VSEBINA – CONTENTS PREGLEDNI ^LANKI – REVIEW ARTICLES Nekaj mo`nih postopkov izdelave hidrofobnih in oleofobnih polimernih membram ter njihova uporaba Hydrophobic and olephobic membrane usage and production processes D. Drev, J. Panjan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51 Liquid metal/ceramic interfaces in dental practice and jewellery manufacturing Staljene vmesne povr{ine med kovino in keramiko v zobozdravstvu in pri izdelavi nakita K. T. Rai}, R. Rudolf, A. Todorovi}, D. Stamenkovi}, I. An`el . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59 IZVIRNI ZNANSTVENI ^LANKI – ORIGINAL SCIENTIFIC ARTICLES Influence of MnS inclusions on the corrosion of austenitic stainless steels Vpliv vklju~kov MnS na korozijsko odpornost avstenitnih nerjavnih jekel ^. Donik, I. Paulin, M. Jenko . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67 AES and XPS characterization of titanium hydride powder Preiskave pra{ka titanovega hidrida s spektroskopijo Augerjevih elektronov in rentgensko fotoelektronsko spektroskopijo I. Paulin, D. Mandrino, ^. Donik, M. Jenko. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73 Fracture characteristics of the Cr-V ledeburitic steel Vanadis 6 Prelomne zna~ilnosti ledeburitnega Cr-V-jekla Vanadis 6 P. Jur~i, B. [u{tar{i~, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77 Surface modifications of maraging steels used in the manufacture of moulds and dies Modifikacija povr{ine jekla maraging in uporaba pri izdelavi kokil in utopov F. Cajner, D. Landek, V. Leskov{ek . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85 Two numerical models of the solidification structure of massive ductile cast-iron casting Numeri~na modela strukture strjevanja masivnega duktilnega `elezovega ulitka K. Stransky, J. Dobrovska, F. Kavicka, V. Gontarev, B. Sekanina, J. Stetina . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93 STROKOVNI ^LANKI – PROFESSIONAL ARTICLES Correlation between the corrosion resistance and the hardness scattering of structural metals treated with a pulsed electric current Korelacija med odpornostjo proti koroziji in raztrosom trdote pri konstrukcijskih materialih, obdelanih s pulziranjem elektri~nega toka A. Babutsky, A. Chrysanthou, J. Ioannou, I. Mamuzic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99 18. MEDNARODNA KONFERENCA O MATERIALIH IN TEHNOLOGIJAH, 15. – 17. november, 2010, Portoro`, Slovenija 18th INTERNATIONAL CONFERENCE ON MATERIALS AND TECHNOLOGY, 15–17 November, 2010, Portoro`, Slovenia . . . . . . 103 ISSN 1580-2949 UDK 669+666+678+53 MTAEC9, 44(2)49–102(2010) MATER. TEHNOL. LETNIK VOLUME 44 [TEV. NO. 2 STR. P. 49–102 LJUBLJANA SLOVENIJA MAR.–APR. 2010 D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH MEMBRAM TER NJIHOVA UPORABA HYDROPHOBIC AND OLEPHOBIC MEMBRANE USAGE AND PRODUCTION PROCESSES Darko Drev1, Jo`e Panjan2 1In{titut za vode RS, Hajdrihova 28c, 1000 Ljubljana, Slovenia 2Univerza v Ljubljani, Fakulteta za gradbeni{tvo in geodezijo, Jamova ul. 2, 1000 Ljubljana, Slovenia darko.drev@gmail.com Prejem rokopisa – received: 2009-09-17; sprejem za objavo – accepted for publication: 2009-12-07 Pri nekaterih vrstah uporabe so zelo za`elene hidrofobne in oleofobne lastnosti membran. To daje tak{nim membranam odlo~ilno prednost pred drugimi materiali. Na podlagi prikazanega postopka izdelave PTFE-membran je razvidno, da njihove pore ne morejo biti manj{e od pribl. 0,01µm. To pomeni, da tovrstne membrane ne moremo uporabljati pri zelo finih filtracijah (ultrafiltracija, reverzna osmoza). Poleg politetrafluoretilena imajo zelo izra`ene hidrofobne in oleofobne lastnosti tudi polisiloksani. Tako kot obstajajo omejitve pri izdelavi PTFE-membran, so tudi pri polisiloksanih (silikonih). Polisiloksanskih membran `al ni mogo~e izdelati po klasi~nem postopku inverzije faz. Kot zelo grobi nadomestki za drage PTFE-membrane so se uveljavili tudi razni oleofobni in hiodrofobni premazi na osnovi mehanskih pen in emulzij. V obeh primerih mora biti v osnovno matrico vgrajen ustrezni hidrofobni in oleofobni material (na primer fluorkarbon). Odprtost strukture in velikost por pri tovrstnih materialih {e zdale~ ni primerljiva s PTFE-membranami. V tem ~lanku je opisan nov postopek izdelave asimetri~nih polisiloksanskih membran, ki je osnova za re{itev nastalega problema. Na podlagi tak{nega postopka bi bilo morda mo`no izdelati tudi zelo fine asimetri~ne polisiloksanske membrane, ki bi imele veliko primerjalno prednost pred vsemi drugimi polimernimi materiali. Klju~ne besede: membrana, oleofobno, hidrofobno, filtracija Hydrophobic and oleophobic membrane properties are very desirable for certain filtration process, giving them an advantage over other materials. A polytetrafluoroethylene (PTFE) membrane production process is presented, revealing a minimum membrane pore size of approximately 0.01µm. It follows that such membranes are not suitable for very fine filtration uses (ultrafiltration, reverse osmosis). In addition to PTFE, polysiloxanes are also marked by enhanced hydrophobic and oleophobic properties. However, polysiloxane membranes suffer from severe production process limitations as their production has thus far not been possible via classical phase inversion processes. Thus, several oleophobic and hydrophobic coatings based on mechanic foams and emulsions have been introduced as a crude replacement for expensive PTFE membranes. These require the introduction of a suitable hydrophobic or oleophobic material such as fluorocarbon into the polymer matrix. However, the degree of the resulting membrane opennes and pore size are not comparable to those of PTFE membranes. In this paper, I present a new production process for asymetric polysiloxane membranes that provides the basis of a solution to this problem. This process could allow the creation of very fine asymetric polysiloxane membranes that would have a large comparative advantage to polymers existing today. Keywords: membrane, oleofobic, hydrophobic, filtration 1 UVOD C. E. Reid in E. J.Breton sta leta 1959 objavila ~lanek o izdelavi in uporabi polimernih membran po postopku inverzije faz 13. V njem sta objavila odkritje, da se lahko uspe{no izlo~ajo elektroliti iz vodnih raztopin z mem- brano iz acetata celuloze. Po tem postopku se tudi danes izdelujejo zelo fine membrane za ultrafiltracijo, nano- filtracijo in reverzno osmozo. Tehnologija izdelave PTFE–membran, ki so na trgu poznane kot membrane Gore-tex, izvira iz patenta W. L. Gore and Associates, Inc., Neweark, iz leta 1971 5. Pozneje so se na trgu pojavile nekatere podobne PTFE–membrane ter razni nadomestki na osnovi suspenzij, mehanskih pen, foto polimerizacije, radioaktivnega obsevanja folij, itd. Vendar pa do sedaj z nobenim drugim postopkom in materialom niso uspeli izdelati membran s podobnimi karakteristikami, kot jih imajo PTFE–membrane. V ~lanku obravnavamo konkretne preizkuse izdelave hidrofobnih in oleofobnih membran, ki smo jih izvedli na podlagi razpolo`ljivih podatkov (patenti, literatura, podatki proizvajalcev surovin, podatki proizvajalcev teh- nolo{ke opreme, izdelki na trgu, itd.) in lastnih zamisli. K inovativnim prispevkom pri tem lahko pri{tejemo nov na~in ka{iranja PTFE–membran in folij na tekstilno podlogo 8 ter postopek izdelave asimetri~ne mikroporoz- ne polisiloksanske membrane 9. Ker postopka nista bila za{~itena kot EU– ali US–patenta, se lahko prosto upo- rabljata zunaj Slovenije. Predpostavljamo, da pri lamini- ranju PTFE–membran uporabljajo podobno tehnologijo, kot je opisana v tem ~lanku. Asimetri~nih polisiloksan- skih membran pa {e ni na trgu, zato sklepamo, da teh raziskav {e ni nih~e implementiral v praksi. Polimerne hidrofobne in ofeofobne membrane so najpomembnej{e za mikrofiltracijo. Razlog za to je velika odprtost strukture ter hidrofobne in oleofobne lastnosti. V zadnjem desetletju so postale nepogre{ljive v Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 51 UDK 541.64.057:66 ISSN 1580-2949 Pregledni ~lanek/Review article MTAEC9, 44(2)51(2010) obla~ilni in obutveni industriji. Veliko se uporabljajo tudi pri industrijskem odpra{evanju, teko~inski filtraciji, biolo{kih ~istilnih napravah itd. 1,2,14. Zaradi relativno velikih por pa niso primerne za fino filtracijo (ultrafiltra- cija, nanofiltracija, reverzna osmoza). Pri izbiri materialov za izdelavo polimernih membran je pomembno vedeti, kak{no vrsto membrane `elimo izdelati, ter za kaj jo bomo uporabljali. Iz PTFE ne moremo izdelati membrane po postopku inverzije faz, temve~ po postopku biaksialnega raztegovanja folije. Tako ne dobimo zelo finih por, kot jih lahko na primer pri polivinilacetatni ali polisulfonski membrani. Vendar pa imajo PTFE-membrane zelo veliko odprtost povr{ine ter hidrofobne in oleofobne lastnosti. Polimerne mem- brane so v ve~ini primerov premalo mehansko stabilne, zato potrebujejo tudi ustrezen nosilni material 1,2,11. PETF–membrane se najve~krat laminirajo na razli~ne tekstilne podloge 8. Posamezne polimerne folije imajo razli~ne prepust- nosti za razli~ne pline, kar je prav tako pomembna lastnost pri izbiri materiala za membrano 3,4,10,12. V tabelah 1 in 2 so prikazane prepustnosti polimernih folij za kisik, ogljikov dioksid in vodno paro. 2 EKSPERIMENTALNI DEL 2.1 Mikroporozni sloji Pri membranskih filtrih z mikroporoznimi sloji se formira mikroporozna plast neposredno na tekstilno podlago. Na tekstilni material se nana{a mikroporozna polimerna mehanska pena, ki se mora po nanosu ustrezno utrditi. Princip formiranja mehanske pene je relativno enostaven 6,7: na specialnem stroju se pod tlakom ume{ava zrak v polimerno disperzijo. Nastane mehanska pena, podobna smetani za kavo. Kakovost in obstojnost mehanske pene je odvisna od gostote D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... 