvveen weld depth and focusing parameter (i.e. ratio of the vvork distance and the position of maximum povver density plane of EB) obtained experimentall-ly. The "forbiden zones" of enhanced defectiveness are marked by dotted lines. The larger the vvork distance the weld becomes more shallovv, but at the same tirne much more sound. To avoid the porosity of the central part of the vveld, the EB active zone oughut to be located beneath the vvorkpiece surface rather than onto or above this surface. Taguchi method can be also used to optimize the electron beam hardening process. This way it is pos-sible to identify the factors vvhich are the most signif-icant and influencial on thickness and hardness of hardened layer and to found the most suitable pro-cessing conditions vvithin the range of factor levels that vvere chosen. In future many further similar experiments should be performed to establish the allovvable zones of ali process parameters and to prepare the knovvledge representation of both EB vvelding and surface modification processes. Finally, it seems indispensable to create the expert systems based on such knovvledge representation3. 3.3 Optional methods The adoption of special optional methods, vvhich enable to vveld materials of poor weldability allovv to get sound vvelds vvithout commonly encountered de-fects, should significantly extend the area of applica- Focusing Parameter Figure 3: The vveld depth vs focusing parameter and zones of defects appearance. M - magnification, P - EB povver tions of EB vvelding. These optional methods consist mainly in space-time shaping of EB povver density distribution. Horizontal electron beam. Overhelming majority of EB vvelds is made by vertical EB. The depth of EB penetration is restricted not only by EB povver densi-ty but also by disturbances in dynamic equilibrium of vapour channel as well by the highly turbulent flow of molten metal18. Therefor, the thick-section vvelding is usually performed by using horizontal EB (Fig. 4) at horizontal or vertical vvelding direction4'15'19. This method enables to stabilize the molten metal flovv and keep the vapour channel open during vvelding. In order to prevent the excessive outflovv of molten metal the cover metal strap is usually applied along-side the lovver edge of vvelding gap. a) b) Figure 4: VVelding vvith horizontal EB at horizontal (a) or vertical (b) vvelding direction. v - vvelding speed, 1 - vvorkpiece, 2 - electron gun Modulated electron beam. Deeply penetrating, high-power pulsed EB may be utilized to meet the grovving demand for very high quality vvelding of thick-section vvorkpieces. The main advantage of this vvelding method consists in lovvering the average EB povver density and minimizing the melted metal volume. This method of vvelding is recommended for vvelding martensitic steels'516. Astigmatic electron beam. The strongly astigmatic EB can be used to improve the vveld root defective-ness18. This astigmatism is intentionally caused by applying of additional quadrupole lens. Close to the vvorkpiece surface the EB cross-section resemble the elongated ellipse of longer axis parallel to vvelding direction (Fig. 5). In vicinity of the channel bottom the ellipse longer axis is perpendicular to the vvelding direction and prevents the creation of spikes. The quadrupole lens is usually connected in feedback loop and is made active only during špike formation, vvhat can be revealed as sudden drop of backscat-tered electrons current. Inclined electron beam. The permanent inclination of EB axis in vvelding plane improves the flovv of molten metal and facilitate the degassing of resultant vveld bead24. Figure 5: Strongly astigmatic EB. 1 - focusing lens, 2 - quadrupole lens Oscillatlng electron beam. This method consists in very fast deflections of EB during vvelding. Different types of oscillations are used: parallel or perpendicular to the vveld gap, elliptical, circular, "X" type, parabollic and so forth. By selecting the optimum frequency and amplitude, it is possible to obtain a sound vveld vvith a narrovv bandshaped bead6'8'15'23'28-45. Oscillatlon of electron beam active zone. In order to deposite the EB energy more uniformly inside the vapour channel the oscillations of EB active zone alongside the channel axis are used. These oscillations rely on applying the additional AC component to DC component of magnetic lens current815. Doubly-deflected rotating electron beam. The combination of circular oscillations and double de-flection of EB (Fig. 6) enables to control the position of EB active zone independently of focusing lens22. The mutual displacement of EB focus and the vvaist of EB due to double-deflection serves as a tool for the fine adjustment of EB active zone shape and position. Figure 6: Doubly-deflected rotating EB. 1 - quadrupole deflection system, 2 - vvaist of EB due to double deflection, 3 - EB focus Tandem and quasitandem electron beam. Some disadvantages caused by ordinary EB vvelding, such as humping, spiking and root porosity can be over-come by using tandem or quasitandem EB as shown in Fig. 7. This method utilizes two electron beams at the same tirne. One beam is a conventional vertical EB. The other one is used as the auxiliary sub-beam for the repairment of vvelding defects. This auxiliary beam may be directed either on the rear vvall of vapour channel eliminating this way the defects of active zone2 38, ot into the channel eliminating the spiking and root porosity. The auxiliary beam may be also directed onto the molten pool, as shovvn in Fig. 7b, changing the flovv of molten metal so that it flovvs smoothly backvvard240. This way the humping phenomenon may be suppressed. Some modification of this method consist in splitting the EB into tvvo beams using double deflecting unit (Fig. 8). The povver density of both beams can be adjusted by changing the pulse length. Electron beam current decay in the fade-out re-gion. It vvas stated28 that the position of EB active zone has a significant influence on the shape of vveld intensity of root defects in the fade-out region. There are many different techniques of defects suppres-sion in the vveld fade-out region. The best result may be obtained by proper control of the beam current decay rate. A composite beam current decay curve vvas designed vvhich allovved to eliminate the voids and diminished root porosity during fade-out (Fig. 9). b) UP 3 150 I|m\| 100 50 \ \ 4 5 V \ - - N -- 1 1 7 '— —A 0 10 20 30 40 50 t|s| Figure 7: Tandem EB vvelding. 1 - auxiliary EB, 2 - main EB. 3 - active zone porosity, 4 - main vapour channel, 5 - auxiliary channel, 6 - solidification metal layer, v -vvelding speed a) © Figure 8: Ouasitandem EB vvelding. (a) - reheating of the root zone by auxiliary EB, (b) the principle of EB splitting; 1 - auxiliary EB, 2 - main EB, 3 - double-deflection system, v - vvelding speed, H - inclination angle of auxiliary EB Figure 9: The vveld fade-out region. (a) - root defects in the fade-out region, (b) - a composite beam current decay curve; 1 - vveld cross-section, 2 - root porosity, 3 - fade-out region, 4 - optimum EB current decay, 5 - conventional EB current decay 4. Improvements in EB vvelding machine construction One of essential requirements in EB vvelding is to predict and reproduce the vveld geometry particular-ly in čase of frequently encountered partial-penetra-tion vvelds. To satisfy this requirement the continu-ous improvement in performance of EB vvelding machines of moderate povvers and voltages as vvell as of higher povvers, for thick-section vvelding, will be necessary. The overall purpose of equipment developments has been to satisfy the user needs, vvhich can be summarised as improvements in: EB vvelding machine resistance against electrical breakdovvn in region of EB formation and against internal and external electromagnetic interference, process control, user friendliness and possibility of main sub-assemblies monitoring, reliability of povver suppliers and CNC systems. 4.1 Vacuum environment When applied in the heavy industry EB vvelding technology is faced to the problem of the volume of vvorking chamber vvhich should contain the vvork-piece to be vvelded. There are few possibly ways to overcome this problem (Fig. 10): non-vacuum ver-sion of EB vvelding process27 31,3236, local vacuum chamber vvhich only contains the seam of vvork-piece1533, movable chamber vvith local vacuum15, "air-to-air" type of vacuum chamber1. The non-vacuum vvelding has been in industrial use mostly for automotive applications requiring vveld depth of less than 10 mm in a single pass. Recently the non-vacuum beams vvere generated at an accelerating voltage of 270 kV vvith sustained povver levels of in axcess of 110 kW. The penetration depths of more then 50 mm in steel and 40 mm in copper vvere achieved32. Fig. 11 illustrates the rapid degradation of electron beam penetration depth over the range 104 - 103 hPa caused by direct air-scatter-ing. Most EB vvelding processes are currently per-formed in partial vacuum conditions, since the vacuum system is cost-effective in this čase and the vacuum chamber pump-dovvn tirne does not limit significant^ the production throughput30. i pD i , 1 i" \- Figure 10: Vacuum chambers: (a) local, (b) movable, (c) "air-to-air"; 1 - sliding louver, 2 - vvelding gap. 3 - electron gun, 4 - air lock, 5 - continuously vvelded bimetallic ribbon l«1 10- 1(|3 p |hl'a| partial vacuum Figure 11: Effect of air pressure in vacuum chamber on penetration depth. hr - relative vvelded depth, pa -atmospheric pressure 4.2 Accelerating voltage and EB povver Contemporary EB vvelding machines utilize the electron guns of moderate povver of 10 - 150 kW vvhich operate in the range of 50 - 200 keV. The rea-son for increasing EB povver and accelerating voltage is to increase the maximum allovvable vveld depth. Hovvever, as could be seen in Fig. 12, the penetration range for electrons of above mentioned energies is 10 -100 pm only35. It means that the thermal energy is deposited essentially on the surface of materials. Therefore, the increasing of EB povver density over to =106 W/cm2 does not improve essen-tially the penetration depth and instead leads to surface vaporization of the vvorkpiece. The dramatic im-provement in penetration depth is possible only by increasing the EB energy to =10 MeV. At this energy the EB vvill interact vvith material in a way fundamen-tally different from the conventional EB. The electron range is novv of the order of 1 cm, what means that material is heated volumetrically. The energy re-quired to melt the material is delivered by several pulses of very high energy EB. Moreover, at EB energies greater than fevv MeV the gas-scattering of electrons drops dramatically (Fig. 13) causing a significant improvement in penetration of EB through the air. Additionally, if the peak EB current is also increased to approximately 1 kA, the self-pinch phenomenon becomes strong enough to resist the spreading of EB during its propagation through the atmosphere. This type of "pinched" EB enables to vveld thick-section vvorkpieces (up to 30 cm) at atmospheric pressures. Besides, such high energy EB can be used also as large-area, nearly in-stantaneous, volumetric heat source for surface treatment. 10 o 10-2 t lil -* 10 -<> ------- Aluminium / x Iron / // s 0,01 0,1 1 10 MeV 10(1 Kleclron energ\ Figure 12: Electron range in iron and aluminium 4.3 Povver supplies The traditional HV transformer rectifiers are stili in use. Nevertheless, the adoption of solid state inver-tors typically of a 5000 Hz base frequency seems to be very promising91232. This approach greatly simpli-fies the design as vvell as cost of the high voltage transformer and cable terminations. Moreover, these povver supplies facilitate the discharge control. It is possible to svvitch off the povver inverter vvithin 20 lis of detection of the current rise31. For the traditional EB vvelding machines vvhich operate at voltages less than 200 kV the standard high voltage DC povver supplies are adequate. Hovvever, for equipment of higher voltages and povvers other high voltage tech- tron beam pulsed accelerators has been developed, designated to accelerating potentials of few MV, beam currents of thousands of amperes, pulse dura-tion of tens to hundreds of nanoseconds, kilojoules of EB energy, and instantaneous povver of gigavvatts to terravvatts. These accelerators has been developed to a state in vvhich the parameters required for vvelding and materials treatment can be met vvith minimal additional development43. 4.4 Guns and cathodes Predominantly the conventional triode guns are used, but recently the return to diode gun configura-tion is also observed31. The diode electron gun en-ables to precise control of EB current over the full op-erating range vvithout risk of sudden current increase during breakdovvn. The displacement of EB vvaist is also reduced. In čase of very high energy EB (=MeV) the diode gans are used. Depending on the applica-tion and the type of accelerator, the anode of this gun can be made out of grid or thin foil, or it may have an aperture to allovv the EB injection into the beam transport pipe. The cathode of very high current diode gun is often made out of graphite. Very high electric field (1 MV/cm) causes the so called explo-sive emission. No external heating of the cathode is required43. The design of an electron gun by numerical methods taking into account the influence of many various electrical and geometrical parameters takes a lot of tirne and vvork but can easily be done employing modem computer technique. The detailed examina-tion of the emittance of the beam and application of mentioned earlier design of experiments have proved to be useful tool in improving the high povver electron guns applied in EB vvelding machines111314. 5. Conclusions As deseribed above the development of advanced electron beam technology vvill be promoted in two main direetions. The first consists in the enhance-ment of joint quality and reliability. Progress in this area vvill require intensive research of fundamental phenomena, clarifying the vvelding and surface treatment processes. The adoption of special methods of EB vvelding and surface treatment as well as imple-mentation of nevv materials destined for EB vvelding should significantly extend the scope of applications and performance of the process. The second direc-tion of EB development is connected vvith the thick-seetion vvelding and surface hardening of large vvork-pieces vvith complex shapes at atmospheric pressure. High energy EB of extremely high povver may be utilized to meet the grovving demand for high quality vvelding of thick and large vvorkpieces of re-fractory and fusible materials at or even above atmospheric pressures and for large area surface treatment of components vvhose large size and com-plex shapes make more conventional surface treat-ments time-consuming and expensive. EB vvelding Electron range Figure 13: Rate of electron energy loss in air at atmospheric pressure Figure 14: High-voltage accelerating system. ML - povver supply for magnetic lens, 1 - electron gun povver supply, 2 - EB accelerating povver supply niques are required and non-pulsed povver accelerators have been developed241 (Fig. 14). EB currents greater than 0,1 A cause in this type of accelerators serious beam instabilities39. Thus, these accelerators are not expected to scale to the higher voltage and povver levels at vvhich the pulsed povver options is most attractive. In this čase the high energy elec- and surface treatment can offer the comprehensive range of versatility and can be stili regarded as very competitive vvith other conventional methods. 6. References 1 P. Anderl, J. Koy, W. Scheffels, Proc. 4th CISFFEL (Cannes 1988), 733 2 Y. Arata, ibid. 1, 21 3 P. Bonnin, M. Aji, ibidl, 325 4 B. Dalmastri, R. Festa, Weld. Int., 10, 1989, 878 51. Decker, C. Oestmann, Bander Bleche Rohre, 30, 1989, 2, 22 6 K. Depner, F. Eichhorn, W. Janseen, K. Iversen, M. Engberding, ibld. 1, 301 7 M. Dijols, J-C. Goussain, G. Colombe, P. Grenier, Proc. 5th CISFFEL (LaBaule 1993), 339 8 J-L. Doong, J-M. Chi, J-R. Hvvang, Fatig. Fract. Eng. Mat. Struct., 13, 1990, 3, 253 9 C. D. A. Eccleston, ibid.1, 795 10 J. Felba, K. P. Friedel, Proc. Int. Conf. EBT (Varna 1991), 336 11 J. Felba, K. P. Friedel, W. Sielanko, S. VVojcicki, A. Wymyslowski, Proc. 4th Int. Conf. EBT (Varna 1994), 63 12 J. D. Ferrario, S. P. Kyselica, G. S. Lavvrence, Ibid. 1, 69 13 K. P. Friedel, J. Felba, ibid.1, 77 14 K. P. Friedel, J. Felba, ibid.11, 69 15 D. Fritz, IIW Doc. no IV-453-88, 1988 16 D. Fritz, Proc. Int. Conf. "Special Technologie '90" (Pilzen 1990), 115 171. M. Frolov, J. G. Kutsan, V. P. Morotchko, V. V. Gu-movsky, S. N. Kovbasenko, B. F. Jakushin, ibid.10, 172 18 H. Irie, S. Tsukamoto, ibid.1, 123 19 A. A. Kaidalov, V. J. Lokshin, O. K. Nazarenko, ibid.16, 45 20 A. M. Kosechek, ibid. 16, 3 21 A. M. Kosechek, ibid.16, 134 22 S. N. Kovbasenko, K. A. Sukach, M. L. Zhadkevich, J. G. Kutsan, Svar. Proizv., 6, 1988. 5 23 S. N. Kovbasenko, J. G. Kutsan, B. N. Shipitsyn, ibid. 10, 318 24 O. K. Nazarenko, Mat. Manuf. Process. 7, 1992, 2. 285 25 M. Ohara, H. Homma, T. Inoue, ibid.1, 227 26 B. E. Paton, et al., Weld. Int., 6, 1989, 466 27 V. I. Perevodchikov, S. I. Gusev, M. A. Zavialov, V. F. Martynov, ibld. 7, 99 28 C. S. Punshon, ibid.1, 287 29 P. J. Ross, Taguchi Techniques for Quality Engineering. McGravv-Hill, New York, 1988 30 R. Roudier, Ph. Dard, G. Sayegh, ibid.1. 767 31 J. D. Russell, ibid.7. 13 32 A. Sanderson, ibid.7. 91 33 G. Sayegh, Proc. 4th Int. Symp. IWS (Osaka 1982). 127 34 F. Shibata, Trans. JWS, 22, 1991, 2. 18 35 A. C. Smith, W. M. Fawley, E. E. Nolting, Proc. HEEB Weid. and Mat. Proc. (Cambridgre, USA. 1992), 1 36 J. Sommeria, A. Metz, ibid.1. 813 37 H-D. Steffens, E-R. Sievers, Mat. wiss. werkst. Tech.. 21 1990 235 38 M.Tomie, N. Abe, X-Y. Yao. Y. Arata. Trans. JWRI. 17. 1988, 2, 19 39 M. Tomie, N. Abe, Y. Arata, Trans. JWRI, 18, 1989. 1, 31 40 M. Tomie, N. Abe, Y. Arata. Trans. JWRI, 18, 1989, 2, 13 41 M. Tomie, N. Abe, Y. Arata, Trans. JWRI, 19, 1990, 2, 15 42 Y. Tomita, K. Koyama, K. Tanabe, ibid.7, 331 43 B. N. Turman, M. G. Mazarakis, E. L. Neau. ibid.35, 44 44 J. M. Vitek, S. A. David, Metali. Trans. 4.. Phys. Metali. Mat. Sci., 11 A, 1990, 7, 2021 45 S. Zolotovsky. ibid.1, 147 Operational Aspects of Experiences in Vacuum Technology by Production of High Quality Stainless and Alloyed Steels Praktične izkušnje pri uporabi vakuumske tehnologije pri izdelavi visoko kvalitetnih nerjavnih in legiranih jekel Koroušič B'., IMT, Ljubljana, A. Rozman, Metal d.o.o. Ravne, J. Triplat, Acroni d.o.o Jesenice, J. Lamut, Met & Mat. Department, University Ljubljana Highlights of the technological scheme of manufacturing quality steels as stainless steels in the vacuum by the duplex EAF (electrical are furnace) + VOD (vacuum oxygen decar-burization) process Is presented here. Speclal attentlon is given to the stage of vacuum oxidation of carbon and chromium and to the use stage of degassing in vacuum. In short feature the VOD computer program is expalained. vvhich enables the effective prepara-tion of the complete technology for VOD - process from 15 to 100 tonnes. Key vvords: production of high quality stainless steels by VOD process Predstavitev tehnološke sheme izdelave kvalitetnih jekel v vakuumu kot so nerjavna jekla po duplex postopku EOP (elektro-obločna peč) + VOD (vakuumsko kisikovo žilavenje). Poseben poudarek je podan na procesno stopnjo vakuumske oksidacije ogljika in kroma ter izkušnje pri uporabi vakuuma. Na kratko bo predstavljen tudi računalniški program oziroma rezultati modeliranja, ki omogočajo hitro in učinkovito pripravo kompletne tehnologije VOD procesa za peči kapacitete 15 do 100 ton. Ključne besede: izdelava kvalitetnih, nerjavnih jekel po VOD postopku Introduction The steel industry faces a number of problems related to an inereasing demand for clean steel im-posed by the consumer industry as vvell as inereasing costs pressures. For these reasons, the steel industry is in the past reasoning its steelmaking tech-niques vvhich resulted in a subdivision into tvvo phas-es, namely: • metallurgy in the melter and, • metallurgy In the ladle. Accordingly, most of the metallurgical vvork shifted from the melter to the ladle vvhich became a metallurgical reactor. The great number of techniques de-veloped in the secondary metallurgy is based on the application of complex vacuum systems. Before choosing among these techniques, the steelmaker must be perfectly avvare of the objeetives to be achieved and the requirements to be satisfied by the steel product and reasonable in priče for the cus-tomer, vvhile stili generating a profit for the steelmaker. The technology routes for secondary metallurgy are in no way a standardized task, even vvhere the ' Prof. Dr. Blaženko KOROUŠIČ, IMT. Lepi pot 11, 61000 Ljubljana underlying basic processes are identical. The condi-tions differ too much from shop to shop. The state of technology by production of high quality stainless steels in Slovenia steelvvorks vvill be reported in this lexture, highlighting the vacuum processes using a typ čase - stainless steel production as an example. General overvievv of vacuum processes The principles of secondary metallurgical processes used the vacuum can be brierfly deseribed as fol-lovving remarks: 1. Early in the 1950s, the vacuum degassing of molten steel vvas developed. The purpose was to re-duce the hydrogen content in steel and particulary in forging ingots. The best results vvere achieved by stream degassing during ladle - to - ladle or ladle -to - ingot teeming1. 2. In 1955 appeared almost simultaneously the RH and DH processes vvhich led to the development of degassing facilities vvith much higher throughputs. 3. In 1956 the first argon injeetion trials vvere carried out - the method of stirring and heating vvhilst degassing - to enable large reaction volume necessary and the comparatively high temperature losses in-curred led to the development of nevv advanced metallurgical techniques as: ASEA-SKF, FINKL, VAD, VOD combining nevv induetive stirring, nevv refraeto- ®J Figure 1: Processing units in the secondary metallurgy Slika 1: Procesne naprave v ponovčni metalurgiji ry lining materials, a sliding gate valve and oxygen lancing into the vacuum degassing unit29. With the industrial application of the above processes, especially vacuum steelmaking equipment passed through several stages of development, see (Fig. 1). The metallurgical possibilities of vacuum steelmaking have dramatically changed the role of the melting furnace. No longer is it necessary to use a particular type of furnace for the production of partic-ular grades of steel. In other vvords, ali furnace can produce nearly ali steels if equiped vvith a secondary steelmaking facility. Therefore, the demand for a high availability of the vacuum units is one of the essential factors and the modular concept has proved to be useful in the se-lection of suitable equipment for each application. Steel production in Slovenia steelvvorks vvas based for more than 120 years on melting in hearth fur-naces. During the 1960's and 1970's, the production moved progressively from open heart to electric are furnace melting. Vacuum refining began in Steelvvork Ravne in 1970 vvith the installation of the vacuum degassing unit for the removal of hydrogen, particularly for the forging shop. The process did not find a vvide application. Vacuum refining in Ravne began in 1984 vvith the installation of two VAD/VOD units (20 and 50 tonnes) and some years later also in Jesenice vvith tvvo 90 tonnes VOD-units vvith the aim to manufacture more sophisticated steel grades (very lovv S, O, H, C con-tents and higher contents in alloys - first of ali stainless, dynamo, tool and other high alloyed steels)10. 1 Some fundamental aspects of metallurgical processes in vacuum An atmosphere at reduced pressure is one of the purest environments possible. Most of the metallurgical processes are made up of heterogeneous reac-tions, and the operation in vacuum inereases the dri-ving force for interphase mass transfer vvhich is thereby accelerated. Vacuum processing, compared to air handling, re-moves the difficulties due to the pick up of oxygen, ni-trogen and hydrogen contained in the atmosphere. On the other hand, vacuum processes are concerned vvith the removal of hydrogen, oxygen, carbon and unvvanted non-metallic inclusions from the melt. These benefits are often obtained at the priče of same side effects as, for example, the difficulties caused by interaetion betvveen melt and refractory lin-ings. Furthermore, just the development of new re-fractory lining led tovvards a successful application of vacuum treatment of the molten steel. With introduction of basic linings (magnesite or dolomite) into vacuum ladles began the era of vacuum metallurgical processes. Activity of oxygen (ppm) - p(CO) = l bar --P h 35 30 25 20 0.50 16 1 4 1 2 10 O! m c d) TJ 700 units). Due to the action of high povver density EB, deep melting of the substrates occured. The layers be-came thick (0,6 - 0,7 mm) and significantly less hard (HV 0,1 < 500 units). The results of comparative vvear tests are presented in graphical form in Fig. 6. An average linear vvear of 3 samples produced by each method vvas a mea-sure of abrasion as function of a friction path. Layers obtained by oscillating EB method vvere extremely abrasion resistant. 4 Conclusions 1. The application of oscillating EB method allovv to produce thick, hard, vvear-resistant, Ni-base surface !ayers vvith satisfactory adhesion to the low-alloyed structural steel. 2. It is possible to control the appearance, structure and properties of the modified surface by process parameters: - the thickness of layers and vvidth of HAZ rise vvith the increase of EB povver, - the depth of melted substrate zone rises in a sim-ilar way, vvhich leads to the dilution of Ni-base alloy by Fe and the disappearance of Cr segregation, - the hardness of layer decreases due to formation of Ni-Fe-Cr solution. 3. The surface layer obtained by the presented method by optimal parameters (21 kV accelerating voltage, 45 mm/min. movement speed, 45 mA beam current) shovved a better vvear resistance in comparison to samples produced by conventional methods. The financial support by the State Committee for Scientific Research is greatly acknovvledged. 5 References 1 Y. Arata, Plasma, eleetron and laser beam technology, Am. Soc. of Metals, Metals Park, OH, 1986 2 M. Tomie, N. Abe, M. Yamada, S. Noguchi, Trans. JVVRI, 19, (1990), 1 3 J. Senkara, J. Mater. Sci. Lett., 10, (1991), 1078 4 Y. Q. Yan, J. Senkara, W. Wlosinski, Surf. Coat. Technol., 48, (1991), 211-217 61000 LJUBLJANA, LEPI POT 11. POB 431 SLOVENIJA INSTITUTE OF METALS Telefon: 061/1251-161, Te!efax: 061 213-780 AND TECHNOLOGIES p.o. VACUUM HEAT TREATMENT LABORATORY Vacuum Heat Treatment Vacuum Heat Treatment is recognised as a high quality cost effective and ultra clean method for processing a wide range of components and materials currently in use in today's industry. The range of our equipment enables us to heat treat most sizes of load, from small batches to work up to 350 mm diameter, 910 mm high, and weight up to 380 kg. ADVANTAGES • Clean, bright surface finish • Minimal distortion • Minimal post treatment operations, e.g., grinding or polishing Five years of continual investment has ensured that VHTL maintains it position as market leader in the field of high quality sub-contract metal processing. We operate the latest generation of IPSEN VTTC furnace capable of processing components up to 350 mm in diameter, which in addition to our high pressure, rapid quenching facilities increases the range of materials suitable for Vacuum Heat Treatment. TYPICAL APPLICATIONS • Bright Annealing • Bright Stress Relieving • Hardening/Tempering • Brazing/Hardening/Tempering • Solution Treatment Demagnetisation Degassing Diffusion Treatments Sintering QUALITY ASSURANCE Quality is fundamental to the IMT philosophy. The choice of process, ali processing operations and process control are continuously monitored by IMT Quality Control Department. The high level of quality resulting from this tightly organised activity has been acknowledged by government authorities. industry and International companies. Influence of Fracture Toughness on Vacuum Hardened HSS Vpliv lomne žilavosti na vakuumsko toplotno obdelano hitrorezno jeklo Leskovšek V.,1 B. Ule, A. Rodič, IMT Ljubljana Fractures. macro-chipping and micro-chipping are ali effects by which cutting edges are de-stroyed. The ability of a steel to resist these phenomena is knovvn as its toughness HSS hovvev-er. possess an appreciable ductility, although the notched or even unnotched specimens tested m the pendulum test are not sensitive enough to discriminate betvveen high and low levels of toughness. Therefore, it becomes important to use a method of testing vvhich can detect small vari-ations m ductility. To establish the fracture toughness, the round-notched tensile specimens vvith a fatigue crack at the notch root ivas used. Fatiguing vvas done in as soft annealed condition After that. the vacuum heat treatment for the achievement ofoptimal vvorking properties vvas carried out and the fmal testing vvas performed. Our experiments confirm that the correlation based on the round-notched tension test can be successfully used to caiculate the critical fracture toughness On the basis of the above-mentioned experimental results. we vvere abie to compose a diagram vvhich simultaneously scoops the technological parameters of vacuum heat-treatment the mechanical properties and the micro structure of vacuum heat-treated HSS M2. Key vvords: fine blanking tool. fracture toughness, hardness, vacuum heat treatment Lomi, makrookruški in mikrookruški so vzrok propadanja rezilnih robov. Sposobnost jekla da se upira tem pojavom, pa je poznana kot žilavost. Hitrorezno jeklo ima upoštevanja vredno duktilnost četudi preizkušanci z zarezo ali celo celo brez zareze pri Charpyjevem preizkusu niso dovolj selektivni, da bi nam omogočali določitev krhke oz. žilave narave loma. Za krhke