52 Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 Tabela 2: Prepustnost razli~nih polimernih folij za H2O, O2 in CO2 pri 25 µm, 90 % RV in 38 °C (3) Table 2: Permeability of different plastic films for H2O, O2 and CO2 at 25 µm, 90 % RV and 38 °C Folija Prepustnost H2O pare(g/m2 × 24 h) Prepustnost O2 (105 cm3/m2 24 h Pa) 25 µm Prepustnost CO2 105 cm3/m2 24 h Pa) 25 µm PES 25–30 40–50 300-350 PC 77–93 4559 27351 Ionomer 25–35 6000–7000 6000–7000 EVA 50–60 11000–14000 40000–50000 Najlon 1,1 40–80 507 1925 CA 100–320 2000–3000 15 702 Regenerirana celuloza 5–15 677 987 Orientirani PS 70–150 4500–6000 13169 PVDC 1,5–5,0 8-25 51 Meh~aniI PVC 15–40 Trdi PVC 30–40 150–350 450–1000 Orientirani PE 7 2000–2500 7500–8500 PP 10–12 3748 10130 PE HD 5 1600–2000 30000–40000 PE LD 15–20 650–8500 30000–40000 Pojasnilo kratic: PTFE = politetrafluoretilen PES = poliester EVA = etilvinilacetat CA = celulozniacetat PS = polisulfon PVDC = polivinilidenklorid PVC = polivinilklorid PP = polipropilen PE HD = visokomolekularni polietilen PE LD = nizkomolekularni polietilen Najlon 6, Najlon 11 = vrsti poliamida Tabela 1: Prepustnost kisika pri razli~nih polimernih folijah 4 Table 1: Oxygen permeability of different plastic films 4 Polimerni sloj Prepustnost 109 cm3 cm/(s cm2 kPa) Dimetilsilikonska guma 45,01 Fluorsilikon 8,25 Nitrilna guma 6,37 Naravna guma 1,80 Butilna guma 0,105 Polistiren 0,090 Polietilen HD 0,075 Najlon 6 0,003 Polietilentereftalat 0,0014 Politetrafluoretilen 0,003 disperzije, povr{inske napetosti, viskoznosti in drugih razmer pri delovanju stroja. Postopek izdelave filter medijev na bazi tekstilne podlage in mikroporoznega polimernega sloja obsega naslednje tehnolo{ke faze 11: • pripravo polimerne disperzije • formiranje polimerne mehanske pene (slika 1) • nanos polimerne mehanske pene na tekstilno podlago (slika 2) • su{enje in kalandriranje (slika 2) • utrjevanje (kondenzacija) polimernega nanosa (slika 3) Na sliki 4 je elektromikroskopski posnetek mikropo- rozne mehanske pene, ki je bila izdelana na podlagi navedene recepture in opisanega postopka. Primer recepture za izdelavo mikroporoznih prema- zov: • 400 mas. d. DICRYLAN FL • 300 mas. d. DICRYLAN 7331 • 300 mas. d. DICRYLAN PMC • 100 mas. d. Helizarinweiß RTN • 100 mas. d. SCOTCHGARD FC 251 • 100 mas. d. DICRYLAN–STABILIZATOR FO Vme{avanje zraka: 210 g/L Su{enje: 70–110 °C Kondenzacija: 150 °C 2.2. Izdelava PTFE–membran Po tem postopku se izdelujejo membrane iz polimer- nega prahu PTFE (5). V patentu (W. L. Gore) so za{~i- teni le osnovni principi, tehnolo{ki postopek pa je obse`nej{i in obsega naslednje faze: • ovla`enje PTFE–prahu z ustreznim drsnim sred- stvom; • homogenizacija navla`enega materiala; • oblikovanje predoblikovanca ustreznega profila z ekstrudiranjem; • oblikovanje folije s kalandriranjem; • biaksialno raztegovanje folije; • odstranjevanje drsnega sredstva iz folije; • utrditev folije. V tabeli 3 so prikazani materiali, ki smo jih uporab- ljali pri laboratorijskih in pilotnih preizkusih. Na sliki 5 pa je prikazana kompletna tehnolo{ka shema izdelave PTFE–membran. Tabela 3: Uporabljene surovine za izdelavo PTFE–membran Table 3: The raw materials used to manufacture PTFE–membranes Proizvajalec PTFE–prah Drsno sredstvo DU POINT Fluoropolymer Division Teflon 669 N, 62 N, 636 N ICI Plastics Division Fluon, CD1 HOECHST AG Werk Gendorf Hostaflon TF 2028, TF 2027, TF 2025 SHELL Industrial Chemicals Shellsol T, Shellsol K, Shell Sinarol II V odvisnosti od raztegovanja folije so lahko pore v membrani manj{e ali ve~je 5. V tabeli 4 so podane raz- li~ne vrste Gore–tex–membran, ki so nastale zaradi razli~ne stopnje raztegovanja. Postopek biaksialnega raztegovanja je najzahtevnej{i del proizvodnje PTFE, saj D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 53 Slika 3: Kondenzacija Figure 3: Condensation Slika 2: Nanos, su{enje in kalandriranje Figure 2: Application, drying and calendering Slika 1: Shematski prikaz stroja za formiranje mehanske pene Figure 1: Plan of the machine for mechanical foam Slika 4: Fotografija mikroporozne mehanske pene, izdelane na osnovi poliakrilatne disperzije in fluorkarbona, pove~ava 190–kratna Figure 4: Photo of microporous mechanical foam produced on the basis of polyacrylate dispersion and pluorocarbon, magnification 190–times je potrebna velika hitrost raztegovanja za dosego ustrez- ne poroznosti (podatki iz patenta). Tabela 4: Prikaz strukture in lastnosti razli~nih Gore–tex–membran Table 4: Illustration of the structure and properties of different Gore–tex–membranes velikost por µm debelina mm poroznost % pretok zraka* ml/ (min cm2) prestop vode** bar 0,02 0,080 50 2,3 24,00 0,20 0,060 78 72,0 2,75 0,45 0,080 84 190,0 1,35 1,00 0,080 91 370,0 0,48 3,00 0,025 95 930,0 0,13 5,00 0,025 95 3.870,0 0,07 10,00 – 15,00 0,013 98 11.300,0 0,03 * prepustnost zraka pri P 0,01 bar **minimalna prepustnost vode Na sliki 6 je prikazan vzorec PTFE–membrane, ki smo jo raztegovali samo v eni smeri. Na sliki 7 pa je prikazan vzorec Gore–tex–membrane, ki je laminiran na poliestrskem filcu. S slike je razvidno, da je membranska struktura delno zaprta z lepilom. To pomanjkljivost smo odpravili tako, da smo na tekstilna vlakna nanesli v zelo tanki plasti ustrezno lepilo in na tako pripravljeno pod- logo laminirali PTFE–membrano 8. 2.3 Koagulacijski postopek izdelave polimernih mem- bran S spreminjanjem tehnolo{kih parametrov lahko izde- lamo polimerne membrane z razli~nimi karakteristikami. ^e polimerno raztopino nanesemo na povr{ino ali na nekak{en drug na~in, formiramo plast polimerne razto- pine (v obliki cev~ice) tako, da je na eni strani izpare- vanje topila, na drugi pa neprepustna plast, se tvori koncentracijski gradient (sliki 8 in 10). Pri koagulaciji zato nastane asimetri~na struktura membrane 1,2. Na tistem delu, kjer je manj topila, nastanejo fine pore, na preostalem delu pa gobasta struktura. Koagulacija se opravlja v netopilu za polimer, ki je isto~asno dobro razred~ilo za topilo, v katerem je raztopljen polimer (slika 9). Pogosto se uporablja kot koagulacijsko sredstvo voda ali vodna para. Za raztap- ljanje polimernih materialov pa se uporabljajo v glavnem organska topila. D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... 54 Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 Slika 8: Izparevanje topila pri sloju tanke polimerne plasti Figure 8: Evaporation in thin layer of polymer layers Slika 6: Na{a PTFE–embrana, ki je raztegnjena samo v eni smeri (elekt. mikroskop, pove~ava 1900–kratna) Figure 6: Our PTFE–membrane stretched in one direction only (elect. microscope, magnification 1900–times) Slika 7: Gore-tex PTFE–membrana na PES–filcu 500 g/m2, (elekt. mikroskop, pove~ava 1900–kratna) Figure 7: Gore-tex PTFE–membrane felt in the PES 500 g/m2, (elect. microscope, magnification 1900–times) Slika 5: Postopek izdelave PTFE–membran Figure 5: The process of manufacturing PTFE–membranes V laboratorijskih razmerah smo izdelovali membrane po naslednjem postopku: • priprava polimerne raztopine, • nana{anje tankega sloja na ravno plo{~o (stekleno), • izparevanje pod kontroliranimi pogoji, da se tvori koncentracijski gradient, • koagulacija (voda, organska topila), • ekstrakcija organskih topil (voda, organska topila), • su{enje v su{ilniku. Na sliki 10 je shematsko prikazan laboratorijski po- stopek izdelave asimetri~nih polimernih mambran, ki smo ga uporabljali pri laboratorijskih preizkusih. 2.4 Mo`ni postopek izdelave asimetri~nih polisiloksan- skih membran Po tem postopku formiramo asimetri~no membrano z inverzijo faz. Mehansko stabilnost pa dose`emo kasneje z zamre`enjem 9. Zato je pomembno, da uporabljamo tak{en postopek zamre`enja, da se ne po{koduje asime- tri~na struktura membrane. Polimerne verige imajo navadno od 200 do 800 enot. Kon~ne skupine (X) so najpogostej{e: • (OH) za kondenzacijsko zamre`enje • (CH2 = CH–) za adicijsko zamre`enje • (CH3–) za peroksidno zamre`enje Kot reaktivne skupine (Y) se najpogosteje uporab- ljajo (CH2 = CH–)– skupine. Zamre`evalci morajo imeti najmanj tri funkcionalna mesta, preko katerih se opravlja zamre`evanje. Najpogostej{i zamre`evalci so: • metil–vodik–siloksani z ve~ vodikovimi atomi v molekuli • tri ali tetraalkoksilani • triaminoalkilsilani Reakcijo zamre`enja spro`imo s katalizatorji, ki skraj{ajo reakcijski ~as. Kondenzacijsko zamre`enje, ki se odvija pri sobni temperaturi ob prisotnosti zra~ne vlage, za na{e namene ni primerno, ker nastanejo pri tem reakcijski produkti, ki lahko po{kodujejo relativno nestabilno strukturo membrane. Pri tem postopku se uporabljajo linearni silikoni s kon~nimi skupinami –OH. Kot zamre`evalci pa se uporabljajo: • H– silikoni • estri silicijeve kisline Kot katalizator se najpogosteje uporablja organski kompleks kositra, ki je raztopljen v ustreznem topilu. Kondenzacija poteka tako, da se odcepljajo vodik, voda in alkohol. Reakcijska hitrost je odvisna od reaktivnosti zamre`evalca, katalizatorja in temperature. Pri tempera- turi med 120 °C in 180 °C je ~as zamre`enja med 30 s in 3 min. Zamre`enje poteka tudi pri sobni temperaturi, vendar pa je ~as bistveno dalj{i (do 24 h). Postopek peroksidnega zamre`enja je najstarej{i. Pri tem postopku je lahko veriga siloksan s CH3 – kon~nimi skupinami ali pa s CH2 = CH– stranskimi skupinami. Zamre`evalec v tem primeru najve~krat ni potreben, ker peroksidni katalizator aktivira CH3 – skupino siloksanske verige in nastane preko nje zamre`enje. Kot katalizatorji se najpogosteje uporabljajo razli~ni peroksidi (benzoil- peroksid). Adicijski postopek je najprimernej{i za dokon~no utrditev relativno nestabilne strukture membrane. Pri tem postopku se uporabljajo linearni silikoni s kon~nimi vinilnimi skupinami (CH2 = CH–). Zamre`evalci imajo proste – H atome. Kot katalizatorji se uporabljajo ~isti metali. Pri tem postopku ne nastanejo stranski produkti, ki bi lahko po{kodovali membransko strukturo. Adicijsko zamre`enje se za~ne pri temperaturi med 90 °C in 100 °C in se nato med 140 °C in 170 °C mo~no pove~a. Reakcijska hitrost zamre`enja je pri 150 °C od 15s do 30 s in pri 120 °C od 30s do 60 s. Zamre`enje se lahko dose`e tudi radioaktivnim obsevanjem. Radikali, ki so potrebni za zamre`enje, se namre~ lahko dobijo tudi s pospe{enimi elektroni. Tak{en postopek je primeren predvsem pri kontinuirnem tehnolo{kem postopku zamre`enja silikonske membrane. D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 55 Slika 9: Diagram koagulacije polimerne membrane Figure 9: Diagram of coagulation of polymeric membrane Slika 10: Shematski prikaz izdelave polimernih asimetri~nih membran po postopku inverzije faz Figure 10: Production of asymmetric polymeric membranes by phase inversion process Postopek zamre`enja poteka po naslednjem meha- nizmu: Tehnolo{ki postopek izdelave sestavljajo naslednje tehnolo{ke faze (9): • tehtanje in me{anje komponent; • razred~itev do primerne viskoznosti; • odstranitev mehur~kov, ~e je raztopina preve~ viskoz- na (vakumiranje); • oblikovanje tankega sloja raztopine (ravna plast, formiranje cevnega profila); • v laboratorijskih razmerah je najenostavnej{e nana- {anje z ustreznim no`em na stekleno plo{~o; • su{enje oziroma tvorba koncentracijskega gradienta (odvisno od vrste topil, debeline sloja, temperature itd.); • koagulacija (inverzija faz), pri ~emer nastanejo asi- metri~ne polisiloksanske (silikonske) membrane; • odvisna je od vrste polimernih materialov, topil v polimerni plasti ter vrste koagulacijskega sredstva; • zamre`enje nastale polimerne membrane; • odvisno je od polimernih materialov in vrste zamre- `enja. Preizkuse smo izvedli v laboratorijskih razmerah. Na ravno stekleno plo{~o smo nanesli pribl.0,1 mm debel sloj polisiloksanske raztopine. Pri sobni temperaturi smo spreminjali ~as su{enja v odvisnosti od razmerja topil metiletilketon/toluen. V primeru, da je topilo samo toluen, je ~as su{enja bistveno dalj{i kot pri raztopini s precej{njim dele`em metiletilketona. Ko je nastal za`elen koncentracijski gradient, smo izvr{ili inverzijo faz v cikloheksanonu, ki je netopilo za polisiloksan. Cikloheksanon pa je dobro topilo za toluen in metiletil- keton. Nastalo membrano smo nato osu{ili in s postop- nim dvigovanjem temperature tudi utrdili z zamre`e- njem. Veliko enostavneje lahko poteka postopek izdelave asimetri~nih polisiloksanskih (silikonskih) membran pri izbiri drugih topil in vodi kot koagulacijskem sredstvu. 