materiale med katere spada hitrorezno jeklo, je pomembno, da izberemo metodo preizkušanja, ki zazna že majhne spremembe duktilnosti jekla ter je selektivna in reproduktivna. Poleg standardnega načina merjenja lomne žilavosti na preizkušancih, ki so dovolj debeli, da je izpolnjen pogoj ravninskega deformacijskega stanja, uporabljamo tudi nestandardni način merjenja lomne žilavosti s cilindričnimi nateznimi preizkušanci z zarezo po obodu. Problemi pri ustvarjanju razpoke v korenu zareze, so nas navedli na idejo, da metodo za določevanje lomne žilavosti s pomočjo cilindričnih preizkušancev z zarezo po obodu modificiramo. Doseženi rezultati so pokazali, da je modificirana metoda tudi dovolj selektivna. Osnovni namen modifikacije je, ustvariti razpoko kontrolirane globine v korenu zareze na mehko žar jenih cilindričnih preizkušancih z zarezo po obodu Pred pulz i rane cilindrične preizkušance zatem vakuumsko toplotno obdelamo, temu pa sledi natezni preizkus. Na osnovi rezultatov dobljenih s pomočjo modificirane metode, smo uspeli na istem diagramu zajeti mehanske lastnosti, tehnološke parametre vakuumske toplotne obdelave in mikrostrukturo vakuumsko toplotno obdelanih preizkušancev iz hitroreznega jekla M2. Ključne besede: orodje za precizno štancanje, lomna žilavost, trdota, vakuumska toplotna obdelava 1. Introduction A carefully selected vacuum heat treatment process improves the basic characteristics of HSS M2. The required vvorking hardness and fracture toughness of HSS is determined mainly by the hard-ening and tempering temperatures, depending on the alloying1. With the optimai vacuum heat treatment process, the best possible combination of fracture toughness and hardness, and therefore, wear resistance, is reached. Vojteh LESKOVŠEK. dipl. inž,, IMT. Lepi pot 11, 61000 Ljubljana The design calculations of HSS tools must consid-er the material strength, vvith a special emphasis on fracture toughness, because of the danger of brittle tool fracture. Fracture toughness is defined as the ability of a material under stress to resist the propa-gation of a sharp crack. To establish the fracture toughness of HSS in hardened and tempered conditions, a non-standard testing method vvith small-scale specimens vvas developed. This method in-volves the introduction of a sharp crack at the notch root, in our čase, by pulsating round-notched tension specimens, thermal treatment and tensile testing. A high level of hardness makes round specimens greatly sensitive to notches, so the test can fail due to unsuccessful pulsating. When successful prepul-sating, a fatigue crack is performed at the notch root of the specimen. The method was modified vvith the formation of a circumferential crack of defined depth at the root of the machined notch on soft annealing specimens, than a tensile test was performed after vacuum heat treatment. Our experiments confirm that the measurements based on the modified round-notched tension test can be successfully used to determine the fracture toughness. 2. Basic characteristics of high speed tool steel M2 Due to the higher vvear resistance of HSS, they are nowadays used also for fine blanking, cold vvorking and deep dravving tools, especially in long series. Tool steels must vvithstand compressive stresses and abrasive or adhesive vvear, vvhile have a suffi-cient toughness to resist chipping and failure. HSS have better resistance to vvear in comparison to cold vvork tool steels because of the increased hardness of the matrix, and of the carbide phase. The carbide phase in the matrix of HSS increases the vvear resistance vvhich is relative to the total vol-ume of carbides, and also to their hardness. The vvear resistance in HSS is mainly determined by vanadium carbides vvhich have a micro-hardness of 2200 to 2400 HV2 3, (Fig. 1). —i_i ■_i_i i i VC (OPel-jCj (FeWMo)6C. FČ3C. Matrixhigh Matnx speed steel carbon steel Figure 1: The comparative hardness of carbides found in tool steels2 Slika 1: Primerjalne vrednosti trdot karbidov, ki jih najdemo v hitroreznih jeklih Hovvever, it must not be forgotten that HSS have a greater hot hardness. Even if the vvork pieces are plače into the tools vvhile cold, the vvorking tool sur-faces become hot. Fractures, macro-chipping and micro-chipping can destroy the cutting edges. The ability of a steel to resist these phenomena is knovvn as toughness. The toughness that can be achieved by HSS is limited by the defects in the steel (carbide segregations and bands inclusions etc.). When the steel is subjected to a load, stress concentrations can appear around the defects and cause a tool fracture, unless the stress concentrations are relieved by a local plastic flovv on the micro scale. The ability of the matrix to undergo plastic flovv can be altered vvithin wide limits by varying the hardness. Thus, the defects in the steel determine the maximum toughness vvhich can be achieved. On the other hand, the heat treatment determines the toughness degree actually achieved vvithin the limits set by the defects. Vacuum heat treatment is one of the most important operations in the manufacturing of tools. Therefore, vvhen HSS are used for cold vvorking processes, the situation is met by choosing low hard-ening temperatures and tempering temperatures be-lovv the peak secondary hardening temperature, to improve fracture toughness, cutting edge strength, vvear resistance and dimensional stability. It is possi-ble to exert a positive influence on ali the parameters by vacuum heat treatment vvhich is carefully select-ed to suit the HSS is determined by a choice of vari-able tempering temperatures, it is often impossible to optimise the mechanical properties, e.g. fracture toughness. A general recommendation for tools that require good impact strength, such as fine blanking tools, is that they should be hardened from temperatures as low as 1050°C1. By this treatment, resistance to tempering is diminished. For tools subjected to high pressures in service, a previous tempering at about 600°C1 is recommended. 3. Experimental procedure 3.1 Material and treatment parameters The test material selected was a conventional high-speed steel (HSS) of the AISI M2 type of the same melt. The chemical composition of the steel ex-amined is listed in Table 1. Table 1: Chemical composition of HSS examined (in wt.-%) Material C Si Mn Cr Mo W V Co AISI M2 0.87 0.29 0.30 3.77 4.90 6.24 1.81 0.53 Cylindrical round-notched tensile specimens vvith a diameter of 10 mm vvere machined from soft an-nealed bars vvith a Brinell hardness of 255. Specimens vvere fatigued to produce a sharp circumferential crack at the notch root, then austeni-tized in a vacuum furnace at temperatures of 1050°C, 1100JC, 1150°C and 1230 C respectively, quenched in a flovv of gaseous nitrogen at a pressure of 5 bar abs. and double tempered one hour at temperatures 510°C, 540°C, 570DC and 600"C respec-tively. 3.2 Mechanical tests The geometry of cylindrical round-notched pre-cracked tensile specimens, prepared according Dieter's recommendation4 is shovvn in Fig. 2. Our previous investigations56 confirmed that such small-scale specimens can be successfully used for the analysis of the relationship betvveen the mi-crostructural variations and the fracture toughness of the investigated steels. Accordingly to Grossmann's concept of hardenability, the formation of the uniform microstructure along the crack front is possible because of, the axial symmetry of the cylindrical specimens, in comparison vvith the conventional CT-specimens, vvhere this condition is not fulfilled. vvhere oys is the yield stress of the investigated steel. This requirement (2) vvas fulfilled on ali our measurements. The fracture surface of the cylindrical round-notched and precracked specimens vvas ex-amined in SEM at low magnification. As is shovvn in Fig. 3, the fatigue crack propagation area vvas sharply separated from the circular central part of the fast fracture area, so that the diameter d of this area vvas easily measured. - P Figure 2: The geometry of a cylindrical round-notched and precracked tensile specimen Slika 2: Nestandardni cilindrični natezni preizkušanec za merjenje lomne žilavosti z zarezo po obodu ter utrujenostno razpoko v korenu zareze For a round-notched precracked specimen, the stress intensity factor is given by Dieter4 as KI =-^-(-1.27+ 1.72 D/d) (1) vvhere d is the radius of the uncracked ligament after fatiguing, P is the applied fracture load, and D is the outer diameter of the cylindrical specimen. In the ex-periments, it is essential for the outer diameter of the specimen in order to obtain a state of plain strain at fracture. In order to apply linear-elastic fracture mechanical concepts, the size of the plastic zone at the crack tip must be small compared vvith the nominal dimen-sions of the specimen. The size requirement for a valid KIC test is given by Shen VVei et. al.7 as D > 1.5 (K|C/crys) (2) Figure 3: Fracture surface of cylindrical round-notched and precracked tensile specimen vvith the circumferential fatigue crack propagation area sharply separated from the circular central fast-fractured area Slika 3: Prelomna površina cilindričnega nateznega preizkušanca z obodno zarezo, s kolobarjastim področjem propagacije utrujenostne razpoke, ki je ostro ločeno od osrednjega, naglo zlomljenega dela. Premer (d) naglo zlomljene prelomne površine lahko izmerimo z optičnim mikroskopom 4. Results and discussion 4.1 Microstructural characterisation The microstructure develops in dependence on the hardening temperature, as vvell as the austenite grain size and the residual austenite content of the initial samples. Metallographic examination of specimens shovv that the austenite grain size of ali specimens vvhich vvere gas quenched from the austeniti-zation temperature 1050 to 1230°C vvas 21 to 8 SG, (Fig. 4). The content of residual austenite in as queched condition vvas determined by X-ray diffraction. The absolute accuracy of the determination of the residual austenite contents vvas ± 1 vol%. The HSS AISI M2 steel is fine-grained, right up to high hardening temperatures, and exhibits a residual austenite contents betvveen 21 and 30 vol%. - Tj, The microstructure after hardening at 1050 C, SG 21. Mag. 500x. Mikrostruktura po kaljenju s temperature 1050°C, SG 21. pov. 500x. The microstructure after hardening at 1100 C. SG 18. Mag. 500x. Mikrostruktura po kaljenju s temperature 1100 C. SG 18. pov. 500x. The microstructure after hardening at 1150 C, SG 13. Mag. 500x. Mikrostruktura po kaljenju s temperature 1150 C, SG 13. pov. 500x. The microstructure after hardening at 1230 C. SG 8. Mag. 500x. Mikrostruktura po kaljenju s temperature 1230 C. SG 8. pov. 500x. Figure 4: Microstructures vvith a different austenite grain size from vacuum hardened specimens from M2 steel Slika 4: Mikrostruktura in velikost austenitnih zrn, vakuumsko kaljenih vzorcev z različnih temperatur austenitizacije Figure 5: The microstructure shovvs carbides and tempered martensite vvith an austenite grain size of 13 SG (TA:1150 C) and 8 SG (TA:1230°C) Slika 5: Mikrostruktura karbidov in popuščanega martenzita z velikostjo avstenitnih zrn SG 13 (TA:1150 C) in SG 8 (Ta:1 230° C) The carbide particles in ali the specimens were alike in size and position, vvhich vvas due to their ori-gin: ali the specimens issued from the same metal-lurgical melt vvhich vvas submitted to the same hot plastic transformation. The carbides looked like strips, and had a size of 1 -20 (.im, (Fig. 5). The resid-ual austenite contents are, vvith reference to temper-ing parameters, belovv 1 vol% in ali samples. After metailographic etching, a stronger marking of the austenitic grain boundary could be noticed. L. Calliari Austenitizing temperature Figure 6: The microstructure of vacuum-hardened and tempered specimens examined by SEM Slika 6: Mikrostruktura vakuumsko kaljenih in popuščenih vzorcev, posnetih na SEM pri 10 000 kratni povečavi Table 2: Vacuum heat treatment parameters and mechanical properties of prepulsating round-notched tesion specimens Vacuum heat treatment Fracture Hardness toughness Spec. Hardening( C) Tempering( C) HRc KIC (MNm32) No. 2 min. 2 x 1h 01-02 1050 510 60.0 18.78 03-04 1050 540 60.5 18.26 05-06 1050 570 58.7 15.80 07-08 1050 600 52.8 16.43 09-10 1100 510 61.8 17.28 11-12 1100 540 62.2 15.69 13-14 1100 570 61.3 15.49 15-16 1100 600 55.0 16.99 17-18 1150 510 60.7 18.26 19-20 1150 540 63.3 13.14 21-22 1150 570 63.2 14.70 23-24 1150 600 57.8 15.63 25 1230 510 62.5 17.77 26 1230 540 65.0 10.55 27 1230 570 65.5 12.08 28 1230 600 63.0 12.95 et al.8, compared vacuum and conventional heat-treated samples of AISI M2, and found that the re-sults of over 100 tests did not point out noticeable dif-ferences among the samples treated vvith the two different procedures. Neither systematic data nor re-lationship vvith the treatment parameters are yet available on this subject. The microstructure of the specimens examined by SEM at a higher magnification (Fig. 6) confirmed a carbide precipitation on the austenite grain bound-aries for HSS M2 at different austenitizing and tem-pering temperatures. The quantity of fine carbide particles decreased vvith the increase of austenitizing temperatures. In addition, it vvas also noticed that at higher austenitizing temperatures, particularly at 1230 C (last column in Fig. 6), the larger carbide particles in contacts of austenite grains seemed to covering the boundaries of the neighboring grains because of variable surface tension on the matrix-carbides boundary. The microstructure of the specimens vvas of martensite type. The eventual presence of small quantities of retained austenite (1 to 5 vol%)8 examined by optical microscopy, vvere too small to estimate vvithout fail in such a heterogeneous microstructure. This phenomena can be attributed to the fact that the heating rate vvas lovver in the vacuum furnace than in the salt bath. By heating the pre-pulsating round-notched tension specimens betvveen 1050 C and 1230 C, diffusion processes in the vacuum furnace took longer than in the salt bath, vvhich can possibly explain why, after metallograph-ic etching, a more intensive marking of the austenitic grain boundary can also be noticed. 