3 DISKUSIJA Hidrofobne in oleofobne membrane imajo velike prednosti pred drugimi polimernimi in kompozitnimi membranami. Zaradi hidrofobnosti zadr`ujejo vodo tudi z relativno velikimi porami. Oleofobnost pa zagotavlja, da se na povr{ini ne nabirajo razne ne~isto~e. Pri obla- ~ilih se tak{ne membrane uporabljajo za vodo nepre- pustna obla~ila, ki pa prepu{~ajo paro (Gore–tex). Pri filtrih za odpra{evanje prepre~ujejo nabiranje oblog na filterskem mediju. Podobno velja tudi pri membranskih modulih 14, ki se uporabljajo pri komunalnih ~istilnih napravah. Asimetri~ne membrane z zelo fino strukturo bi lahko imele celo vrsto novih aplikacij, ki jih drugi polimerni materiali ne morajo ponuditi. Te aplikacije bi lahko temeljile na zelo veliki prepustnosti plinov. Polisiloksanska folija brez membranske strukture je ve~ kot 100–krat bolj prepustna za CO2, O2, N2 kot druge. ^e bi imela folija {e asimetri~no membransko strukturo, bi bila ta prepustnost {e mnogo ve~ja. Ker pa je polisiloksan hidrofoben, bi prepre~il prepu{~anje vode. Morda je mo`no na osnovi asimetri~ne polisiloksanske membrane izdelati umetne {krge, ki bi omogo~ale uporabo kisika iz vode in odvajanje ogljikovega dioksida v njo. Sposobnost prepu{~anja plinov in zadr`avanje vode, ki jo imajo hidrofobne membrane, bi lahko upo- rabili tudi v raznih drugih separacijskih postopkih. ^e imajo membrane {e oleofobne lastnosti, so {e toliko bolj uporabne. Zato ni ~udno, da so v zadnjem ~asu za~eli uporabljati membranske filterske module na osnovi PTFE, PVDT in {e nekaterih drugih membran v kombi- naciji z biolo{kimi ~istilnimi napravami s suspendirano biomaso 14. V tem primeru se dose`e bistveno ve~ja koncentracija biomase v bazenu in je zato potreben mnogo manj{i volumen. Poleg tega ne potrebujemo bazena za kasnej{e usedanje ter posebne stopnje sterili- zacije. Membrana z velikostjo por 0,2 µm zadr`i poleg velikih delcev tudi vse bakterije. 4 SKLEP Glavna slabost sedanjih hidrofobnih in oleofobnih membran je njihova omejenost za zelo fino filtracijo. Razlog za to je v tehnologiji izdelave tovrstnih membran. Zelo fine membrane z veliko odprtostjo strukture se namre~ izdelujejo po postopku inverzije faz. Po tem postopku `al ni mo`no izdelati PTFE–membrane. Tudi polisiloksanskih membran z asimetri~no strukturo por {e ni na trgu. Morda bo postopek, ki je opisan v tem ~lanku, spodbudil proizvajalce membran, da bi razmi{ljali o predlagani tehnolo{ki re{itvi. 5 REFERENCE 1 Handbook of Industrial Membrane Technology, Edited by: Porter, M. C., 1990 William Andrew Publishing/Noyes D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... 56 Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 2 Rösler, H. W., Membrantechnologie in der Prozessindustrie – Poly- mer membranwerkstoffe, 77, No5, 2005, Chemie Ingenier Technik 3 Müller, K., Kunststoffflaschen und Verschlüssen-Messung und Modellierung der Stofftransportvorgänge, Ph. Thesis, Technische Universität München, 2003 4 Zhang, H., The permeability characteristics of silicone ruber, society for the advancement of material and process engineering, 2006 5 W. L. Gore and Associates Inc., Patent za izdelavo PTFE membran, 1972, Auslegenschrift 21 23 316 6 Wilson, A. J., Foams, Physics, chemistry and structure, 1989 7 Kornev, K. G., Neimark, A. N., Rozhkov, Foam in porous: thermo- dynamic and hydrodynamic peculiartis, Advances in colloid and interface science, 82 (1999) 8 Drev, D., Nov na~in ka{iranja PTFE membran in folij na tekstilno podlogo. Patent 9200375, z dne 18. 10. 1994. Ljubljana: Urad Repu- blike Slovenije za varstvo industrijske lastnine, 1994 9 Drev, D., Postopek izdelave asimetri~ne mikroporozne polisiloksan- ske membrane. Patent 21544. Ljubljana: Urad RS za intelektualno lastnino, 2005 10 Zheng, J. P., Charbel, P. T., Kwok, H. S., Microporous silicon as a light trapping layer for photodiodes, Electrochemical and Solid-state Letters, 3 (2000) 7, 338–339 11 Drev, D., Filtri za otpra{ivanje s mikroporoznim polimernim slo- jevima = Dust filtering with microporous polymer layers. Polimeri, 18 (1997) 5–6, 228–232 12 United States Environmental Protection Agency, Office of water, membrane filtration guidance manual, EPA 815-R-06-009, Novem- ber 2005 13 Glate, J., The early history of reverse osmosis membrane develop- ment, Desalination, 117 (1998), 297–309 14 Pinnekamp, J., Weitergehende Reinigung in kommunalen Kläran- lagen mittels MBR-Technologie, Institut für Siedlungswasser- wirtschaft der RWTH Aachen, oktober 2008 D. DREV, J. PANJAN: NEKAJ MO@NIH POSTOPKOV IZDELAVE HIDROFOBNIH IN OLEOFOBNIH POLIMERNIH ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 51–57 57 K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE AND JEWELLERY MANUFACTURING STALJENE VMESNE POVR[INE MED KOVINO IN KERAMIKO V ZOBOZDRAVSTVU IN PRI IZDELAVI NAKITA Karlo T. Rai}1, Rebeka Rudolf2, Aleksandar Todorovi}3, Dragoslav Stamenkovi}3, Ivan An`el2 1University of Belgrade, Faculty of Technology and Metallurgy, Karnegijeva 4, 11120 Belgrade, Serbia 2University of Maribor, Faculty of Mechanical Engineering, Smetanova 17, 2000 Maribor, Slovenia 3University of Belgrade, School of Dentistry, Clinic for Prosthodontics, 11000Belgrade, Serbia karlo@tmf.bg.ac.rs Prejem rokopisa – received: 2009-10-19; sprejem za objavo – accepted for publication: 2009-12-12 Metal-ceramic fusing has been the essential step in obtaining materials that benefit from both ceramic and metal constituents, i.e. where the combined properties of metal and ceramic layers are desirable. When considering fusing methods, soldering and active metal brazing are the most effective. These processes involve braze melting and flowing between the two pieces of material. In the first part the phenomena occurring on the boundary between the ceramics and the active filler metal during the metal-ceramics joining are discussed. Three interconnected sub-processes are considered: (1) wetting of the ceramic surface, (2) chemical reactions at the interface and (3) diffusion with a moving interface. Then, the appearances at the grain boundary grooves of the ceramic surface are presented as phenomena on the catalytic surface. In the second part, examples from dental practice and jewellery manufacturing are used for comparative analysis. Finally, we discuss the composition and properties of the soldering and brazing alloys used for dental practice and jewellery manufacturing. Key words: metal-ceramic bonding, brazing, soldering, dentistry, jewellery Zlitje kovine in keramike je temeljna stopnja za nastanek materiala, ki pridobi lastnosti obeh sestavnih delov, kjer je `elena kombinacija lastnosti plasti kovine in keramike. Od metod zlitja sta lotanje in aktivno spajkanje najbolj u~inkoviti. Procese sestavljata taljenje in pretok taline med obema materialoma. V prvem delu razpravljamo o procesih na meji med keramiko in aktivnim polnilnim materialom med spajanjem kovine in keramike. Obravnavamo tri povezane procese: (1) omo~ljivost kerami~ne povr{ine, (2) kemi~ne reakcije na vmesni povr{ini in (3) difuzijo s premikajo~o se mejno povr{ino. Nato obravnavamo nastanek `lebov na mejah zrn na kataliti~ni povr{ini. V drugem delu uporabljamo primere iz dentalne prakse in izdelave nakita za primerjalno analizo. Na koncu je razprava o sestavi in lastnostih spajk in lotov, ki se uporabljajo v dentalni praksi in pri izdelavi nakita. Klju~ne besede: vezava keramika-kovina, lotanje, spajkanje, zobozdravstvo, nakit 1 INTRODUCTION The joining of metals to ceramics has been widely practiced in fabricating structural components to utilise the favourable characteristics of the engineering cera- mics. Ceramics offer a huge array of mechanical, thermal and electrical properties.1–5 These range from low thermal conduction in ceramics such as alumina, to high thermal conduction, e. g. diamond, and from low density electrical insulators to superconducting ceramics. Piezo-ceramics offer the almost unique ability to convert electrical energy into mechanical movement, and vice versa. Materials such as beta-alumina and zirconia, which exhibit ionic conduction, are used as sensors. These materials have found application in a wide range of industries, as well as in medicine. A common thread of the applications is that, at some point, the ceramic component interfaces with something else – another part of the component made from another material, usually metal. At this point, some kind of attachment or joint is required. Soldering and active metal brazing are the most effective when considering fusing methods. This process basically involves braze melting and flowing between the two pieces of material. This is commonly referred to as šwetting’ and is absolutely critical – particularly when brazing ceramics. There are many materials that can be fused together to produce joints between materials – those that melt above 450 °C are classed as brazes, and materials melting below 450 °C are called solders. There are two basic methods to encourage wetting: to apply something to the surface of the ceramics so that the braze will wet, or, to put something in the braze that will induce wetting. Surface treatments include metallisa- tion, metal coating and metal hydride treatment, while braze modification involves a process known as active metal brazing. In either case the actual brazing operation takes place either in a controlled atmosphere, such as nitrogen or argon, or a vacuum better than 10–4 N m–2. During active brazing or soldering, samples are heated in a furnace to melt braze or solder which is placed between them. The active element in braze or solder, such as Ti, reacts chemically with the compo- nents to enable wetting and the formation of a strong bond. While active brazing can produce strong bonds, Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 59 UDK 669:666:616.31:739 ISSN 1580-2949 Pregledni ~lanek/Review article MTAEC9, 44(2)59(2010) the high reactivity of the active metal requires the use of a vacuum furnace in an inert atmosphere to prevent oxidation. In addition, the high melting temperature of active brazes can lead to large thermal stresses on cooling, and fracture or de-bonding of the ceramic. Reactive solders reduce mismatch in thermal strains on cooling by lowering the process temperatures, but significant thermal stresses can still be present. Ideally, a strong metallic bond could be formed at/near room temperature, to avoid large thermal stresses between the components. 2 LIQUID METAL/CERAMIC INTERFACES The phenomena occurring on liquid metal/ceramic interfaces are of the greatest importance for metal-cera- mics joining techniques. An interdisciplinary task in which different fields of materials science must parti- cipate together with transport phenomena is shown in Table 1.4–9 Therefore, in this paper, various approaches are explained which are necessary to treat transport phe- nomena on ceramic surfaces. In that sense, three steps are considered: (1) wetting of real ceramic surface, (2) chemical reactions at the interface and (3) diffusion with a moving interface. Table 1: Aspects of investigation Tabela 1: Teme raziskave GENERAL VIEW – physical interface processes (reversible physical forces) – chemical interface processes (irreversible reactions) – phase diagrams – thermodynamic design transport phenomena (macro scale) (i) momentum  slow motion of liquid metal (ii) heat  inner/outer (unsteady and/or steady state transport) (iii) mass  diffusion with homo/heterogeneous chemical reaction (unsteady state phenomena) MICROSRUCTURAL VIEW ceramic liquid metal – surface structure, pores, cracks – structure – grain size – composition – non-bonded areas transport phenomena (micro scale) diffusivity: at the interface and into the bulk mass transfer during homo/heterogeneous appearances ATOMISTIC VIEW ceramic liquid metal (covalent or ionic binding) (metallic bonding) – phase boundary structure – bonding forces – interfacial energy – epitaxy basic equations of transport phenomena  electronic structure of metal/ceramics interface 2.1 Wettability of real ceramic surface In real systems, wettability is quite complex.3–5,7–13 Even sometimes, something as simple as surface rough- ness can have complex and contradictory effects on wettability. Furthermore, chemical segregation in both the solid and liquid phases can have a huge effect on surface and interfacial energies and, hence, wettability. Also, it is difficult to estimate the extent of wetting because interfacial reactions change the wetting charac- teristics as time elapses. The relationship between wetting and interfacial reactions is not well understood, i.e. molten copper shows a high contact angle and low wettability against ceramics such as Si3N4, SiC and ZrO2, although molten aluminium exhibits a low contact angle and high wettability against ceramics. On the other hand, the wetting of an MgO/Al system progresses through three phases: (I) vibratory phase or a chemically quasi- equilibrium phase; (II) constant phase and (III) decreasing phase or a chemically reactive phase. The rate in the decreasing of the contact angle in phase III depends on temperature and the bulk and/or surface properties of the ceramics, since the decrease is caused by the interfacial reactions. Most wetting studies are based on the contact angle, which was first defined empirically by Young. This equation is a balance of the horizontal forces due to surface tensions acting upon a liquid drop in contact with a solid. Young observed that in most cases when a liquid is placed on a solid, the liquid remains as a drop with a definite contact angle between the liquid and solid phases (Figure 1). The physico-chemical principles underlying wetting action are to be found in standard texts. If the liquid metal is (or is not) under pressure, it will come to rest on a real (porous and/or rough) ceramic surface in some position, determined by the advancing smooth surface- liquid metal contact angle and the shape of the pores in the surface. The surface of practical importance is usually grids formed of roughly spherical atom positions and/or grain boundaries. One can start from the well known Young-Dupré Eq. (1) and the relationship given by Cassie and Baxter,13 modified with two coefficients (hLS and hSA),7–9 Eq. (2): Wad = M (1 + cos ) (1) K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... 60 Materiali in tehnologije / Materials and technology 44 (2010) 1, 59–66 Figure 1: Sessile drop configurations assumed for the most reactive wetting models showing a complete reaction product layer Slika 1: Konfiguracija premi~ne kapljice, privzeta pri najve~ modelih reaktivnega omo~enja, ki prikazuje plast reakcijskega produkta cos  = hLS f1 cos  – h SA f2 (2) where, Wad – work of adhesion, J/m2; M – surface energy of the metal; : apparent contact angle (measured by sessile drop method);  – advancing (receding) contact angle (a not yet well understood parameter); f1 – the total area of solid-liquid interface; f2 – the total area of liquid-air (or furnace atmosphere) interface. Areas f1 and f2 can be derived from the value  and the dimensions of the grids formed of roughly spherical atom positions (lattice image) and/or surface grain boundaries. (f1 and f2 are the reflections of real or simulated surface image), hLS – coefficient of net liquid-solid heterogeneity influence, hSA – coefficient of net liquid-air (furnace atmosphere) heterogeneity influence. The microstructural parameters that may have an influence on the properties of the metal/ceramic interface (reflected on hLS and hSA) are: (i) structure of the hetero- phase boundary, (ii) characteristic defects such as steps, faces and dislocations at or to the interfaces, (iii) possible chemical reactions, (iv) reaction products, (v) chemical gradients and segregation of impurities at the interface.14 While the wettability models predict whether wetting will occur for a given system, they do not predict the rate of wetting. Theoreticall modelling the spreading kinetics is complex. Even for nonreactive systems, modelling is complicated by surface irregularities, surface contami- nation, the formation of precursor films, and other surface anomalies. The spreading kinetics of reactive systems is further complicated by interfacial reactions at the solid-liquid and liquid-vapour interfaces (e.g., oxidation). The verification of a spreading mechanism is further hindered by a lack of experimental data. The interfacial reactions and the spreading kinetics are very rapid, making the collection of spreading kinetics and reaction-rate data challenging. Typical spreading kinetics for copper-titanium alloys on alumina is shown in Figure 2.15 2.2 Chemical reactions at liquid metal/ceramic interface To explain the action of a polycrystalline surface it is assumed that reactant molecules of liquid metal are somehow changed, energized, or affected, to form intermediates in the regions close to the surface. In one theory the intermediate is viewed as an association of a reactant molecule with a region of the surface, in other words the molecules are somehow attached to the surface. In another theory molecules are thought to move down into the liquid metal close to the surface and be under the influence of surface forces. In this view the molecules are still, but are nevertheless modified. In still a third theory it is thought that an active complex, a free radical, is formed at the solid surface. This free radical then moves back into the main liquid stream, triggering a chain of reactions with fresh molecules before finally being destroyed. In the contrast with the two theories which consider the reaction to occur in the vicinity of the surface, this theory views the surface simply as a generator of free radicals, with the reaction occurring in the main body of liquid metal.7–9 On the other hand, in order to obtain knowledge of the reactivity, it is very important to respect the changes of the standard free energy for the formation of nitrides, carbides and oxides against temperature. From these thermodynamic data, when metal was placed in contact with nitride, carbide, or oxide ceramics, we are able to predict whether the metal could decompose the ceramic to form a nitride, carbide, or oxide phase. 2.3 Diffusion with a moving interface The observed process, in conjunction with chemical reactions at a solid/liquid boundary, involves diffusion steps. One result of these transient processes is the motion of the boundary between the phases.8 In the general situation, two phases are in contact as in Figure 3. The moving phase boundary is at x = X, and at this boundary Co and C* = C(X, t) represents the equilibrium K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 61 Figure 3: The concentration profile during transient diffusion8 Slika 3: Profil koncentracije med tranzientno difuzijo8 Figure 2: An example of rapid spreading kinetics in the spreading data for 1.2 g samples of copper-titanium alloys on alumina at 1120 °C15 Slika 2: Primer hitre kinetike o {irjenju 1,2 g zlitine baker-titan na korundu pri 1120 °C15 concentration of, e.g. Al2O3, that coexist in ceramics and melted metal respectively at the temperature under consideration. In the insert metal phase, Fick’s second law applies: dC/dt = D d2C/dx2, x > X (3) where C is the Al2O3 concentration during transient diffusion, D the diffusion coefficient of the Al2O3 in melted metal, assumed to be independent of compo- sition. The next condition is that the concentrations on either side of the interface are related by the equilibrium expression of the form: Co = KC*, where K is the partition ratio between the phases. The material balance at the interface takes the form: –D (dC/dx)x=X = (C* – Co) dX/dt (4) Equ. 4 describes the locus of X with time. Relation- ship between Co and C*:  0.5 [C*/(C*  Co)] =  exp 2 erf  (5) when X = 2 (2D)0.5 (6) is well known. The function of ,  exp 2 erf  may be evaluated from standard handbooks. So the D can be calculated if Co, C*, X and t are measured. The illustra- tive calculated surface diffusivities for selected metal/ alumina systems with groove widths (w) from 1 µm to 10 µm, are given in Table 2. Table 2: Surface diffusivities for selected metal/alumina systems10 Tabela 2: Povr{inska difuzivnost za izbrane sisteme kovina/korund10 System Temperature (K) Surface Diffusion,wD (m3 s–1) Ni/Al2O3 1773 4.4 · 10–19 Au/Al2O3 1373 1.8 · 10–24 Cu/Al2O3 1423 2.8 · 10–22 Al/Al2O3 1373 1.1 · 10–19 Groove widths (w) 3 PHENOMENA AT GRAIN BOUNDARY GROOVES Various processes may cause resistance to the overall reaction which occurs at the liquid-solid interface.6,9 For an element of single groove between two grains (Figure 4) one can visualize: (i) Liquid metal film resistance. Reactants diffuse from the main body of the liquid metal to the exterior surface of the solid ceramic specimen. (ii) Groove diffusion resistance. Reactants move through the groove into the ceramic. Most of the reaction takes place within the groove. (iii) Surface phenomena resistance. At some point in their wandering reactant molecules become associated with the surface of the ceramic. They react to give products which are then released to the liquid phase within the groove. (iv) Groove diffusion resistance for products. Products then diffuse out of the groove. (v) Liquid metal film resistance for products. Products then move from the mouth of the groove into the main liquid stream. (vi) Resistance to heat flow. For reactions accompanied by heat release or absorption the flow of heat into or out of the reaction zone may not be fast enough to keep the liquid metal/solid ceramic speci- men isothermal. If this happens, the specimen will cool or heat, strongly affecting the rate. Thus the heat transfer resistance across the liquid film or within the ceramic could influence the rate of reaction. Usually, the reaction layer grew rapidly in the first minute. After approximately a few minutes, further increase in the reaction layer thickness becomes gradual and parabolic in time. In this way, the growth of product layer can be observed as a two-step process: an initial rapid thickening rate and a second parabolic rate, assumed to be diffusion controlled.11 Rapid initial layer growth can be related to the substrate/ceramics surface roughness. On the other hand, slower growth is assisted by capillarity. It is well known that capillary mass transport manifests itself with developing a grain boundary groove on the surface of a polycrystalline material wherever a boundary intersects an interface between a solid and another phase. Mass transport can involve several mechanisms: interfacial diffusion, volume diffusion on either side of the inter- face, and interfacial reaction (solution-precipitation). Depending on the physical characteristics of the system and groove size, one of these mechanisms will be rate controlling, resulting in characteristic groove shapes and growth kinetics. The rate of reaction for the solid ceramic substrate/ liquid metal sample may depend on: 1. Surface kinetics, which may change with develop- ment of the grain boundary grooves. 