4.2 Mechanical tests Experiments9 vvere performed on 28 prepulsating round-notched tension specimens, (Table 2), heat-treated in an IPSEN VTTC-324 R single chamber vacuum furnace vvith uniform high pressure gas quenching. In the follovving, the assumption is made that the values KIC are determined by the above-mentioned method. The obtained values of KIC are very similar to those obtained by G. Hoyl10, vvho determined the fracture toughness KIC for HSS M2 steel, e.g. 18,3 MNm3;' for sample austenitized 4 minutes at 1220 C and tempered 1 hour at 510 C, by conventional methods. Belovv a hardness level of about 50 HRc, fracture toughness is dependent only on the hardness of the sample10. At higher levels of hardness, the fracture toughness for M2 varies in-versely vvith the austenitizing temperature, as shovvn in Fig. 7. G.Hoyl10 discovered that above a hardness level of 60 HRc, fracture toughness is in-sensitive to most metallurgical factors. The effects of tempering betvveen 510 and 600 C on fracture toughness are shovvn for M2 in the same figure. As expected, there is a minimum of toughness values corresponding to the hardness peak. The net effect of tempering is attributed to a combination of stress relief and a reduction in ductility due to the secondary hardening effect. F T1050 + H1050 ' FTIIOO o H1100 * FT1150 0 H1150 a FT1230 * H1230 65 H d 55 n e 45 H R 35 Figure 7: The effect of austenitizing and tempering temperature on fracture toughness and hardness of M2 steel (FT-fracture toughness; H-hardness) Slika 7: Vpliv temperature austenitizacije in popuščanja na lomno žilavost KIC in trdoto, vakuumsko toplotno obdelanega hitroreznega jekla M2 (FT-lomna žilavost, H-trdota) In examining the evolution of tempering10, it vvas found that there vvas a peak value of fracture toughness for low tempering temperatures, (belovv 500r C). The obtained values are similar to those obtained in tempering at the conventional temperature, 25 C above the peak of the secondary hardening temperature. This is considered as advantageous, but as the effect is due to retained austenite that could transform later, the under-tempered tools could be susceptible to dimensional instability in ser-vice, vvhich is unacceptable for fine blanking tools. On the basis of the experimental results, it vvas possible to dravv the diagram shovvn in Fig. 8 vvhere the technological parameters of the vacuum heat-treatment, mechanical properties and microstructure of the vacuum heat treated specimens are simulta-neously combined. From the diagram in Fig. 8. it is also evident that the fracture toughness for the tempering temperatures 540 C, 570 C and 600 C, respectively, in-creases vvith the decrease of hardening temperature in agreement vvith observations in reference 10. On the other hand, for the tempering temperature of 510' C, it vvas found that the fracture toughness values vvere very close, though slightly higher than for 600 C, irrespective of the hardening temperatures. On the basis of the curves in Fig. 8. it can be reliably assumed that the HSS M2 hardened from low austenitizing temperatures and tempered at 510 C can achieve the optimal combination of hardness and fracture toughness. Fracture toughness (MNm-3/2) 500 5Ž0 540 560 580 600 Tempering temperature (°C) --- TT 510 °C —t— TT 540 °C -*— TT 570 °C -o— TT 600 °C Austenite grain size SG Figure 8: The influence of austenite grain size on the fracture toughness of HSS AISI M2, (TT-tempering temperature, TA-hardening temperature, HRc-hardness at 510 C) Slika 8: Vpliv velikosti austenitnega zrna na lomno žilavost hitroreznega AISI M2, (TT-temperatura popuščanja, TA-temperatura kaljenja, HRc-trdota po pop. na temperaturi 510 C) The relationship betvveen fracture toughness and austenite grain size, f.i. SG grade 8, shovvs us that at the tempering temperatures betvveen 510 and 600 C, the obtained values KIC are from 17,77 to 10,55 MNm3 2 and the difference is quite important in practice. Different fracture toughness at equal austenite grain size or the nearly constant fracture toughness of HSS M2 hardened from different austenitizing temperatures, f.i. from 1050 to 1230 C, and double tempered at 510 C, is in accordance vvith the hypothesis that the austenite grain size is not the dominant parameter effecting the fracture toughness of HSS M2. The result of the present investigation is useful for the optimisation of vacuum heat treatment for different HSS tools submitted to tensile impact stress dur-ing use vvhere an optimal combination of hardness and fracture toughness are decisive. 4.3 Tool life tests Long production runs have underlined the impor-tance of an improved fine blanking tool life, (Fig. 9). On the basis of the experimental results shovvn in Fig. 8. it vvas found that the optimum vacuum heat treatment of fine blanking HSS M2 tools needs at least tvvo preheating stages (650 and 850°C respec-tively), a variable hardening temperatures (usually 1050 to 1150l C) and double tempering at the same temperature (510 C/1 hour). Figure 9: Fine blanking tool for a ratchet vvheel Slika 9: Orodje za precizno štancanje zobnika varnostnega pasu The life of a fine blanking tool varies considerably depending on the size and design of the blank, the type of blanking steel, and care and maintenance. To establish tool life, vve selected three tools for a fine blanking ratchet vvheel. The punches and dies vvere from HSS AISI M2. Blanks vvith a material gauge of 4 mm vvere from AISI C 1045 blanking steel in a spheroidized-annealed condition. The punches and dies vvere stress-relieved at 650 C 4 hours and vacuum heat-treated in a single chamber vacuum furnace vvith uniform high pressure gas quenching. Depending on the alloying and the condition of austenitization (1100°C/2min), the final hardness of 61.5 HRc and the final fracture toughness of 17.28 MNm 3'2 vvas reached after double tempering at 510 C for 1 hour (Fig. 8). The vacuum heat-treated punches and dies vvere tested on a 250t triple-action hydraulic press and compared vvith the fine blanking tools for fine blanking ratchet vvheels conventionally heat-treated in a salt bath as follovvs: stress relieved at 650°C/4hour, preheated at 450 C, 650 C and 850 C respectively, austeni-tized at 1100 C/2min and control-quenched to 550°C in 5 min, held isothermally at 550°C for 10 min and cooled to 80JC vvith air, follovved by double tempering at 600°C/1 hour. The final hardness of the tools achieved after double tempering at 600°C/1 hour, vvas 58 to 59 HRc, depending on the alloying. The basic trial parameters (such as the cutting force, strikes per tirne unit, temperature, greasing etc.) vvere constant during the experiments, and did not affect the final results. The differences observed in fine blanking tools could only originate in the punches and dies themselves. 15.000 ratchet vvheels vvere made vvith each tool and edges exam-ined in a binocular microscope to estimate the dam-age. Minor defects vvere observed on the vacuum heat-treated punches and dies and those conven-tionally heat-treated in a salt bath. Figure 10: Defects on the punch cutting edge Slika 10: Poškodbe na rezilnem robu pestiča The wear of the punches and dies increases vvith the operation tirne. SEM observations show that the starting vvear of the cutting edges of the punches and dies could not be easily determined. The material shovved not only surface fatigue, but also adhesion and abrasion, (Fig. 10). The experiments shovved that higher vvorking hardness (61.5 HRc) and improved fracture toughness of vacuum heat-treated punches and dies - particularly those double-tempered at 510° C - had significant ef-fect on the defects on the cutting edges. Compared to conventionally heat treated tools tempered at 600UC vvas the vacuum heat-treated tool life by 15 to 20%, greater. During the tool tests, no effects vvere found that could be related to the in service dimen-sional instability due to the later-transformed re-tained austenite. Namely, X-ray structural analyses shovved that the content of retained austenite did not exceed 1 vol% in ali the specimens. 5. Conclusions The exact significance of fracture toughness as it affects HSS properties and service behavior is not thoroughly understood, and there are differences in behavior betvveen grades and process routes. Hovvever, the modified method for the establishment of fracture toughness improved by IMT, appeared to be a successful method for establishing the fracture toughness of HSS. The measurements of vvear in the cutting edges of punches and dies shovv that double tempering at 510°C/1 h after vacuum heat treatment prolongs the life of fine blanking tools for ratchet vvheels by 15 to 20%, compared to conventionally heat-treated tools vvhich vvere hardened at the same austenitizing temperature and tempered tvvice at 600°C. It seams that it is not the type of process vvhich substantially af- fects the tool life, but first of ali, the proper choice of hardening and tempering temperatures. The vvear resistance of punches and dies cannot only be described as a material property, but as a property of a complex tribological system. Yet, it is proven that the vvear resistance depends, above ali, on the microstructure of the tool material and on its physical and chemical properties. The presented results, obtained by the evaluation of daily confirmed data, justify the use of modem vacuum heat-treatment technology and the use of the nevvest high-performance HSS steels to achieve great improvements in tool lives and overall econo-my. 6. References 1 K. E. Thelning: Steel and its Heat Treatment. Bofors Handbook, 1975, 313 -319 2 R. VVilson: Metallurgy and Heat Treatment of Tool Steels, McGravv-Hill Book Company UK, 1975, 69 3 H. Czichos: Tribology, Elsevier Scientific Publishishing Company Amsterdam, 1978 4 G. E. Dieter,: Mechanical Metallurgy, McGravv-Hill. 1986, 358 5 B. Ule, D. Kmetic, A. Rodič: Rudarsko-Metalurški zbornik, 1989, 3, 509-519 6 V. Leskovšek, B. Ule, A. Rodič, D. Lazar: Vacuum. 43. 1992, 5-7, 713-716 7 Shen Wei et.al.: Engineering Fracture Mechanics. 16, 1982, 1, 69-92 8 L. Calliari, A. Molinari, E. Ramous, G. Torbol. G. Wolf: Vacuum and Conventional Treatment of Tool Steels, Instituto per la Ricerca Scientifica e Technologica 38050 Povo (Trento) ltaly, 49-57 9 V. Leskovšek, B. Ule, A. Rodič: Metals Ailoys Technologies, 27, 1993, 1-2, 195-204 10G. Hoyle, High Speed Steels, Buttervvorth & Co. Ltd. 1988, 143-146 Characteristics of Cemented Carbide Particles/Structural Steel Vacuum Brazing Joint Značilnosti vakuumskega spoja zrn karbidne trdine s konstrukcijskim jeklom Šuštaršič B., V. Leskovšek, A. Rodič, IMT, Ljubljana The strength of the vacuum brazing joint betvveen cemented carbide particles and the structural steel base depends on microstructural characteristics of the hard metal/braze interface formed during vacuum brazing. These are determined by the selected brazing agent, the structural steel base, as vvell as the vacuum brazing procedure. The R&D vvork concerning the procedure of manufacturing these types of grinding tools is introduced, vvith a strong emphasis on the characteristics ofthe brazing agent used. the vacuum brazing procedure, as vvell as the resulting microstructural characteristics of the brazing joint betvveen the hard metal particles and the structural steel base. Key vvords: vacuum brazing, cemented carbide particles, structural steels, Cu-based brazing alloy-powders, microstructural features Trdnost vakuumskega spoja zrn karbidne trdine s konstrukcijskim jeklom je odvisna predvsem od mikrostrukture mejne plasti kovinska osnova/karbidna trdina, nastale med vakuumskim trdim spajkanjem. Le-ta pa je odvisna od izbrane kovinske osnove (konstrukcijskega jekla), vrste karbidne trdine in tehnoloških parametrov spajkanja. V pričujočem prispevku je predstavljeno razvojno raziskovalno delo vezano na postopek izdelave brusnih plošč, ki so sestavljene iz jeklenih plošč na katere so nanešena groba zrna karbidne trdine. Poudarek je na opisu mikrostrukturnih značilnostih nastalega spoja v odvisnosti od uporabljene spajke, karbidne trdine in jekla, kakor tudi pogojev vakuumskega spajkanja. Ključne besede: vakuumsko trdo spajkanje, zrna karbidne trdine, konstrukcijska jekla, spajke na osnovi Cu, mikrostrukturne značilnosti 1. Introduction Tungsten carbides vvith Co matrix (WC-Co) are old and well-known composites, referred to as cemented carbides or hard metals. They are also well-known under the trade mark name VVIDIA. The most important application of cemented carbides is in the production of machining tools, as vvell as in the production of vvear resistant parts or layers in many fields of application. VVaste cemented carbide parts of vvorn tools (inserts, knives, drills, cutters, savvs, punches, etc.) can be ground and the resulting relatively rough and sharp edged particles of the WC-Co composite can be used for the manufacturing of grinding vvheels vvhich are fast and simply mounted on the electric drill. The cemented carbide particles are uniformly deposited on a clean surface of a steel base (grinding vvheel) and the diffusion bonding of particles vvith the steel base can be obtained by different methods of brazing. The most convenient brazing method is vacuum brazing. In vvood, stone-cutting and leather industry, as vvell as in rubber industry, it is often necessary (by a fast ■ Mag. Borivoj ŠUŠTARŠIČ dipl. inž., IMT. Lepi pot 11, 61000 Ljubljana and simple procedure) either to clean or to rub off the surface of the semi-finished products during individual steps of the manufacturing procedure. Stock-farming is gradually substituting for the hard and time-consuming trimming of hoofs by manual grinding. These kinds of grinding vvheels are also appropriate tools for housevvork applications. Currently, this type of grinding tools is also increasingly in demand on the Slovenian market. The researchers of the Institute of Metals and Technologies in Ljubljana, Slovenia, have many years of experience in different R&D fields concerning vvear resistant materials (metal povvder manufacturing, heat treatment, vacuum brazing, material development, investigations and testing, etc.). Therefore, a decision vvas made at the IMT Ljubljana to develop a procedure of manufacturing these types of grinding tools. For the brazing of sintered cemented carbides (WC-Co composites), either vvith tool or structural steel base as a brazing agent, technically pure cop-per (usually OFHC Cu 99.9 mass %) povvder is com-monly used123. The joining of WC-Co composites and steel base vvith pure Cu and the diffusion processes occurring in this connection, as vvell as the formation of different phases at the interface hard metal/braze and braze/steel base vvas the subject of several previous investigations1'4. The common statement of these is that, during brazing at the hard metal/braze interface, thin Co and Fe rich layer of an intermetallic compound is formed. Voids are also generated in the border zone of the hard metal vvhich greatly reduce the strength of the joint. Commercial producers56 of Cu based brazing agents often recommend Cu-Ni based povvders as a proper material for the brazing of cemented car-bides/steels couples. Simultaneously, our own povvder preparation experiments7 shovv that by the vvater atomization of Cu-2%Ni alloy, a nearly spherical povvder vvith relatively high apparent density and flovvabil-ity, as vvell as a proper particle size distribution vvith the mean particle size betvveen 40 in 70 pm can be obtained. Thus by the optimization of process para-meters of vvater atomization78-9 (especially of vvater pressure, tundish nozzle diameter and superheat) it can be concluded that the manufacturing of brazing povvder vvith required morphological properties is possible. VVater atomization experiments shovv7 that Cu-2%Ni alloy povvder has even slightly better morphological properties in comparison to pure Cu based vvater atomized povvder. In spite of the fact that vvater atomized povvders have a relatively high oxygen content in prepared Cu-2%Ni alloy povvder, good-quality joint betvveen hard metal and steel is expected, if de-oxidising atmosphere (vacuum) and good morphological properties of the povvder are taken into con-sideration. In the present article, the research vvork concerning the procedure of manufacturing these types of grind-ing tools is introduced, vvith a strong emphasis on the characteristics of the brazing agent used, the vacuum brazing procedure, as vvell as the resulting mi-crostructural characteristics of the brazing joint betvveen the cemented carbide particles and the struc-tural steel base. Considering the relatively vvell investigated process of multiphase diffusion4 during brazing vvith pure Cu, the aim of our vvork vvas also to find out if a similar phenomenon occurs in brazing vvith Cu-Ni alloy povvder. The influence of a steel base chemical composition is also considered. 