2. Grain boundary groove (GBG) resistance which sets up internal concentration gradients. The effects of K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... 62 Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 Figure 4: Phenomena around an element of single groove between two grains9 Slika 4: Pojavi okoli elementa enostavnega `leba med dvema zrnoma9 "volume diffusion" and "surface diffusion" on groove development are given in ref.10 3. Interface T or temperature gradient at the liquid metal/solid ceramics interface. This is caused by heat release or absorption during reaction. 4. Film diffusion resistance or concentration gradients across the liquid metal film. The surface kinetics and grain boundary groove (GBG) diffusion cannot be treated as steps in series as they enter the rate equation together. So for a 1st order reaction on a flat plate specimen, analysis gives the rate of reaction: –rA’’’ = k’’’CAs  (7) with  = (tanh MT)/MT (8) Where, : effectiveness factor = (actual reaction rate within GBG)/(rate if not slower by GBG diffusion), MT: Thiele modulus, MT = L(k’’’/Deff)1/2; Deff: effective diffu- sion coefficient in porous solids [m3liquid(metal)/msolid(ceramic) s], L: characteristic size of ceramic sample/specimen, usually flat plate [L = (volume of sample)/( exterior surface) = thickness/2]. The "effective diffusivity" must be measured experi- mentally; it depends generally on the concentration of active species, temperature and on the groove structure. The actual mechanism for diffusion in grooves is complex, since the groove dimensions may be smaller than the mean path of the diffusing molecules. We have no GBG resistance when: MT < 0.4 here  = 1 and –rA’’’ = k’’’CAs; and strong GBG diffusion effects when: MT > 4 here  = 1/ MT and rA’’’ = [(k’’’ Deff)1/2 CAs]/L (see Figure 5). The presented description is similar to the phenomena on a catalytic particle6. 5 DENTAL PRACTICE Dental materials may be classified as preventive materials, restorative materials or auxiliary materials.16–21 Preventive materials include pit and fissure sealants, sealing agents that prevent leakage, materials that are used primarily for their antibacterial effects, and liners, bases, cements, etc. In some cases a preventive material may also serve as a restorative material that may be used for a short-term application. Restorative materials consist of all synthetic compo- nents that can be used to repair or replace tooth structure, including primers, bonding agents, liners, cement bases, amalgams, resin-based composites, hybrid ionomers, cast metals, metal-ceramics, ceramics, and denture poly- mers.16,21 These materials can also be designed as controlled-delivery devices for the release of therapeutic or diagnostic agents. Restorative materials may be used for temporary, short term purposes (temporary cements, and temporary crown and bridge resins) or for longer- term application (dentin bonding, agents, inlays, onlays, crown, removable dentures, fixed dentures, and ortho- dontic appliances). These materials may further be classified as direct restorative materials or indirect restorative materials, depending on whether they are used –intraorally to fabricate restorations or prosthetic devices directly on the teeth or tissues or extraorally, in which materials are formed indirectly on casts or other replicas of the teeth and other tissues. An ideal resto- rative material would be biocompatible, bond with other visible tissues, exhibit properties similar to those of tooth enamel, dentin and other tissues, etc. Auxiliary dental materials are substances which are used in the process of fabricating dental prostheses and appliances, but which do not become part of these devices. These include acid-etching solutions, impre- ssion materials, casting investments, gypsum cast and model materials, dental waxes, acrylic resin for impre- ssion and bleaching trays, acrylic resins for mouth guards and occlusion aids, and finishing and polishing abrasives. K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 63 Figure 6: Schematic illustration of a cross-sectional view of a natural anterior tooth and supporting tissues (left) in combination with restorative dental materials (right) Slika 6: Shemati~en prikaz prereza naravnega zoba in nosilnega tkiva (levo) v kombinaciji s restavrativnim dentalnim materialom (desno) Figure 5: Effectiveness factor () v. Thiele modulus (MT) 6 Slika 5: Faktor u~inkovitosti () v odvisnosti od modula Thiele (MT) 6 Figure 6 is a schematic cross-section of a natural tooth and supporting bone and soft tissue (left) in combi- nation with restorative dental materials (right). Under healthy conditions, the part of the tooth that extends out of adjacent gingival tissue is called a clinical crown, and that below the gingival tissue is called the tooth root. The crown of a tooth is covered with enamel and the root is covered with cementum, and it consists of dentin and tissue within one or more root canals. The surface of implantable biomaterials is in direct contact with the host bone and soft tissue and it plays a critical role in determining its biocompatibility, func- tional compatibility, osteoinduction of bone and osseo- integration of implants. The goal of regenerative therapy around orthopaedic and dental implant devices is to create a suitable environment in which the natural biolo- gical potential for functional regeneration of the bone and of the periodontal ligament can be stimulated and maximised. In implantology, wide arrays of biomaterials have been used so far with variable success rates.17–20 Metals are probably the oldest form of materials used for dental implants and the most common type of materials used so far. Superior fracture and fatigue resistance caused them to become the materials of choice in traditional load-bearing applications. A large spectrum of metals and alloys were used as endosseous implants, in orthopaedics, cranio-maxillofacial surgery and dentistry (Table 3). Gold-based alloys were among the first alloys to be used for implants, probably because these alloys were available in dentistry, but also because they promoted the fibrous interface with a bone, i.e. a distant osteogenesis with a short lifespan. Cobalt-chro- mium alloys were also developed and used as endosseous implants. However, the fundamental problem with these metals and alloys was the fibrous response which they promoted with the bone. By today’s standards, none of these materials achieve osseointe- gration, probably because of their corrosion effected by the living tissue and the release of elements into the tissue. Today, these metals have largely been replaced by titanium and titanium alloys. Bioceramics are a group of ceramics which are biologically active materials rich in calcium and phosphate. Hydroxyapatite and tricalcium phosphate are similar in composition to bone and teeth and can be used for the augmentation of alveolar ridges and filling bony defects. The various forms of bioceramics are Single crystals (Sapphire), Polycrystalline (Hydroxyapatite) Glass (Bioactive glass) Glass ceramics (Ceravital) Composites (Stainless steel reinforced Bioglass). There are about four types of bioceramics: INERT: Attached by compact morphological fixation, e.g. Alumina, Carbon POROUS: Attached by vascularization through pores, e.g. Porous Alumina. SURFACE ACTIVE: Directly attach by chemical bon- ding with bone, e.g. Bioglass, Hydroxyapatite RESORBABLE Designed to be slowly replaced by bone, e.g. Tricalcium Phosphate (TCP). Ceramic coatings on dental implants, such as Calcium phosphates (hydroxyapatite- HA) appear to have better biological response than cpTi or Ti alloy alone, even if their clinical predictability remains contro- versial. Coatings seem to promote faster bone adaptation, higher bone implant strength, better osteoblast precursor activity, bone growth around dental implants and thus, bonding of the bone to the implants, i.e. osseointe- gration. HA coatings were considered highly because they enhance osseointegration, and they lead to the formation of more mineralised extracellular matrix (ECM) and to faster bone formation with respect to Ti substrates alone. The philosophy was that HA has an K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... 64 Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 Table 3: Example of dental and medical materials and their applications Tabela 3: Primeri dentalnih in medicinskih materialov in njihove uporabe MATERIALS PRINCIPAL APLICATIONS Metals and Alloys 316 Stainless Steel-CP-Ti, Ti-Al-V, Ti-Al-Nb, Ti-13Nb-13Zr, Ti-Mo-Zr-Fe Co-Cr-Mo, Cr-Ni-Cr-Mo Ni-Ti Gold Alloys Silver products Platinum and Pt-Ir Hg-Ag-Sn amalgam Fracture fixation, stents, surgical instruments, bone and joint replacement, fracture fixation, dental implants, pacemaker, encapsulation Bone and joint replacement, dental implants, dental Bone plates, stents, orthodontic wires Dental restoration Antibacterial agents Electrodes Dental restorations Ceramics and Glasses Alumina Zirconia Calcium Phosphates Bioactive glasses Porcelain Carbons Joint replacement, dental implants, dental Various parts of dental replacement Bone repair and augmentation, surface coatings on metals Bone replacement Dental restorations Heart valves, percutaneous devices, dental implants advantage over smooth Ti surfaces having: (i) a bioactive surface versus an inert Ti surface; (ii) higher bond strength of the bone to the implant; and (iii) increased bone-to-implant contact. In spite of reports about overstressing, rapid bone-resorption adjacent to HA-coated implants, short-term survival rates (from 6 months to 6 years), or other causes of failures, there is still evidence of positive effects of HA coatings on osseointegration. Recently, other ceramic coatings, such as titanium nitride (TiN) and titanium carbide (TiC), have been proposed for implantology. Most dental solders are either gold-based or silver- based alloys (Table 4). These alloys contain special elements such as tin to encourage a lower melting range and better flow. Gold-based solders are used primarily to solder cast alloys, whereas silver solders are used more in orthodontic applications. A variety of gold-based solders are available with various melting ranges to meet specific applications. The manufacturer of the cast alloys generally specifies the compatibility of solders with their alloys, and it is common for a manufacturer to offer both a series of dental casting alloys and compatible solders together. The strength, hardness, and corrosion of dental solders also depend on the composition. In general, solders with a higher melting range are stronger and harder than lower fusing solders. 5 JEWELLERY MANUFACTURING The joining of many different materials is necessary in jewellery manufacture (the metal frames of rings, earrings, necklaces, etc. with different kinds of stone – precious, semi-precious, synthetic stones, coral, and pearls). One of the phases in jewellery manufacturing is placing the precious stone into the metal frame. The stones are initially cut into the desired sizes and shapes, and then polished. The metal frame cast is ground and polished and the stones are joined to it by adhesives, soldering, or by mechanical clamping.22 The attachment between stone and metal becomes weaker over time, because of elastic deformation. If the stone is soldered to the metal frame the joint between stone and metal is stronger. Different stones have different mechanical and physical properties and it is necessary to find an appropriate combination of solder components for each stone. For example, sapphire is an AlB2BOB3B mineral. Sapphire can be soldered after bonding or brazed without bonding. If the sapphire is brazed to titan without bonding, an alloy with 72 % Ag could be used. Diamond could be brazed to steel with L-Ag40Cd, L-Ag30Cr, or L-Ag20Si. There are two groups of precious stones according to their price: first-order stones (diamond, ruby, sapphire and emerald) and second-order stones (aquamarine, topaz, zircon, onyx, quartz topaz, etc.). Many of them have similar chemical compositions, but their colours are different. The chemical composition of diamond is 99.95 % carbon. Emerald is a green mineral made of alumi- nium, and its colour comes from chrome or vanadium. Sapphire and ruby are minerals of AlB2BOB3B. Sapphire could have a blue colour from titan and iron, and violet from vanadium. Ruby could have a red colour from chrome or a chestnut colour from chrome and iron. The chemical compositions of stones could be very important when joining different kinds of stones. Many stones in the second group are very sensitive to heat. The result of heating the stone above its critical temperature could be K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 65 Table 4: Examples of solders used in dental practice Tabela 4: Primeri lotov, ki se uporabljajo v dentalni praksi Name of the solder, (Chemical composition, %) Melting interval Form of solder Operation temperature Colour Application Auroker L 1040 PF (79.0 % Au, 17.0 % Ag, 3.0 % Pt, 1.0 % Zn) 980 – 1040 °C Strip (0.25 1) mm 1040 °C yellow soldering metal-ceramics gold alloys 7 Aurodur HL 750 (73.0 % Au, 11.0 % Ag, 2.0 % Pt, 11.0 % Zn, 3.0 % Cu) 710 – 750 °C Strip (0.25 1) mm 750 °C yellow soldering dental alloys and alloys for metal-ceramics application Auroker L 1060 (72.0 % Au, 5.0 % Ag, 12.0 % Pd, 2.0 % Zn, 8.0 % Cu, 1.0 % In) 1010 – 1070 °C Strip (0.25 1) mm 1060 °C white before firing ceramics Midor L (58.0 % Au, 21.0 % Ag, 18.0 % Cu, 2.0 % Zn, 1.0 % Sn) 790 – 820 °C Plate 0.33 mm Strip (0.25 1) mm 820 °C Lightyellow joining gold alloys for casting Silopal L 890 (26.0 % Au, 16.0 % Ag, 30.0 % Cu, 15.0 % Pd, 13.0 % Zn) 850 – 890 °C Strip (0.25 1) mm 890 °C white soldering alloys based on the silver-palladium, before firing (baking) and the midfielder element for non-precious and precious metals Witex lot (39.0 % Co, 22.0 % Cr, 19.0 % Ni, 10.0 % Fe, 5.0 % Si, 3.0 % Mo, 1.0 % B, 1.0 % V) Wire Ø 1 mm 1180 °C white joining of materials for mobileprosthetic, CoCr alloys a change in colour. Turquoise has a critical temperature at 250 °C, when its colour then changes from sky blue to green. At the critical temperature interval of tanzanite (from 400 °C to 500 °C), the nuances of yellow gra- dually disappear and the blue colour is intensified. Other precious stones can also express changes in their colours (quartz, topaz, amethyst, aquamarine, etc.). Therefore, it is very important during the joining of stones to the metal frame, that the stones are not heated above their critical temperatures. The composition of brazing alloys used for jewellery applications is Au-Ag-Cu, with the addition of metals such as zinc and cadmium. Cadmium is very toxic, especially its vapour, during the melting process, and is replaced with non-toxic elements. Cadmium could be substituted with the following metals: tin (Sn), indium (In), gallium (Ga), and zinc (Zn). Also, some solders without cadmium have very high concentration of Ni (5 % to 9 %). According to EU Directive 76/769 EES (January 2000), the concentration of nickel in jewellery cannot exceed the maximum of 0.05 %. This EU Directive is a prerequisite for the prevention of allergic reactions caused by nickel. Nano-foil, with properly selected solder or braze components (without Ni), placed between metal parts could prevent allergic reactions to nickel and could enable a high quality of metal-metal joints. 6 CONCLUSIONS Attempts to investigate the phenomena occurring on liquid metal/ceramics interfaces during metal ceramic joining in dental practice and jewellery manufacturing are still in their early stages. In this paper we carefully explain only the basic concepts and approaches necessary to treat wetting with diffusion phenomena accompanied by chemical reactions. The attention is confined to the typical examples taken from the actual practice and literature. Nowadays, in dentistry, attention is increasingly focusing on the extensive range of alternative materials. These new materials include titanium and cobalt/nickel base alloys and all ceramic crowns. The latter have excellent aesthetic properties, but do not have the long-term clinical approval that gold has. For example, zirconia has only passed clinical tests during the last 8 years. In addition, the CAD/CAM techniques associated with the use of these materials are, in many cases, prohibitively expensive. Moreover, the problems of poor aesthetics often associated with porcelain-fused-to-metal techniques might be addressed through the use of extremely high gold content alloys, which have recently been patented. It is considered that, as CAD/CAM technology progresses, special high gold alloys should be developed which are well suited to milling and grinding operations and sufficiently high strength for long span bridges and small cross sections. Besides long term clinical approval and longevity, the most important advantages of gold alloys are easy workability, biocompatibility, aesthetics and maximum range of indications. ACKNOWLEDGEMENT This paper is part of Eureka project E!4213 NANO-FOIL. The authors gratefully acknowledge the Ministry of Higher Education, Science and Technology of the Republic of Slovenia and the Ministry of Science and Technological Development of the Republic of Serbia. 7 REFERENCES 1 Z. Samard`ija, R. B. Marinenko, S. Bernik, M. Kosec, M. ^eh, Mater. Technol., 37 (2003), 1–2 2 T. Kopar, V. Ducman, Mater. Technol., 37 (2003), 1–2 3 S. Bernik, M. Podlogar, N. Daneu, A. Re~nik, Mater. Technol., 42 (2008), 2 4 G. Elssner, G. Petzow, ISIJ Int., 30 (1990), 1011 5 A. Meier, D.A. Javernick, G. R. Edvards, JOM, February (1999), 44 6 O. Levenspiel, Chemical reaction engineering 1999, (John Wiley & Sons, Inc.) 7 K. T. Raic, Unsteady or starting phenomena at liquid metal/ceramic interfaces, Advances in Science and Technology, 32 (2003), 725–733 8 K. T. Raic, An estimate of Si3N4 diffusion into Cu based filler metal, Ceramics International, 26 (2000), 19 9 M. Mihailovic, T. Volkov Husovic, K. T. Raic, Micro- and nano- scale wetting of reactive metal at metal-ceramic interface, Advances in Science and Technology, 45 (2006), 1526–1531 10 E. Saiz, R. M. Cannon, A. P. Tomsia, Acta Mater., 47 (1999), 4209 11 M. D. Baldwin, P. R. Chidambaram, G. R. Edwards, Metallurgical and Materials Transactions A, 25A (1994), 2497 12 Y. Nakao, K. Nishimoto, K. Saida, ISIJ Int., 30 (1990), 1142 13 A. B. D. Cassie, S. Baxter, Trans. Faraday Soc., 40 (1944), 546 14 W. Mader, M. Ruhle, Acta Metall., 37 (1989), 953 15 A. Meier, P. R. Chidambaram, G. R. Edwards, J. Mat.Sci., 30 (1995), 3791–3798 16 R. G. Craig, J. M. Powers, Restorative Dental Materials, 11th ed., Ed., Mosby, Inc., 2002 17 Gross U, H. J. Schmitz, V. Strunz, Surface activities of bioactive glass, aluminium oxid and titanium in a living environment. Ann NY Acad Sci, 523 (1988) 211–26 18 W. Q. Yan, T. Nakamura, M. Kobayashi, H. M. Kim, F. Miyaji, T. Kokubo, Bonding of chemical treated titanium implants to bone. J Biomed Mater Res., 37 (1997) 2, 267–75 19 M. Weinlaender, Bone growth around dental implants. Dental Clin North Am 35 (1991), 3, 585–601 20 Hench LL, Wilson J. Surface-active biomaterials. Science, 226 (4675) (1984), 630–6 21 D. Stamenkovi} et al. Stomatolo{ki gradivni materijali, 2003, ZUNS, Beograd 22 K. T. Raic et al., Nanofoils for Soldering and Brazing in Dental Joining Practice and Jewellery Manufacturing, 1st International Conference on Materials and Technology sponsored by FEMS and IUVSTA,13-15 October (2008), Portoro`, Slovenia, published in Mater. Technol., 43 (2009)1, 3 K. T. RAI] ET AL.: LIQUID METAL/CERAMIC INTERFACES IN DENTAL PRACTICE ... 66 Materiali in tehnologije / Materials and technology 44 (2010) 2, 59–66 ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC STAINLESS STEELS VPLIV VKLJU^KOV MnS NA KOROZIJSKO ODPORNOST AVSTENITNIH NERJAVNIH JEKEL ^rtomir Donik, Irena Paulin, Monika Jenko Institute of Metals and Technology, Lepi pot 11, SI-1000 Ljubljana, Slovenia crtomir.donik@imt.si Prejem rokopisa – received: 2009-09-19; sprejem za objavo – accepted for publication: 2009-12-21 The influence of alloy sulfur on the crevice solution chemistry and the nature of the surface in the crevice during initiation were investigated. Two austenitic steels with different amounts of sulfur in the alloy, AISI 316L and SS 2343, were compared using electrochemical measurements with an additional standard test – ASTM G48. The electrochemical measurements, i.e., a potentiodynamic measurement, a linear polarization and a Tafel polarization, were performed in a typical corrosion solution of 3.5 % NaCl and 0.1 % NaCl. The results show the difference for the steels in a low concentration of Cl– ions, while in the higher concentration of the Cl– ions the steels act very similarly. The MnS inclusions were analyzed and characterized using high-resolution SEM with EDX elemental analyses. The SE images show large differences for these materials in terms of the amount of inclusions. From the EDX measurements it was confirmed that the inclusions consist of Mn and S, and very few other types of inclusions were found. Keywords: corrosion, stainless steel, MnS inclusion, electrochemical corrosion tests, ASTM G48 Raziskovali smo vpliv vsebnosti `vepla na korozijske lastnosti dveh avstenitnih nerjavnih jekel. Z elektrokemijskimi meritvami in s standardnimi preizkusi jami~aste korozije ASTM G48 smo posku{ali ugotoviti razlike med dvema preiskovanima materialoma. Preiskovali smo dve po sestavi podobni nerjavni jekli tipa AISI 316L in SS 2343 z razli~nima vsebnostma `vepla. Elektrokemijske meritve: potenciodinamske meritve, linearne polarizacije in Tafelove polarizacije, smo naredili v tipi~ni raztopini za korozijske preizkuse 3,5 % in 0,1 % NaCl. Pri nizkih koncentracijah kloridnih ionov v raztopini lahko opazimo odmike med lastnostmi materialov, medtem ko prihaja pri visokih koncentracijah kloridnih ionov do tega, da imata materiala zelo podobne korozijske lastnosti. S pregledom pod vrsti~nim elektronskim mikroskopom visoke lo~ljivosti z EDX analizatorjem smo dolo~ili obliko in vrsto vklju~kov ter