2. Experimental vvork As a metal base for rough and sharp edged particles of the WC-Co composite, tvvo different steels (soft lovv-carbon structural steel Č.0561 or DIN W.No.: 1.0570 vvith 0.2 mass % C and low-alloy hardening steel Č.4733, Ravne Steel Plant VCMo150 or DIN VV.No.: 1.7228 vvith 0.5 % C, 1 % Cr and 0.2 % Mo) vvere selected. From the selected steels, round plates of standard grinding vvheels dimensions (diameter approximately 115 mm and thickness 2 h- 6 mm) vvere made. For deposition, a mixture of metal povvder (fraction 45 : 75 pm of Cu-2 % Ni), binder and cemented carbide (ISO G10/G20 vvith 6 + 10 % Co and nominal Vickers hardness 1500 HV) particles in size 1 -h 3 mm vvas prepared. After the deposition of the mixture on the steel base, the samples vvere vacuum brazed in IPSEN vacuum heat treatment furnace (type VTTC-324R). Optimal brazing conditions vvere achieved in vacuum 10"3 mbars, in temperature range betvveen 1100 and 1200°C, at heating rate = 20°C/min. and at cool-ing rate = 3°C/min. These brazing conditions ensure that the cemented carbide particles and the steel base remain in solid state during brazing, vvhile the braze is melted. This ensures good vvetting of cemented carbide particles and steel surface vvith the brazing agent. The applied metal povvder and the cemented carbide particles vvere examined by optical and scanning electron microscope. The apparent density and flowability of the prepared povvders vvas determined by HalPs apparatus in accordance vvith MPIF standards10 (Metal Povvder Industries Federation Standards No.: 03 and 04). The oxygen content in metal povvder vvas also determined. The chemical composition of selected steels vvas checked by X-ray fluorescence (ARL 3460 Metal Analyser). The hardness of individual components of grinding vvheels vvas determined before and after brazing. Samples from manufactured grinding vvheels (see Fig. 1) vvere then cut out for metalographic investigations and analysis vvith electron micro-probe analyzer (EPMA). 2cm i-1 Figure 1: Macroscopic photo of grinding vvheels manufactured at IMT Ljubljana Slika 1: Makroskopski posnetek na IMT Ljubljana izdelanih brusnih plošč 3. Results and discussion The vvater atomized Cu-2 % Ni povvder prepared at IMT Ljubljana, vvhich vvas used as the brazing agent in our experiments has mainly almost spherical particles (see Fig. 2). The povvder particles are coated vvith a thin oxide film. Metalographic examinations also shovv internal porosity of some particles. The cel-lular solidification structure of povvder particles (see Fig. 3) results from the high cooling rate (= 105 - 107 K/s) obtained during vvater atomization. The size of cells strongly depends on povvder particle size be-cause the cooling rate primarily depends on the particle size formed during atomization. High apparent density and good flowability are the basic features of a high-quality brazing agents. The diffusion of Fe is primarily influenced (sup-pressed) by the carbon content in steel4, vvhereas the influence of other alloying elements (Cr, Mn, etc.) has not yet been analysed in detail. In studying diffusion processes it has to be consid-ered that mutual solubility of presented components is very small or practically equal to zero (solubility of braze in cemented carbide). Namely, the solubility of the main individual elements presented in the brazed components (Fe, Cr, Co, W) in braze (Cu) is relative-ly small. For example, Co solubility in Cu at brazing temperatures is = 8 %, solubility of Cr is = 0.65 % and Fe solubility is = 5 %. But, W and C are practically in-soluble in liquid Cu. Figure 3: Optical micrograph of rapid solidified cellular microstructure of vvater atomized Cu-2%Ni povvder particles vvith noticed internal porosity Slika 3: Posnetek hitro strjene celične mikrostrukture delcev vodno atomiziranega prahu Cu-2%Ni z opazno notranjo poroznostjo Cemented carbide particles vvere analysed vvith EPMA. Electron images show the expected sharp edged particles and the expected distribution of W and Co contents. In trace, Fe, Ni, Mn, Si and Cr are also present. The determination of qualitative and quantitative distribution of C is not possible vvith our EPMA. 3.1 Characteristics of cemented carbide particies/structural steel brazing joint First, the brazing results analysis and a discussion of the brazing non-alloyed soft structural steel/ce-mented carbide vvith Cu-Ni braze vvill be presented. The selected steel has a fine grained ferrite-pearlite microstructure vvith approximately 15% of pearlite (see Fig. 4), low carbon content and Brinell hardness 205-215 HB. Fig. 5 shovvs the cross section steel/braze/cemented carbide particle vvith elements distribution obtained vvith EPMA. It can clearly be seen that Fe diffused to the interfaces braze/cement-ed carbide particles, vvhile W mostly remains in cemented carbide particles. Therefore, one can con-clude that the diffusion of W is negligible, or rather, that it is limited to the diffusion to the interface cemented carbide/braze. Ni is found in braze. With Ni enriched regions are also interfaces steel/braze and braze/cemented carbide. Because of regular povvder particle shape, the pre-pared povvder has a relatively high apparent density (=4.4 g/cm3) and good flowability (=18 sec./50 g). Figure 2: SEM micrograph of particle shape of vvater atomized Cu-2%Ni povvder used as vacuum brazing agent Slika 2: SEM posnetek delcev vodno atomiziranega prahu (spajke Cu-2%Ni) uporabljenega za vakuumsko trdo spajkanje karbidnih zrn z jekleno osnovo As it has already been mentioned, the vvater atomized povvders have a relatively high oxygen content (0.25 mass % of 02 in prepared Cu-2%Ni alloy povvder). This could be primarily attributed to particle surface oxidation during vvater atomization and surface adsorption of oxygen molecules during the handling vvith povvder. In Cu based alloy-powders, oxygen sol-ubility and oxides from slag" should also be taken into consideration. Individual contributions to the over-all oxygen content of povvder stili have to be established in our future investigations. The brazing procedure is carried out in relatively high vacuum and therefore the high oxygen content of povvder has little influence on the brazing quality. Because of the carbon content in the steel and in the cemented carbide, the reduction of oxygen vvith carbon is also possible. The Cu-Ni alloying system is one of complete solid solubility and therefore it can be concluded that the used braze is a substitutional solid solution. Besides the braze, vvhich is liquid at the brazing temperature, the diffusion processes (solid > liquid > solid) are influenced by the selected metal base (steel). Namely, because of the concentration gradi-ents (and the driving force for diffusion is the gradient of the chemical potential), Fe and other alloying elements presented in the metal base diffuse across the braze at the interface braze/cemented carbide. An increased amount of Co is noticed at the interface cemented carbide/braze and partially at the interface braze/steel. From this point of view it can be concluded that Co diffuses from the matrix of the cemented carbide particles to the interface cemented carbide/braze and across the braze to the interface braze/steel. The diffusion of Mn and Si is insignificant. On the basis of previous investigations124, element Figure 4: Microstructures of cemented carbide particle/structural steel joint of prepared grinding vvheel Slika 4: Mikrostruktura vakuumskega spoja zrn karbidne trdine s konstrukcijskim jeklom distribution obtained vvith EPMA and metallographic examination (see Fig. 4 and 5) it can be concluded that at the interface cemented carbide/braze a thin layer (= 20 ; 40 um) probably of an intermetallic com-pound enriched vvith Co and Fe vvas formed. Unfortunately, the determination of the C and O distribution is not possible by our EPMA. Therefore, data of other investigators have to be considered4. The diffusion of C from steel across the braze as vvell as by the decomposition of cemented carbide particles and then the diffusion of W and C to the interface cemented carbide/braze is possible. Studies of the Co-Cu-Fe-W quaternary system at the brazing temperatures and their subsystems vvith C have shovvn that the formation of a number of very stable compounds is possible. In addition to u phase (Co,Fe)7W2(Co,Fe,W)4, the formation of cementite phases M3 ([Co,Fe,W]3C), M6C and M12C that have large negative Gibbs' energies is also possible. Energetically the most stable phase is M12C, but a spot analysis (WDX) of the Fe/Co-containing com- pound phase at the interface cemented carbide/-braze shovved4 that at shorter brazing times espe-cially phase M3C is formed. The Gibbs' energy of this cementite phase is more than tvvice smaller than Gibbs' energy required for the formation of WC Therefore it can be concluded that WC from cemented carbide particles during brazing is decomposed and released atoms of C and W diffuse across the Co matrix to the interface cemented carbide/braze. C and Co reach Fe, vvhich is present at interface. faster than W. Namely, the diffusion coefficient of C in Fe is very large because of its interstitial solubility and Co is the base of the solid solution of cemented carbide matrix. The diffusion of W is the slovvest process, therefore it can be concluded that this process controis the formation of a compound layer and for very short brazing times energetically the most favourable process is the formation of a stable W-poor cementite phase. Therefore in the first stage of brazing, C and Fe can diffuse from the steel across the braze to the interface braze/cemented carbide, because of their large con-centration gradients. By the formation of cementite phase, the chemical potential of C is decreased and the conditions for WC decomposition, the diffusion of C and W to the interface cemented carbide/braze, as vvell as for the continuing grovvth and development of the compound layer are fulfilled. Some authors4 ex-plained the inferior strength of the vacuum brazed joint by the formation of a thin compound layer during brazing and appearance of microporosity behind this layer. The void formation during brazing vvith pure Cu, could be explained by the faster diffusion of Co from the cemented carbide matrix in comparison vvith the diffusion of other possible elements in the cemented st<>,-! braze Figure 5: Elements distribution at the cemented carbide particle/structural steel joint of the prepared sample Slika 5: Porazdelitev elementov v vzorcu izdelane brusne plošče na spoju zrno karbidne trdi ne/konstrukcijsko jeklo cemented carbide cemented carbide * - carbide matrix (Kirkendall's effect). The metallo-graphic examination of our samples (brazing joints) shovvs that microporosity is not present. On the basis of this, it can be concluded that the diffusion of Co is suppressed or, vvhich is more probable, that Fe and especially Cu and Ni occupy the Co emptied sites. Namely, Ni, Fe and Co can form the system of a com-pletely solid solution and therefore these elements can substitute for the Co in cemented carbide matrix and, as it vvas mentioned before, the micro-probe analysis of brazed samples really shovved a region enriched by the Fe, Cu and Ni behind the formed compound layer. At the interface steel/braze, the formation of a sim-ilar compound zone as observed at the interface braze/cemented carbide is not evident. The border zone steel/braze is enriched by Cu, Ni and Co (see Fig. 5). Co and Ni form a system of completely solid solubility vvith Fe, and solubility of Cu in y Fe at brazing temperatures is = 7.5%. Therefore, from this point of view and from the morphology of the border zone (see Fig. 4) it can be concluded that an intercrys-talline diffusion of Cu is occurred, the solution of the mentioned elements in Fe is carried out and a homo-geneous solid solution in the system Fe-Co-Ni-Cu at the interface steel/braze is therefore formed. Fig. 6 presents schematically a brazing joint vvith estab-lished diffusion directions of individual elements and formed border zones at the interface steel/braze and braze cemented carbide particles. The grinding experiments vvith prepared grinding vvheels from the soft low-carbon structural steel and with Co enriched layer, Ni > Fe structural braze steel 0,2 % C ' (Cu - 4% Ni) (ferrite * 15% pearlite) compound zone (Me j:) enriched wit.h Co and Fe and Ni - Cu+Ni i islands of Co*re I ° 1 : o o o intercrystalline diffusion of Cu Cu - r i rt>, Co - poor layer Figure 6: Schematic presentation of the brazing joint and the diffusion flows of the individual elements vvith generated border zones (sample of grinding vvheel prepared by brazing of structural steel and cemented carbide particles) Slika 6: Shematični prikaz vakuumskega spoja zrno karbidne trdine/konstrukcijsko jeklo z difuzijo posameznih elementov in nastalima mejnima conama cemented carbide particles vvith Cu-Ni brazing agent shovved that during the brazing a sufficiently strong diffusion bonded joint is formed. The average hardness of the steel base after brazing is 140 HN^, of cemented carbide particles 1475 HV, and of the braze 100 HV01. 