njihovo sestavo. Pri SEM-posnetkih lahko opazimo razlike med materialoma po koli~ini vklju~kov in potrdimo z EDX-analizo, da vklju~ki vsebujejo elementa Mn in S ter da skoraj ni prisotnih drugih vklju~kov. Klju~ne besede: korozija, nerjavna jekla, MnS-vklju~ki, elektrokemijski korozijski preizkusi, ASTM G48 1 INTRODUCTION The effect of sulfur on the corrosion behavior of austenitic stainless steel is manifested through the behavior of the sulfide inclusions due to the low solu- bility of sulfur in ferrous metals.1–5 Sulfide inclusions have been recognized as preferential sites for localized corrosion for almost 100 years, for steels, and for almost 60 years for stainless steels.1,6–13 The effect of an alloying element can be manifested through effects on the passive film, the local solution chemistry, or the interfacial electrochemical kinetics.10,14–23 The effects of alloyed sulfur on the localized corrosion in steels have been ascribed by various workers to each of these mecha- nisms. In our stainless steel, SS 2343, the sulfur is added to enhance lubrication through inclusions in the oil-free processes.2,22–24 Many theories have proposed that changes in the occluded chemistry are the key due to the effects of the dissolution products of the MnS inclusions.2,23 Because most dissolved sulfur species have been observed to enhance metal dissolution, many sulfur species (i.e., sul- fate, thiosulfate, elemental sulfur, and aqueous sulfide) have been predicted to result from MnS inclusion dissolution in the occluded region. Two categories of sulfur species may be considered, depending on the origin of the species: (a) sulfur species in the environ- ment (H2S, HS–, S2O3, SO2, HSO4–, SO42–) and (b) sulfur in the material (sulfur in solid solution, sulfur segregated at the surface or in the grain boundaries, and sulfur in the sulfide inclusions).4,25 In the case of the MnS sulfate inclusions, although their determination effect on the corrosion of steel and stainless steel has been extensively studied, the exact mechanism is not known. Sulfate (SO42–) was predicted by Eklund to be the dominant dissolved sulfur species originating from sulfide inclusions via the electrochemical oxidation of the metal sulfide in the inclusions to sulfate, with the simultaneous formation of acid.26 The alteration of the chemistry of the stainless-steel surface in the vicinity of the sulfide inclusions, and the subsequent effects on the interfacial kinetics, has also been proposed as playing an important role in localized corrosion initiation. It was also observed that the corrosion initiated primarily at the inclusion edges, and concluded that the bare metal initially exposed upon inclusion dissolution became passivated and did not Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 67 UDK 669.14.018.8:620.193 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(2)67(2010) corrode until the precipitation of a MnCl2 salt film (found using AES) occurred to restrict the mass transport. Smialowska27 proposed that localized corro- sion initiated by the formation of a thick salt layer at weak spots in the passive film (such as might be expected at inclusions), followed by hydrolysis of the salt leading to local acidification, inclusion dissolution and exposure of the bare, unprotected metal. Marcus et al.27 proposed that regardless of its origin (i.e., from the alloy or from the solution), sulfur enhances metal dissolution via the formation of an adsorbed sulfur layer, which leads to weakening of the metal-metal bonds on the surface, thereby reducing the activation energy barrier for metal dissolution. The important aspect of the corrosion is the forma- tion of the passive film on the top of stainless steels. After the formation of this passive film on the bare metal surface, a slow dissolution of the metal cations conti- nues, involving dissolution at the passive film surface and transport of the ions through the oxide. The question then arises as to where the bulk impurities (e.g., sulfur) go during the process. Above the critical concentration of sulfur at the interface (metal/oxide) the breakdown of the passive film was observed.4 There is much evidence for the damaging effect of the sulfur species over a wide range of corrosion-related service failures. The relation between the sulfur-induced corrosion mechanism presented in areas of practical importance can be rationalized on the basis of (a) the source of the sulfur, (b) the transport process to the metal surface, and (c) the condition of the reduction of the sulfur species into the harmful chemical state of sulfur. Although it is clear that sulfide inclusions act as preferential sites for localized corrosion, the mechanism by which this occurs has remained unclear. In the present work two stainless steels with different amounts of sulfur in the alloy were studied. The study was conducted using different electrochemical techni- ques: potentiodynamic measurements, linear polarization and Tafel polarization. MnS inclusions were investigated using SEM with EDX measurements.28.29 2 EXPERIMENTAL Two types of stainless steel were investigated. Their compositions were confirmed using an Optical Emission Spectrometer (OES) as follows (weight percent): 17.0 % Cr, 10.0 % Ni, 1.5 % Mn, 0.40 % Si, 0.035 % P, 0.002 % S, 0.015 % C and the remainder Fe, denoted as AISI 316L; and 16.5 % Cr, 10.5 % Ni, 1.5 % Mn, 0.38 % Si, 0.041 % P, 0.025 % S, 0.020 % C and the remainder Fe, denoted as SS 2343. The experiments for the electrochemical measure- ments were carried out in 3.5 % sodium chloride solution with and without 0.5-M Na2S, and in 0.1 % sodium chloride solution. The standard test ASTM G48 was performed in a 6 % FeCl3 solution at room temperature. All the chemicals were from Merck, Darmstadt, Ger- many. The test specimens were cut into discs of 15 mm diameter. The specimens were abraded with SiC emery paper down to a 1000-grit prior to the electrochemical measurements. The specimens were then embedded in a Teflon PAR holder and employed as a working electrode. The reference electrode was a saturated calomel elec- trode (SCE, 0.242 V vs. SHE) and the counter electrode was a high-purity graphite rod. Tafel polarization, potentiodynamic and linear polari- zation measurements were recorded using an EG&G PAR PC-controlled potentiostat/galvanostat Model 263 with SoftCorr computer programs. In the case of the potentiodynamic measurements and the linear polari- zation the specimens were immersed in the solution 1 h prior to the measurement in order to stabilize the surface at the open-circuit potential. Potentiodynamic curves were recorded starting at 250 mV more negative than the open-circuit potential. The potential was then increased, using a scan rate of 1 mV s–1, until the transpassive region was reached. Linear polarization curves were recorded at ±10 mV around the open-circuit potential using a scan rate of 0.1 mV s–1. The Tafel polarization curves were recorded at ±250 mV around the open-circuit potential using a scan rate of 0.5 mV s–1. Some of the samples were passivated in 3M HNO3 for 30 min to prevent localized corrosion on the boldly exposed surfaces. The surface analytical experiments were performed using a high-resolution Scanning Electron Microscope JEOL JSM-6500F. The standard test method ASTM G 48 for the pitting corrosion resistance of stainless steels and related alloys uses ferritic chloride solution. The method is used to compare the resistance of the stainless steels and related alloys in order to increase the resistance to pitting and crevice-corrosion initiation under the specific conditions. In our experiments Method A of standard ASTM G48 – ferric chloride pitting test was used: 6.0 % FeCl3(aq) with test specimens of 25 mm × 50 mm without sharp edges. All the specimens in the test had the same dimensions and were grinded with silicon grinding paper 600. The samples were weighed before and after exposure, and exposed to the solution for 72 h. 3 RESULTS AND DISCUSSION 3.1 Polarization measurements In order to study the influence of the sulfur inclusions in stainless steels, different test solutions were used. The linear polarization curves were recorded for both materials on samples around the corrosion potential (Figure 1). The calculations were performed from linear polari- zation measurements using the equation: Rp = a c/(2.3 Icorr (a + c)). ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... 68 Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 The polarization resistance, Rp, is evaluated from the linear polarization curves by applying a linear least- squares fit of the data around ±10 mV Ecorr. The corro- sion current, Icorr., is calculated from Rp, the least-squares slope, and the Tafel constants, a and c, of 100 mV decade–1. The value of E(I = 0) is calculated from the least-squares intercept. The corrosion rate was calculated by using the following conversion formula: vcorr. = C(EW/d)(Icorr./A), where EW is the equivalent weight of the sample in g, A is the sample area in cm2, d is its density in g/mL, and C is a conversion constant that is dependent upon the units desired. The slopes of both materials were similar, so the corrosion rates of these materials did not differ significantly from each other. The corrosion potential for the stainless-steel AISI 316 L was –220 mV and for SS 2343 it was –230 mV. The potentiodynamic curves were measured in two different solutions: 0.1% and 3.5 % NaCl, with and without the addition of Na2S·9H2O. Figure 2 shows the potentiodynamic curves of the stainless steels AISI 316 L and SS 2343 in 0.1 % NaCl without any pre-passiva- tion of the materials. After 1 h of stabilization at the open-circuit potential, the corrosion potentials (Ecorr) for both materials were around –350 mV. Following the Tafel region both alloys exhibited passive behavior. The passive region is limited by the breakdown potential (Eb), which corresponds to the oxidation of water and the transpassive oxidation of the metal species. The breakdown potential for SS 2343 was approximately 380 mV and for AISI 316L it was 600 mV. Figure 3 shows the potentiodynamic curves of the stainless steels AISI 316L and SS 2343 in 3.5 % NaCl. After 1 h of stabilization at the open circuit, the potential values of Ecorr for both alloys were around –200 mV, although it was slightly higher for AISI 316L. The values of Eb are similar for both steels and are around 200 mV and 250 mV for SS 2343 and AISI 316L, ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 69 Figure 4: Potentiodynamic curves of the stainless steels AISI 316 L and SS 2343 in 0.1 % NaCl and 0.5-M Na2S solution Slika 4: Potenciodinamski krivulji nepasiviranih vzorcev nerjavnih jekel AISI 316 L in SS 2343 v 0,1-odstotni raztopini NaCl v 0,5 M Na2S Figure 2: Potentiodynamic curves of the stainless steels AISI 316 L and SS 2343 passivated samples in a 0.1 % NaCl solution Slika 2: Potenciodinamski krivulji pasiviranih vzorcev nerjavnih jekel AISI 316 L in SS 2343 v 0,1-odstotni raztopini NaCl Figure 3: Potentiodynamic curves of the stainless steels AISI 316 L and SS 2343 in a 3.5 % NaCl solution Slika 3: Potenciodinamski krivulji vzorcev nerjavnih jekel AISI 316 L