3.2 Characteristics of cemented carbide particies/low-alioy hardening steel brazing joint Now, the analysis of brazing experiments vvith the second selected steel vvill be presented. Here, in comparison vvith the first steel, a higher carbon con-tent and a higher content of alloying elements (Cr, Mn, etc.) has to be considered. Fig. 7 shovvs the cross seetion steel/braze/cemented carbide particle vvith element distribution obtained vvith EPMA. It can be clearly seen, that Fe diffuses to the interfaces braze/cemented carbide particles. W remains in cemented carbide particles. Therefore it can be concluded that the diffusion of W is negligible or more precisely that it is limited by the diffusion to the interface cemented carbide/braze. Ni is mostly in the braze. With Ni enriched regions are also interfaces steel/braze and braze/cemented carbide. An inereased amount of Co at the interface cemented carbide/braze as well as at the interface braze/steel is noticed. On the basis of that it can be concluded that Co diffuses from the matrix of cemented carbide particles to the interface cemented carbide/braze and aeross the braze to the interface braze/steel. The presence of Mn at the interface braze/cemented carbide is evident. Because of the presence of Cr in steel, the diffusion of Cr also occurs. The diffusion of Si is insignificant. cemented carbide particle Figure 7: Elements distribution at the cemented carbide particle/low-alloy hardening steel joint of the produced grinding vvheel sample Slika 7: Porazdelitev elementov na spoju zrno karbidne trdine/jeklo za poboljšanje izdelanih brusnih plošč On the basis of previous statements, element distribution obtained by EPMA and metallographic ex-amination (see Fig. 7 and 8) it can be concluded that at the interface cemented carbide/braze, a thin layer (= 20 h- 30 um), probably of an intermetallic compound enriched vvith Co and Fe, or vvhat is more probable, concerning the diffusion of C, of the cementite phase MexC vvas formed. The thickness of the formed layer depends on the thickness of the braze and cemented carbide particles distance from the steel base, re-spectively, as vvell as on the brazing temperature and the soaking tirne at the brazing temperature. In the layer of the braze, at certain places (islands), an increased content of Co and Fe can also be no-ticed. Because of intensive diffusion of Co out of cemented carbide particles, a diminished Co content behind the formed compound layer is observed. But at the same places, an increased content of Cu, Fe and Ni is noticed. Therefore it can be concluded that the diffusion of Ni, Fe and Cu proceeded in opposite direction. Besides Fe, Co, Ni and W (probably also C), Cr and Mn, vvere also found at the interface braze/cemented carbide. Therefore it can be stated that during the brazing of steel base and cemented . braze ", iM cemented carbide lOOum i r l carbide particles vvith Cu-Ni brazing agent (alloy-pow-der), a really very complex compound vvas formed. As it vvas mentioned above, some authors4 have es-tablished that brazing joint betvveen the steel plate and cemented carbide in this region is the vveakest, because of the brittle nature of the compound layer formed during the brazing. At the interface steel/braze seems to form no simi-lar compound layer (as it vvas observed at the interface braze/cemented carbide). Hovvever, the notice-able increase of Co content is also observed, as vvell as the absence of W and the diminishment of Fe content because of its diffusion tovvards the cemented carbide. The border zone steel/braze is also enriched vvith Cu and Ni (see Fig. 7). Therefore, from that and from the morphology of the border zone (see Fig. 8 a) it can be concluded that as the intercrystalline diffusion of Cu proceeds, the solution of the mentioned elements in steel (Fe) is carried out and a homoge-neous solid solution forms in the system Fe-Co-Ni-Cu at the interface steel/braze. »ith Co, Ni t Fe enriched layer luw allny hardening / braze steel y (Cu- 2% Ni) (0,5% C, 1% Cr, 0,2% Mo) / (pearlite+20% ferrite) compound zone (MexC) enriched with Fe, Co, Ni, Cr and Mn cemented carbide particel +4 •i L---Cu.M ' i',.{tinds of Co+Fe I ° 3 | Cr*Hi ^ J_ / intercrystalline diffusion of Cu (Cu-rich, Fe-poor layer) Co - poor, Fe, Cu, Ni - rich layer cemented carbide 100 p m || Figure 8: a) Microstructure of cemented carbide particle/low-alloy hardening steel joint of the grinding vvheels produced at IMT Ljubljana and b) noticed damage (hair-shaped crack) of the cemented carbide particle Slika 8: a) Mikrostruktura vakuumskega spoja zrno karbidne trdine/jeklo za poboljšanje vzorca brusne plošče izdelane na IMT Ljubljana in b) opazne poškodbe (lasne razpoke) karbidnih zrn Figure 9: Schematic presentation of brazing joint and the diffusion flovvs of the individual elements vvith generated border zones (sample of grinding vvheel prepared by brazing of low-alloy hardening steel and cemented carbide particles) Slika 9: Shematični prikaz vakuumskega spoja karbidna trdina/jeklo za poboljšanje vzorca na IMT izdelane brusne plošče z difuzijo posameznih elementov in nastalima mejnima conama Islands enriched vvith Fe and Co in the braze layer are observed, especially at some places, vvhere the cemented carbide particles are very closely located to the steel base. Therefore it can be concluded that the most intensive diffusion of these tvvo elements during brazing (in comparison vvith other active elements) occurs if the diffusion of interstitial C, vvhich could not be analysed by our EPMA, is neglected. Metallographic examinations also shovved that some of the cemented carbide particles used vvere damaged (hair-shaped cracks). After brazing, around these cracks, a thin layer of cemented carbide vvith di- minished Co content is observed. It can be conclud-ed that during the brazing these damaged regions represent an active part for the diffusion of Co from cemented carbide to the interface cemented car-bide/braze and across the braze to the interface braze/steel (see Fig. 8 b). Fig. 9 presents schemati-cally the brazing joint vvith established diffusion direc-tions of individual elements and the formed border zones at the interface steel/braze and braze/cement-ed carbide particles. The grinding experiments vvith prepared grinding vvheels of low-alloy hardening steel and cemented carbide particles vvith Cu-2%Ni brazing agent shovved that during the brazing a strong enough diffusion bonded joint is formed. The microstructure is ferritic-pearlitic vvith the prevailing content of pearlite and rare ferrite grains. The average hardness of the steel base after brazing is 200 HV,, of cemented carbide particles 1980 HV, and of the braze 110 HV01. 3.3 A comparison of both brazing joints For the first brazing couple (structural steel/ce-mented carbide particles) only a marked Co and Fe rich compound layer at the interface cemented car-bide/braze vvas observed. At the second brazing couple (lovv alloy hardening steel/cemented carbide particles), hovvever, the marked enrichment vvith Co at both formed interfaces is observed. Here, the marked Co diminishment in the cemented carbide particle re-gion close to the compound layer formed during brazing is also observed. Because of Cr presence in steel, the diffusion of Cr to the interface braze/cemented carbide occurred. The diffusion of Mn for the first brazing couple is insignificant, but it is evident for the second brazing couple. Therefore, on the basis of our investigations, it can be concluded that the low-alloy lovv-carbon structural steel for the brazing of cemented carbide particles vvith the steel base is a better op-tion. The most important parameters of brazing are the brazing temperature (including heating and cooling rate) and the soaking tirne at this temperature. Too high temperature and too long holding tirne of the brazing couple at this temperature enable the forma-tion of a thicker layer of a hard and brittle compound phase at the interface braze/cemented carbide. It also causes the diminishment of the Co content at the cemented carbide border zone as vvell as formation of voids because of the Kirkendall effect and the resulting vveakness of the brazing joint. The decompo-sition of WC because of C and W diffusion at interface cemented carbide/braze is also possible. Therefore lovver temperatures (close to the melting point of the braze) and shorter soaking times at this temperature are more suitable for both selected steel bases. Hovvever, the brazing conditions have to be selected in such a manner that sufficient strength of the diffusion bonded joint is already formed. 4. Conclusions The usability of Cu-2%Ni vvater atomized povvder for the brazing of cemented carbide and steel base vvas investigated. On the basis of our investigations it can be concluded that vvater atomized Cu-2%Ni alloy-povvder is suitable brazing agent. A vacuum brazing procedure for the manufacturing of grinding vvheels containing cemented carbide particles vvas devel-oped. The investigations shovv that for these types of grinding vvheels, soft structural steel as a metal base seems to be the most convenient solution. The aim of future experiments and investigations is to optimise the brazing procedure as vvell as to devel-op nevv shapes of grinding vvheels. Future investigations should shovv vvhich combination of selected ma-terials (steel/braze/VVC-C) will give a joint vvith the highest strength. We have in mind brazing experi-ments vvith different steels (structural steels, tool steels, etc.), brazes (pure Cu, Cu-Ni vvith different Ni contents, etc.) and different sorts of cemented carbide particles. By metallographic examinations, analysis vvith EP-MA and hardness measurements the brazing joints betvveen structural as vvell as low-alloy hardening steel and cemented carbide particles using Cu-2%Ni brazing agent vvere partly evaluated. The diffusion joint is sufficiently strong and vvithout noticeable defects. Its microstructural feature is a thin layer of in-termetallic compound from the system Fe-Co(-W-Cr-Cu-Ni) or, more probably cementite phase MXC type at the interface braze/cemented carbide formed during the brazing and intercrystalline diffusion of Cu at the interface braze/steel. Hovvever, only more de-tailed and systematical investigations vvith analysers (SEM/EDX, WDX, EAS) sensible for the light elements (O, C) distribution could fully explain the oc-currence of multi phase diffusion and take into ac-count ali thermodynamic and kinetic aspects in order to ansvver vvhat kind of compound is formed at the interface steel/braze and especially at the interface braze/cemented carbide during the brazing. References 1 D. Kmetic et al.: High temperature vacuum brazing vvith the simultaneous vacuum heat treatment, Annual reports of IMT Ljubljana (slovenian language), IMT Report No.: 89-064, November 1991 2 D. Kmetič et al.: Vacuum bonding of tool and structural steels in vacuum heat treatment furnace, Annual reports of IMT Ljubljana (slovenian language), IMT Report No.: Ml 87-032, November 1988 3 E. Claire et al.: Part 3: Production of Metal Povvders and Part 7: Povvder Systems and Applications-Subsections: Cemented carbides and Metal Povvders Used for Brazing and Soldering, Metal Handbook, 9th edition, Volume 7, Povvder Metallurgy 4 M. Stock, K. Hack: Thermochemical Aspects of Multiphase Diffusion during Brazing of Hard Metal, Z. Metallkd. 84, 1993, 11, 759-766 5 Commercial brochure: DEGUSSA (Lote, Lotpasten, Flussmittel), MH 58-4-889 B, Hanau, Germany 6 Commercial catalogue: MAHLER, Brazing in Protective-gas and Vacuum Furnaces, Esslingen, Germany 7 B. Šuštaršič, B. Breskvar, V. Leskovšek, A. Rodič: Microstructural characteristics of vvater atomized Cu-based povvders for brazing, Proceedings of 30th Symposium on Devices and Materials SD'94, Zreče, Slovenia, 239-244 8 J.J. Dunkley: The Production of Metal Povvders by VVater Atomization, Povvder Metaiiurgy International, Vol. 10, No. 1/78 9 J.J. Dunkley, J.D. Palmer: Factors Affecting Particle Size of Atomized Metal Povvders, Povvder Metallurgy International, 1986, Vol. 29, No. 4 10 MPIF: Standard Test Methods for Metal Povvders and Povvder Metallurgy Products, Metal Povvder Industries Federation, Edition 1985/1986, Princeton, Nevv Yersey 11 J.J. Dunkley: The Factors Determining the Oxygen Content of VVater Atomized 304L Stainless Steel Povvder. Reprint of Paper Presented at the National PM Conference, Philadelphia, U.S.A., May 1981, Davy-Loewy R&D Centre Pulsed Plasma Nitriding of Stainless Steel Nitriranje nerjavnega jekla v pulzirajoči plazmi Torkar M.,1 V. Leskovšek, IMT Ljubljana B. Rjazancev, Department of Orthopaedics, Hospital Jesenice A use could be found for pulsed-plasma nitrlded AISI 316L stainless steel in a wlde range of biomedlcal appllcatlons, e.g. femoral and biartlcular heads of joint prostheses. Forged samples of steel vvere pulsed-plasma nitrided at a temperature of 540°C for 24 hours to increase the hardness of the surface. During nitriding, the hardness of the 70 pm thick layer uniformly increased to 958 HV 0.1. The hardness of the steel belovv the nitride layer remained unehanged, 210 HV 0.1. The thickness of the layer depended on the process parameters. The research shovved that nitriding of stainless steel for medical implants in pulsed plasma is feasible. Key vvords: stainless steel, pulsed plasma nitriding, biomedical applications Nerjavno jeklo AISI 316L, nitrirano v pulzirajoči plazmi, bi bilo mogoče uporabiti za biomedicinske namene, npr. za femoralne in biartikularne glave kolčnih vsadkov. Kovani vzorci jekla so bili nitrirani v pulzirajoči plazmi 24 ur na temperaturi 540°C. Med nitriranjem je trdota 70 pm debele plasti enakomerno narasla do 958 HV 0.1. Trdota jekla pod nitrirano plastjo je ostala nespremenjena, 210 HV 0.1. Debelina nitrirane plasti je odvisna od parametrov procesa. Raziskava je pokazala, da je nitriranje medicinskih implantatov v pulzirajoči plazmi izvedljivo. Ključne besede: nerjavno jeklo, nitriranje v pulzirajoči plazmi, uporaba v biomedicini 1. Introduction An increase of the wear resistance and surface strength of different steels and alloys vvith nitride forming elements can be obtained by pulsed plasma nitriding'. Usually, also the corrosion resistance of the surface is increased, vvhile the corrosion resistance of stainless steel is sometimes slightly dimin-ished. Modem nitriding devices and technology, hovvever, maintain or even increase the corrosion resistance2. Nitriding in pulsed plasma is beneficial due to the low temperature (betvveen 350 C and 660" C) of the process and the possibility of influencing the compo-sition of the nitrided layer (y, e, y+e, and diffusion layer)3. The low temperatures of the process maintain the mechanical properties of the material belovv the layer unehanged. The process is ecological friendly and nontoxic. Gases in normal state are nonconductive for elec-trical current. This property is changed by high volt- Dr. Matjaž TORKAR IMT. Lepi pot 11 61000 Ljubljana age or at low pressure when lightning or glovv dis-charging appears, respectively. In both cases, the nonconductive gas is transformed to an ionized plasma vvith a sufficient electrical conductivity. Nitriding in pulsed plasma is based on glovv diseharging pulsed current in a lovv pressure chamber. Electrons are released on the cathodic surface of the sample, sputtering off the surface atoms, and nitrogen ions migrate into the specimen. At a distance of some millimeters above the cathodic surface of the specimen, the ions accelerate and hit the surface vvith high kinetic energy. About 90 % of this energy is transformed to heat, vvhich vvarms the surface up to the nitriding temperature. The heat is controlled by electric povver, and no additional heating is required. Nitrogen ions are highly reactive in the plasma, and iron nitrides start to form on the sputtered surface. Because of the lovv temperature, FeN mole-cules on the surface of the tool or sample decom-pose into lovver nitrides. At the decomposition FeN-^Fe2N^Fe3N^Fe4N, nitrogen is released. A part of the released nitrogen diffuses into the sample and the rest is returned into the plasma. The process Electron ----1> FeN^j \^Adsorption —>n ^>e-Phase —»N — »N "y'-Phase a - Phase © < S Z cr £ N Figure 1: Schematic presentation of the ion nitriding process Slika 1: Shematski prikaz postopka ionskega nitriranja of ion nitriding is shovvn schematically in Fig. 1. In principle, ali materials based on iron can be nitrided, since vvith glow discharging, ions of nitrogen are ac-tive enough to recombine on the surface. In some minutes of treatment, the nitride layer is formed and the steep gradient of concentration accelerates the diffusion of nitrogen into the specimen4. The objective of the present research vvas to check the nitriding of AISI 316L stainless steel in pulsed plasma. 2. Experimental work and results 2.1. Experimental procedure The samples for nitriding vvere machined from the AISI 316L stainless steel bars vvith diameter of 30 mm and the composition of the steel vvas: 0.049 % C, 17.9 % Cr, 13.1 % Ni, 2.5 % Mo, 0.03 % Al and vvere pulsed-plasma nitrided at 540 C for 24 h. After nitriding, the samples vvere prepared for microstructure examination and hardness tests. The hardness vvas measured vvith a Vickers indenter, us-ing a 100 g load and a 10 s load duration. The microstructure vvas examined, after etching vvith Marble's reagent, by optical microscopy, vvhile the fracture surfaces of the nitride layer vvere examined by scanning eleetron microscopy (SEM). 2.2. Experimental results In Fig. 2 a and b, the optical micrographs of unetehed and etehed nitrided layers are shovvn. The nitrided layer is light-gray in an unetehed condition, and only inclusions are found by optical or scanning microscopy. After etching in Marble's reagent, the nitride layer is darker and it looks homogeneous. The optical micrographs (Fig. 3 a and b) of the transverse seetion through the base steel and nitride layer shovv Vickers imprints and a needle serateh, both of vvhich confirm the increased hardness of the nitride layer. The needle serateh, easily visible in the base steel, disappeares vvhen crossing the harder nitride layer. The increased hardness of the nitride layer is connected to the reduetion of its fracture toughness and the appearance of a brittle fracture, vvhereas the fracture of the steel belovv the layer remains duetile. In Fig. 4 a the cracks in the nitride layer, formed at bending the nitrided surface to an angle 180 are shovvn. The fracture is brittle and the crack stops in the interface of the nitride layer-steel, because the base material has a higher fracture toughness. The topology and morphology of the nitrided surface vvith a bending crack is shovvn in Fig. 4 b. The nitride surface can be polished to a high brilliancy, vvhich is important for medical use. as i.e. for femoral and biarticular heads of joint pros-theses. The modification of the surface morphology and topology may be beneficial for an improved biologi-cal performance or improved bone-bonding5 at the uncemented modular stem or cement-bonding at the cemented stem. The base steel shovvs after nitriding a duetile trans-granular fracture (Fig. 5). Hardness measurements shovv an average and homogeneous hardness of 958 HV 0.1 in the layer (Fig. 3 a) and a hardness of 210 HV 0.1 in the base material. The usual hardness transition zone. found in nitrided steels, vvas not found. The results confirm that the process of ion nitriding is suitable for inereasing of the relatively soft AISI 316L stainless steel surface. It is regarded as promising for inereasing the vvear resistance of such materials. 3. Conclusions Ion nitriding of AISI 316L steel in pulsed plasma in-creases the surface hardness from 210 HV 0.1 to 958 HV 0.1, vvhile the hardness of the base alloy is unehanged. (a) ' ' 100/im (b) - - . 1 OO^ni Figure 2 a and b: Optical micrograph of the nitrided layer after nitriding in pulsed plasma for 24 hours at 540 C, (a) - unetched, (b) - etched vvith Marble's reagent Slika 2 a in b: Optični mikroposnetek nitriranega sloja po nitriranju v pulzirajoči plazmi, 24 ur na temperaturi 540 C, (a) - nejedkano, (b) - jedkano v Marble jedkalu Figure 3 a and b: Optical micrograph of the nitride layer, (a)- Vickers indentions and (b)- a needle scratch Slika 3 a in b: Optični mikroposnetek nitriranega sloja, (a)- odtisi merjenja trdote po Vickersu in (b)- raza preko nitriranega sloja, napravljena z iglo The thickness of the layer after 24 hours of nitriding at 540 C is up to 70 pm. The propagation of the crack opened in the nitride layer stopps at the interface nitrided layer-base steel due to higher ductility of the base steel. This indi-cates to a difference in fracture toughness betvveen the nitrided layer and the matrix. The modification of the surface morphology and topology may be beneficial for an improved surface bonding vvith bone or cement. Acknovvledgement The authors are grateful to the Ministry of Science and Technology, Slovenia for supporting this vvork. References 1 H. Hornberg; Glimm-Nitrieren: ein Verfahren zum Nitrieren von Stahloberflachen mit Hilfe einer Glimmentladung, Harterei Tech. Mit., 17, 1962, 2, 82. Figure 4 a and b: SEM micrographs. Nitrided surface vvith bending cracks (a)- bending angle 180 and (b)- detail of the surface Slika 4 a in b: SEM posnetek nitrirane površine z razpokami, nastalimi pri upogibu, (a)- kot upogiba 180 in (b)- detajl površine Figure 5: SEM micrograph. Ductile fracture of the base material Slika 5: SEM posnetek žilavega preloma jekla pod nitriranim slojem 2 R. Chaterjee-Fischer, VVarmebehandlung von Eisen-vverkstoffen, Expert Verlag, Sindelfingen, 1986, 125 3 M. Hempel, A. Kochendorfer and E. Hillnhagen: Gefiigeveranderungen in u- eisen durch lonen-bestrahlung, Arch. Eisenhuttenvv., 33, 1962, 6, 504 4 H. Knupell, K. Brotzmann and F. Eberhard, Nitrieren von Stahl in der Glimmentladung, Stahl u. Eisen, 78, 1958. 26, 1871 5 C. P. A. T. Klein, J. G. C. VVolke, R. C. Vriesde, J. M. A. De Blieck-Hoger Vorst: Cortical bone ingrovvth in grooved implants vvith calcium phosphate coatings: a gap model study, J. Mater. Sci. Mater. Med., 5, 1994, 9&10, 569-574 Hydrogen and Temper Embrittlement of Medium Strength Steel Vodikova in popustna krhkost jekla srednje trdnosti Ule B.,1 V. Leskovšek, Institute of Metals and Technology, Ljubljana The fracture ductility of high strength steel is strongly influenced by the presence of hydrogen, although hydrogen does not significantly affect the yieid strength. The deterioration of fracture ductility is particularly evident in low strain rate tension tests, but less pronounced at conventional crosshead speeds. Microfractographic investigations of fracture surfaces of hydrogen charged high strength 5Cr-1 Mo-0.3V steel from low strain rate tension test indicate that the grovvth and the coalescence of voids in the final stages of the fracture process are partly assisted by the decohesion of interfaces on vvhich hydrogen is adsorbed. Hovvever, such phenomena are not observed in the experimental medium strength steel. Although a strong interaction betvveen hydrogen and temper embrittlement vvas frequently observed in the alloyed steels'' and though the magnitude of such effect vvas directly related to the degree of intergranular phosphorus enrichmenf, such synergy vvas not found in the experimental steel vvith post-martensitic microstructure. Key vvords: high strength steel, medium strength steel. hydrogen and temper embrittlement. fracture ductility, lovv strain rate tension test Na lomno duktilnost jekla z visoko trdnostjo močno vpliva v jeklu prisoten vodik, čeprav slednji nima znatnejšega učinka na napetost tečenja jekla. Poslabšanje lomne duktilnosti je zlasti očitno pri majhni hitrosti natezanja, medtem, ko je pri običajni hitrosti natezanja manj izrazito. Mikrofraktografske preiskave prelomnih površin nastalih pri počasnem natezanju vodičenega jekla 5Cr-1 Mo-0.3V z visoko trdnostjo kažejo, da je rast in zlivanje por v končnih fazah procesa loma deloma podprta z dekohezijo medplastij, na katerih je adsorbiran vodik. Tega pojava nismo opazili pri preiskovanem jeklu srednje trdnosti, čeprav je bila pri legiranih jeklih često opažena močna interakcija med vodikovo in popustno krhkostjo12 in čeprav je bila intenzivnost tega učinkovanja neposredno povezana s stopnjo interkristalne obogatitve s fosforjem,2 pa takšne sinergije nismo našli pri preiskovanem jeklu s postmartenzitno mikrostrukturo. Ključne besede: jeklo visoke trdnosti, jeklo srednje trdnosti, vodikova in popustna krhkost, lomna duktilnost, upočasnjeno natezanje 1. Introduction The adverse effect of hydrogen on mechanical properties has long been recognised in various metallic materials, especially in high strength steels. Hydrogen embrittlement of such steels often in-volves both a change in fracture mode and a reduc-tion in ductility compared vvith the unhydrogenated condition3. It has also been established that the ki-netics of hydrogen embrittlement depends on the strain rate4"6. Hovvever, at constant static load, the Dr. Boris ULE IMT, Lepi pot 11 61000 Ljubljana failure of high strength hydrogen charged steels, knovvn as delayed failure, frequently occurs. This is caused by stress induced segregation of hydrogen and is characterised by the nucleation of a microc-rack vvhich then grovvs until it reaches a critical size, resulting in an abrupt failure. The incubation period, and the tirne to failure, are prolonged vvith a decrease in load until, at a sufficiently lovv load, delayed failure does not occur. Therefore, a threshold stress inten-sity factor Kth can be introduced vvhich is consider-ably lovver than the critical stress intensity factor or fracture toughness K,c of the uncharged steel. Our preliminary investigations confirm that lovv concentrations of hydrogen (< 1 ppm by vveight) in high strength steel have no substantial influence on its fracture toughness measured at conventional strain rate (1 mm min 1). Hovvever, the lovv strain rate (0.1 mm min 1) of hydrogen charged high strength steel decreases the fracture ductility, vvhich indicates the existence of the threshold stress intensity factor at semi-static testing conditions6. Because the interaction betvveen hydrogen and temper embrittlement vvas frequently observed in lovv and medium alloyed steels12, the synergism betvveen both embrittlements vvas investigated in a 5Cr-1 Mo-0.3V steel vvith post-martensitic microstructure. The aim of the present vvork is therefore not only to demonstrate the applicability of lovv strain rate tension test in order to study the hydrogen embrittlement phenomena in high strength steel, but also to study the possible interaction betvveen hydro-gen and temper embrittlement in medium strength steel. 2. Experimental A secondary hardening steel containing 0.38 C, 0.99 Si, 0.38 Mn, 0.012 P, 0.01 S, 5.19 Cr, 1.17 Mo, and 0.23 V (ali wt-%) vvas used. Cylindrical tensile specimens vvith 10 mm dia. vvere machined from a forged rod after being homogeneously annealed and normalised. Some specimens vvere austenised at 980 C for 30 minutes in a vacuum furnace, quenched in a flovv of gaseous nitrogen at a pressure of 0.5 MPa, and then tempered at temperatures of 620, 640 or 670r C. Three separate classes of yield stress, 1220, 1020 and 900 MPa respectively, vvere achieved. The remaining specimens vvere quenched under the same conditions, then tempered tvvice for 2 hours at 710 C vvith intermediate undercooling (yield stress: 668 - 679 MPa, Charpy V-notch impact ener-gy: 72 J), vvhereas some of these specimens vvere additionally tempered for 24 hours at 570°C, i.e. in the temperature range of reversible temper embrittlement78 (yield stress: 648 - 665 MPa, Charpy V-notch impact energy: 37 J). In such cases, the experimental steel had a post-martensitic microstructure. The cathodic charging of tensile specimens vvas performed for 1 h in 1 N sulphuric acid at a current density of 0.3 mA cm 2. The concentration of hydro-gen in some specimens (bulk concentrations) vvas determined using a high temperature (1050 C) vacuum extraction technique vvith gas chromatograph-ic analysis. Tension tests vvere performed after the hydrogen charging of specimens had been com-pleted and the specimens had been exposed to air for 24 hours. This enabled the concentrations of hy-drogen to approach the residual value of approxi- mately 0.8 ppm by vveight vvhich remained almost tirne independent. The tension tests vvere performed at a conventional strain rate, i.e. at a crosshead speed of 1 mn min and at a lovver strain rate, i.e. at a crosshead speed of 0.1 mm min1. The fracture surfaces of the tensile specimens vvere examined in a scanning electron microscope (SEM). 3. Results 3.1 Tensile Properties The results obtained from different strain rate tension tests for both uncharged and hydrogen-charged steel are pointed out in Table I and II. The results in Table I refer to the specimens vvhich vvere quenched and tempered in a temperature range from 670 to 620° C vvith a class of yield stress of approx. 900 MPa, 1020 MPa and 1220 MPa, respectively. Table I: Mechanical properties of uncharged and hydrogen charged steel, quenched and tempered up to high yield stresses. Tabela I: Mehanske lastnosti nevodičenega in vodičenega jekla, ki je bilo kaljeno in popuščano na visoko napetost tečenja Crosshead speed 1 mm min"1 Crosshead speed 0.1 mm min Yield Uniform Reduction Yield Uniform Reduction stress elongation of area stress elongation of area MPa % % MPa % % Uncharged steel 924 8.7 52 910 8.5 51 1010 7.4 51.3 1027 6.5 50.3 1270 6.4 50 1214 6.2 50.3 Hydrogen-charged steel 885 8.4 50.3 899 8.1 47.7 1082 7.2 49.3 1078 6.5 42.7 1209 6.1 47.3 1226 6.0 27.3 The decrease of the crosshead speed at tension had no influence on the mechanical properties of the uncharged steel, vvhereas it essentially influenced the reduction of area in the hydrogen charged steel (approx. 0.8 ppm hydrogen). The loss of ductility vvas more pronounced in steel vvith a higher yield stress. The results in Table II refer to the steel vvith a post-martensitic microstructure, i.e. specimens vvhich vvere quenched and tempered tvvice at 710 C vvith intermediate undercooling, and specimens vvhich Qa = 160.3 t 4.5 kJ m oT< 11 1J2 1 .„3 y - W3,KJ vvhere R.A^ is the reduction of area at conventional strain rate tensile test, i.e. at a crosshead speed of 1 mm min 1 and R.A.0, is the reduction of area at low strain rate tensile test, i.e. at a crosshead speed of 0.1 mm min"1. Table II: Mechanical properties of uncharged and hydrogen charged steel quenched and tempered tvvice at 710 C and of the same steel, additionally tempered for 24 hours at 570°C Tabela II: Mehanske lastnosti nevodičenega in vodičenega jekla, ki je bilo kaljeno in dvakrat po-puščano pri 710°C ter istega jekla, ki je bilo dodatno popuščano 24 ur pri 570 C Crosshead speed 1 mm min1 Crosshead speed 0.1 mm min1 Figure 1: Evaluation of the activation energy of segregation of phosphorus according to Arrhenius equation (from Ref. 7) Slika 1: Izvrednotenje aktivacijske energije za segregiranje fosforja z uporabo Arrheniusove enačbe (Ref. 7) vvere further, additionally, tempered for 24 hours at 570 C, both vvith a class of yield stress of approx. 660 MPa. A weak reversible temper embrittlement of experimental steel, having a high tempered post-martensitic microstructure, vvas produced vvith addi-tional tempering as confirmed by our preliminary in-vestigations of the time-temperature relationship of the Charpy impact energy reduction resulting from such tempering7. An activation energy of about 160 kJ mol"1, vvhich is very close to that for bulk diffusion of phosphorus in ferrite, vvas derived from the slope of a log-log plot of tirne versus reciprocal tempering temperature (Fig. 1). Indeed, it has already been confirmed, using Auger spectroscopy, that in particular phosphorus segre-gates in this type of steel. Romhanyi and covvorkers9 found up to 6 % P and 1 % S at grain boundaries, and a strong carbon peak vvith carbide structure is also remarkable in Auger spectra of such steel, austeni-tized at 1100 C, quenched and tempered at 600 C (Fig. 2). Hovvever, the segregations in our experi-mental steel, quenched from much lovver temperature, vvas not so intense, and the embrittlement phe-nomena could be detected only by impact testing and not at ali at semi-static tensile tests. The results from both Tables are also shovvn in Diagram (Fig. 3), vvhere the tendency for hydrogen embrittlement is formulated in accordance vvith Morimoto and Ashida10: R.A.-R.A.01 Degree of embrittlement =-x 100% (1) R. A. -j Yield Uniform Reduction stress elongation of area MPa % % Yield Uniform Reduction stress elongation of area MPa % % Ouenched and tempered tvvice at 710 C(Charpy V-notch energy:72 J) Uncharged steel 679 12.4 55 668 11.8 55.7 Hydrogen-charged steel 668 11.2 45 661 10.3 43.1 Ouenched and tempered tvvice at 710 C further, additional-ly, tempered for 24 hours at 570 C (Charpy V-notch energy: 37 J) Uncharged steel 665 11.1 51.7 664 11.7 51.7 Hydrogen-charged steel 648 11.1 52.5 640 10.9 53 dN(E) dE " A 3 703 Fe E (eV) Figure 2: Auger spectrum of the intergranular surface of steel vvith 5 % Cr, austenitized at 1100 C, quenched and tempered at 600 C (from Ref. 9) Slika 2: Augerjev spekter z intergranularne površine jekla s 5% kroma, austenitizirano pri 1000 C, kaljeno in popuščano pri 800 C (Ref. 9) rate test and the fracture surface of hydrogen charged low- and conventional-strain rate test specimens of medium strength, i.e. the fracture surface of hydrogen charged steel, either quenched and tem-pered twice at 710 C or of the same steel addition-ally tempered for 24 hours at 570; C, are totally duetile (Fig. 6). Of course, the fracture surface of Figure 5: Detail from Fig. 4 (SEM) Slika 5: Detail iz slike 4 (SEM) Figure 4: Fracture surface of hydrogen charged low-strain rate tensile specimen vvith yield stress of 1226 MPa (SEM) Slika 4: Prelomna površina vodičenega in počasi natezanega preizkušanca z napetostjo tečenja 1226 MPa (SEM) Figure 6: Fracture surface of hydrogen charged low-strain rate tensile specimen vvith yield stress of 640 MPa (additionally tempered for 24 hours at 570 C) (SEM) Slika 6: Prelomna površina vodičenega in počasi natezanega preizkušanca z napetostjo tečenja 640 MPa (dodatno popuščano 24 ur pri temperaturi 570 C (SEM) temper embrittled* hydrogen charged uneharged Figure 3: Degree of embrittlement of hydrogen charged steel as funetion of yield stress Slika 3: Stopnja krhkosti vodičenega jekla v odvisnosti od napetosti tečenja 3.2 Fractography The micromorphology of the typical fracture surface of hydrogen charged lovv-strain rate tensile specimen vvith yield stress of 1226 MPa, is shovvn in Fig. 4 and 5. It can be concluded that the hydrogen-induced fracture is locally duetile, tearing type of fracture vvith some quasicleavage details on the periph-ery of larger and deeper tunnel-type dimples. The fracture surface of hydrogen charged high strength specimens obtained at conventional strain 600 700 800 900 1000 1100 Yietd stress. rfys (MPa) 1200 1300 0,8(um additionally tempered impact specimens as for instance Charpy V-notch specimens - even uncharged - are of mixed mode. After an additional tempering at 570 C for 24 hours, the crack propagation path changes and sporadic intergranular fracture along preaustenite grain boundaries, and quasicleavage fracture details, and single ductile tearings can be regularly observed7. 4. Discussion Lovv concentration of hydrogen in high strength steel have no significant influence on the mobility of the dislocations in earlier stages of the tensile defor-mation process. Hydrogen has almost no effect on Figure 7: Schematic representation of microvoid formation, grovvth, and coalescence along grain boundaries in vvhich hydrogen is adsorbed (from Ref. 13) Slika 7: Shematski prikaz tvorbe mikropor, njihove rasti in zlivanja vzdolž meja zrn, na katerih je adsorbiran vodik (Ref. 13) the yield stress or on the uniform elongation of steel, and it only affects the reduction of area. Hovvever, the reduction of area decreases only if the strain rate is lovv enough to enable the Cottrell atmosphere of the hydrogen atoms pinned on the dislocations to penetrate deep into the plastic zone of the tensile specimens. Since the size / of the plastic zone of a hydrogen charged specimen is approximately half the diameter of the neck (1=3 mm) at fracture, and the crosshead speed v is 1.6 x 103 mm s 1 (0.1 mm min"1), a value of strain rate e = v/l = 5.3 x 104 s1 is obtained. In earlier literature11 higher e values are quoted for stainless steel. Hovvever, the investigations performed by Nakano et aV2 on hydrogen charged steel vvith yield stress of 500 MPa using lovv strain rate measurements shovv that at sufficient concentration of hydrogen in steel the reduction of area asymptotically approaches the lovver value even at a critical strain rate of e = 104 s1, vvhich is of the same order of magnitude as in the present investigations. Microfractographic examinations shovv that hydrogen charged lovv strain rate tensile specimens exhibited some interfacial separation on the fracture surface. The grovvth and the coalescence of mi-crovoids along the grain boundary, schematically shovvn in Fig. 7, are accelerated by separating inter-nal interfaces vvhere hydrogen is adsorbed1314. Microvoid coalescence and the separation of inter-nal interfaces due to adsorbed hydrogen become op-erative vvhen the triaxial stress state in the narrovv neck of the tensile specimen is formed (Fig. 7, se-quences 3 to 5), resulting in the "condensation" of the last stage of plastic deformation in the lovv strain rate tension testing of high strength hydrogen charged steel. Hovvever, such phenomena are not observed in medium strength steel. Although a strong interac-tion betvveen hydrogen and temper embrittlement vvas frequently observed in such alloyed steels12 and though the magnitude of such effect vvas directly re-lated to the degree of intergranular phosphorus en-richment2, such synergy vvas not found in the exper-imental steel vvith post-martensitic microstructure. In studying the influence of bulk and grain boundary phosphorus content on hydrogen induced cracking in lovv strength steel, Dayal and Grabke15 also found that the effect of phosphorus is related to the bulk content and not to the grain boundary concentration. Obviously, in the čase of the experimental steel vvith post-martensitic microstructure, the influence of hy-drogen decreases and becomes more complicated due to some particular effect of the microstructure. In agreement vvith Charbonnierand Pressouyre16these results shovv that the nearer is the actual microstruc-tural state to the state of the thermodynamical equi-librium, the less susceptible is the steel to hydrogen embrittlement. 5. Conclusions The relevance of the lovv-strain rate tension test to establish the hydrogen embrittlement susceptibility of both high and medium strength steel is demon-strated. The applicable formula (1) for the estimation of such susceptibility10, based on the reduction of area measurements at the lovv and the conventional strain rate tensile test, is also successfully adopted. The synergism betvveen hydrogen and temper embrittlement vvas not found in the experimental, addi-tionally tempered steel vvith post-martensitic microstructure. On the contrary, such steel is less susceptible to the influence of hydrogen, due to some particular effects of the microstructure vvhich is close to the thermodynamical equilibrium. References 1 C. A. Hippsley and N. P. Haworth: Mater. Sci. Tech., 4, 1988, 791-802 2 C. A. Hippsley: Mater. Sci. Tech., 3, 1987, 912-922 31. M. Bernstein and A. W. Thompson: Int. Metali. Rev., 21, 1976, 269 "J. K.Tien, A. W. Thompson, I. M. Bernstein and R. J. Richards: Metali. Trans., 7A, 1976, 821 5 M. Hashimoto and R. Latanision: Metali. Trans., 19A, 1988, 2799 6 B. Ule, F. Vodopivec, L. Vehovar, J. Žvokelj and L. Kosec: Mater. Sel. Tech.. 9, 1993, 1009-1013 7 B. Ule, F. Vodopivec, M. Pristavec and F. Grešovnik: Mater. Sci. Tech.. 6, 1990, 1181-1185 8 J. Janovec, P. Ševc and M. Koutnik: Kov. Zlit. Teh.. 29, 1995, 40 9 K. Romhanyi, Zs. Szasz Csih, G. Gergely and M. Menyhard: Kristali Tech., 15, 1980, 471-477 10 H. Morimoto and Y. Ashida: Transactions ISIJ, 23,1983, B-325 11 M. B. VVhiteman and A. R. Troiano: Corrosion. 21, 1965. 53-56 12 K. Nakano, M. Kanao and T. Aoki: Trans. Nat. Res. Inst. Met. (Jpn), 29, 1987, 34-43 13 H. Cialone and R. J. Asaro: Metali. Trans.. 10A. 1979. 367-375 14 H. Cialone and R. J. Asaro: Metali. Trans.. 12A. 1981. 1373 15 R. K. Dayal and H. J. Grabke: Steel Research, 58.1987, 179-185 16 J. C.Charbonnier and G. M. Pressouyre: Residual hy-drogen in steels, 4th International Conference "Residuals and Trace Elements in Iron and Steel", Portorož, Yugoslavia, October 1985, Proceedings, pp. 81-103 A New Concept of Quality Evaluation of High Energy Electron Beam Used in VVelding Nov način ovrednotenja visoko energijskega curka elektronov v varilni tehniki VVojcicki S.,1 IVT, Warsaw The proposition of a new technological parameter, so called "coefficient of electron beam energy consumption", fot the mathematlcal description of the quality of electron beam used In vvelding technology is presented. Some features of this parameter are shown. Key vvords: electron beam vvelding, coeficient of energy consumption Podan je predlog novega tehnološkega parametra imenovanega "koeficient porabe energije elektronskega curka" za matematični opis kvalitete elektronskega curka, ki se uporablja v varilni tehniki. Prikazane so nekatere značilnosti tega parametra. Ključne besede: varjenje, elektronski curek, poraba energije 1. Introduction The high energy electron beam (EB) (Fig. 1) is usually characterized by perveance and by means of three groups parameters: • geometric: the envelop form, crossover size - 2rc, aperture angle - 2a, diameter of the focus spot -2r ■ • energetical: integral povver, energy density in the planeš of crossover and focused spot, brightness, dimensions of the active zone12; • structural: electron density distribution in different EB cross sections, phase diagrams, emittance. The evaluations of these parameters by measurements or computer calculation are difficult in practi-cal systems. In first čase, a special apparatus is re-quired and moreover, the obtained data are deformed by measurement error. In the other čase, it is impossible to simulate ali factors and phenome-na, like ion focusing, aberrations of electron gun etc., that have a big influence on the parameters of EB. Consequently, the calculated parameters of EB are loaded vvith considerable errors also. The anticipa-tion of the technological effects, that could be obtained by using a particular EB is difficult because of unknovvledge of real parameters of EB. For example, in EB vvelding technology it is knovvn1,24 that the relationship betvveen the depth of vvelds h and parameters of EB is given by the empirical formula: Dr. Sc. S. VVOJCICKI Institute of Vacuum Technology Dluga Street 44/50 00-241 Warsaw, Poland vvhere U - accelerating voltage, I - EB current, v -speed of vvelding, M - constant dependent on the thermal properties of vvelded materials, and a, (3, y -experimental coefficients representing the features of EB. Usualy a = 0,3-1,3, p = 0,7-1,0, y =0,3-1,0. Becasuse the value of a, y coefficients vary in a vvide range there is no other way to present technological posibilities of EB /and their technological quality/ than to examine experimental vvelds and to compare their cross sections - as it is shovvn in Fig. 23 - or dravving up the plots shovving relationship betvveen depth of vvelds and EB parameters and others parameters of vvelding. The typical diagram h = f(v) is shovvn on Fig. 3. 2. Coefficient of electron beam energy consumption The aim of this vvork is the presentation of a new idea of quantitative evaluation of the high energy EB quality. This idea is based on the mathematical description of technological effects that EB causes in the materials. In the čase of an EB vvelding technol-ogy, the most important parameter is the possibility of obtaining deep and narrovv vvelds by applying the smallest possible amount of EB energy during the operation. From this point of vievv, vve propose to introduce for description of the EB, the so called "coefficient (K) of EB energy consumption"defined as: (a) (b) gun o 10 20 30 E 50 60 70 80 90 wb (kW) (kV) •Ub (mm/min) 30 70 1 0 500 100 lens Figure 2: Comparison of vvelds made in JVVf by EB with different coefficient of energy consumption KF9.33: a) accelerated voltage Vb = 70 kV, beam povver Wb = 30 kW, vvelding velocity v = 1,33 mm/s, depth of weld h = 80 mm, KFl = 221 J/mm2; b) Vb = 500 kV, Wb = 10 kW, v = 1,33 mm/s. h = 75 mm. KF = 100 J/mm2 40 [mm] 30 Figure 1: Same of the geometric parameters of EB: crossover diameter - 2rc, aperture angle - 2oc, spot diameter - 2r„ 20 K = - Ul v-h mm' (2) 10 vvhere U - accelerating voltage, I - EB current, v - ve-locity of vvelding, h - depth of vveld. The physical interpretation of thistechnological parameter of EB is very simple. It is the quantity of en-ergy vvhich is indispensable to create a vveld vvith unit depth on unit length of vveld. EB vvith lovver coefficient of energy consumption K are better than those vvhich need more energy to make the vvelds vvith the same depth. For example, coefficient K for first EB (Fig. 2) is equal: ▲ \ \ A i < aluminium stainless steel copper ■ ■ A \ . • \ • ■ A A ■ v • ■ • • 10 20 30 40 [mm/s] 60 Figure 3: The typical relationship betvveen vveld depth h and the velocity of vvelding v for aluminium, stainless steel and copper. Electron beam povver P = 8 kW magnetic active zone electron cross over HIGH AND ULTRA-HIGH VACUUM COMPONENTS 416 PAGES 38 CATEGORIES 10 SECTIONS 4 CURRENCIES 2INDEXES... 1 CATALOGUE IKE9 Head Office Caburn-MDC Limited The Old Dairy, The Street, Glynde, East Sussex BN8 6SJ United Kingdom Berlin Caburn-MDC Ostendstrasse 1 D-12459 Berlin Germany Torino Caburn-MDC (Alberta Rava) Str. Molinetti 41, II Molino 10098 Rivoli, Torino Ita ly Lyon Caburn-MDC S.A.R.L. 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