in SS 2343 v 3,5-odstotni raztopini NaClFigure 1: Linear polarization curves of the stainless steels AISI 316 L and SS 2343 in a 0.1 % NaCl solution Slika 1: Graf linearne polarizacije nerjavnih jekel AISI 316 L in SS 2343 v 0,1 % raztopini NaCl respectively. The breakdown potentials for these alloys are relatively low due to a narrowing of the passive region in the presence of a high concentration of the Cl– ions. As a result, the corrosion current values are relatively high, and therefore the corrosion rates are around mm per year (Table 1). Table 1: Corrosion rates calculated from Rp for the stainless steels AISI 316 L and SS 2343 in (10–3 mm per year) Tabela 1: Korozijske hitrosti, izra~unane iz Rp, za nerjavni jekli AISI 316 L in SS 2343 v (10–3 mm na leto) Material vcorr. (10–3 mm/y) NaCl NaCl+Na2S AISI 316 L 140.6 46.75 SS 2343 158.9 53.93 The results mentioned above suggest that the dis- tinction between the selected materials should be investigated in a low concentration of Cl– ions. There- fore, the corrosive activator 0.5M Na2S was added to the solution and the corrosion characteristics were measured. The pH of the test solution evidently increased; there- fore, the corrosion potential was moved to more negative values to approximately –500 mV (Figures 4 and 5), whereas the corrosion current and the corrosion rate do not increase significantly (Table 1). In the presence of S2– ions both steels exhibited better corrosion characteristics. However, we have to take into account that the pH is evidently higher than in the case of the test solution without the addition of the sulfide and that at approxi- mately pH 12 the stainless steels are in a highly passi- vated area. In the 0.1 % solution of NaCl, both steels differed significantly, the passive region was evidently wider for the AISI 316 L compared to the SS 2343 (Figure 2). With the addition of 0.5-M Na2S the values of Ecorr were shifted from approximately –350 mV to –500 mV and –550 mV for the AISI 316L and SS 2343, respectively (Figure 4). The extension of the passive region increased with the addition of the sulfide. The Eb values were shifted to 450 mV and 750 mV for the AISI 316L and SS 2343, respectively (Figure 4). 3.2 Scanning Electron Microscopy Figure 6 shows a SE image of an inclusion on the surface of the SS 2343; the EDX analysis was also performed on the marked areas. The results are shown in Table 2. The EDX measurements inside the inclusion showed the increased amount of Mn and S, with an atomic ratio 1 : 1, which indicates the presence of MnS in the inclusion (Spectrum 1 and 3). Spectrum 2 was performed on the bulk material and its composition confirmed the chemical composition of the standard composition of this type of stainless steel. 3.3 Standard corrosion test Standard corrosion test for the pitting corrosion resistance of stainless steels and related alloys in the ferritic chloride solution, ASTM G 48, was performed for both alloys. The results are shown in Table 3, where the weight loss is shown in absolute mass and also in ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... 70 Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 Figure 6: SE image of an inclusion in the sample SS 2343 with the areas of the EDX measurements Slika 6: SE-posnetek povr{ine vzorca s podro~ij EDX-analize Figure 5: Potentiodynamic curves of the stainless steels AISI 316 L and SS 2343 in 3.5 % NaCl and 0.5-M Na2S solution Slika 5: Potenciodinamski krivulji nepasiviranih vzorcev nerjavnih jekel AISI 316 L in SS 2343 v 3,5-odstotni raztopini NaCl v 0,5 M Na2S Table 2: EDX measurements of the inclusion in SS 2343 in weight percent Tabela 2: EDX-meritve vklju~ka na povr{ini jekla SS 2343 v masnih dele`ih (%) C O Mg Al S Ti Cr Mn Fe Ni Spectrum 1 27.31 1.08 24.34 14.40 0.82 4.66 21.59 3.82 1.98 Spectrum 2 2.25 0.67 16.67 1.77 64.37 10.82 Spectrum 3 29.60 5.68 50.09 4.15 1.30 mass percentage. The results show a slight difference in the weight loss, which could be the result of an increased amount of sulfur manganese inclusions. This small increase in the weight loss for the sample with the higher amount of sulfur inclusions is in the range of the measurement uncertainty and it could also be correlated to either a lower corrosion characteristic of the metal or a difference in the preparation of the samples. Table 3: Results for ASTM G 48 test Tabela 3: Rezultati preizkusa ASTM G 48 m1/g m2/g m/g m/% AISI 316 L 3.0098 2.9195 0.0903 3.0 SS 2343 3.1527 2.9959 0.1568 4.9 4 CONCLUSION The materials AISI 316 L and SS 2343 are chemically almost identical; the only difference is in the amount of sulfur, which is up to 10 times higher in the SS 2343. The results from the electrochemical measurements show a slight difference in the corrosion characteristics in favor of the AISI 316 L. However, the distinction between the selected materials is possible only in the solution containing low Cl– concentrations. In contrast, in the presence of high Cl– concentrations, the test solution is so aggressive that the difference in the chemical composition of both materials is diminished. The Na2S addition in the corrosion solutions increased the pH to around pH 12, and therefore pushed the stainless steel into the passive region so the S2– ions did not play an important role in the differentiation of these to stainless steels. The difference between both materials is expressed by a standard corrosion test, ASTM G 48, where the AISI 316L showed a slightly better corrosion characteristic than the SS 2343 stainless steel. 5 REFERENCES 1 Brossia, C. S.; Kelly, R. G., Occluded solution chemistry control and the role of alloy sulfur on the initiation of crevice corrosion in type 304ss. Corrosion Science 40 (1998) 11, 1851–1871 2 Krawiec, H.; Vignal, V.; Heintz, O.; Oltra, R.; Chauveau, E., Dissolution of chromium-enriched inclusions and pitting corrosion of resulfurized stainless steels. Metallurgical and Materials Transac- tions a-Physical Metallurgy and Materials Science 37A (2006) 5, 1541–1549 3 Krawiec, H.; Vignal, V.; Heintz, O.; Oltra, R., Influence of the dissolution of MnS inclusions under free corrosion and potentiostatic conditions on the composition of passive films and the electro- chemical behaviour of stainless steels. Electrochimica Acta 51 (2006) 16, 3235–3243 4 Sulfur-Assisted Corrosion Mechanisms and the Role of Alloyed Elements. In Corrosion Mechanisms in Theory and Practice, Marcus, P., Ed. CRC Press: 2002 5 Schmuki, P.; Hildebrand, H.; Friedrich, A.; Virtanen, S., The composition of the boundary region of MnS inclusions in stainless steel and its relevance in triggering pitting corrosion. Corrosion Science 47 (2005) 5, 1239–1250 6 Brossia, C. S.; Kelly, R. G., Influence of alloy sulfur content and bulk electrolyte composition on crevice corrosion initiation of austenitic stainless steel. Corrosion 54 (1998) 2, 145–154 7 Donik, C.; Kocijan, A.; Mandrino, D.; Paulin, I.; Jenko, M.; Pihlar, B., Initial oxidation of duplex stainless steel. Applied Surface Science 255 (2009) 15, 7056–7061 8 Kocijan, A.; Milosev, I.; Pihlar, B., The influence of complexing agent and proteins on the corrosion of stainless steels and their metal components. Journal of Materials Science-Materials in Medicine 14 (2003) 1, 69–77 9 Kocijan, A.; Milosev, I.; Merl, D. K.; Pihlar, B., Electrochemical study of Co-based alloys in simulated physiological solution. Journal of Applied Electrochemistry 34 (2004) 5, 517–524 10 Kocijan, A.; Donik, C.; Jenko, M., Electrochemical and XPS studies of the passive film formed on stainless steels in borate buffer and chloride solutions. Corrosion Science 49 (2007) 5, 2083–2098 11 Jenko, M.; Mandrino, D., Application of surface analytical technique HRAES to metallurgy. Strojarstvo 41 (1999) 3–4, 107–110 12 Mandrino, D.; Vrbanic, D.; Jenko, M.; Mihailovic, D.; Pejovnik, S., AES and XPS investigations of molybdenum-sulfur-iodine-based nanowire-type material. Surface and Interface Analysis 40 (2008) 9, 1289–1293 13 Donik, C.; Kocijan, A.; Paulin, I.; Jenko, M., Oxidation of duplex stainless steel at moderately elevated temperature. Materiali in tehnologije / Materials and technology 2009, In Press, Corrected Proof 14 Abreu, C. M.; Cristóbal, M. J.; Losada, R.; Nóvoa, X. R.; Pena, G.; Pérez, M. C., Comparative study of passive films of different stainless steels developed on alkaline medium. Electrochimica Acta 49 (2004) 17–18, 3049–3056 15 Gojic, M.; Marijan, D.; Kosec, L., Electrochemical behavior of duplex stainless steel in borate buffer solution. Corrosion 56 (2000) 8, 839–848 16 Kocijan, A.; Milosev, I.; Pihlar, B., Cobalt-based alloys for ortho- paedic applications studied by electrochemical and XPS analysis. Journal of Materials Science-Materials in Medicine 15 (2004) 6, 643–650 17 Kocijan, A.; Donik, C.; Jenko, M., The corrosion behaviour of duplex stainless steel in chloride solutions studied by XPS. Materiali in Tehnologije 43 (2009) 4, 195–199 18 Kocijan, A.; Donik, C.; Jenko, M., The electrochemical study of duplex stainless steel in chloride solutions. Materiali in Tehnologije 43 (2009) 1, 39–42 19 Olefjord, I.; Clayton, C. R., Surface-composition of stainless-steel during active dissolution and passivation. Isij International 31 (1991) 2, 134–141 20 Olsson, C. O. A.; Landolt, D., Passive films on stainless steels - chemistry, structure and growth. Electrochimica Acta 48 (2003) 9, 1093–1104 21 Stypula, B.; Stoch, J., The characterization of passive films on chromium electrodes by XPS. Corrosion Science 36 (1994) 12, 2159–2167 22 Bastos, I. N.; Tavares, S. S. M.; Dalard, F.; Nogueira, R. P., Effect of microstructure on corrosion behavior of superduplex stainless steel at critical environment conditions. Scripta Materialia 57 (2007) 10, 913–916 23 Wranglen, G., Pitting and sulfide inclusions in steel. Corrosion Science 14 (1974) 5, 331–349 24 Yin, Z. F.; Zhao, W. Z.; Tian, W.; Feng, Y. R.; Yin, C. X., Pitting behavior on super 13Cr stainless steel in 3.5% NaCl solution in the presence of acetic acid. Journal of Solid State Electrochemistry 13 (2009) 8, 1291–1296 25 Marcus, P.; Talah, H., The sulfur-induced breakdown of the passive film and pitting studied on nickel and nickel-alloys. Corrosion Science 29 (1989) 4, 455–463 ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 71 26 Eklund, G. S., Initiation of pitting at sulfide inclusions in stain- less-steel. Journal of the Electrochemical Society 121 (1974) 4, 467–473 27 Lukac, C.; Lumsden, J. B.; Smialowska, S.; Staehle, R. W., Effects of temperature on kinetics of passive film growth on iron. Journal of the Electrochemical Society 122 (1975) 12, 1571–1579 28 Mandrino, D.; Godec, M.; Torkar, M.; Jenko, M., Study of oxide protective layers on stainless steel by AES, EDS and XPS. Surface and Interface Analysis 40 (2008) 3–4, 285–289 29 Paulin, I.; Donik, C.; Jenko, M., Mechanisms of HF bonding in dry scrubber in aluminium electrolysis. Materiali in Tehnologije 43 (2009) 4, 189–193 30 International, A., ASTM G48-03(2009) Standard Test Methods for Pitting and Crevice Corrosion Resistance of Stainless Steels and Related Alloys by Use of Ferric Chloride Solution. ^. DONIK ET AL.: INFLUENCE OF MnS INCLUSIONS ON THE CORROSION OF AUSTENITIC ... 72 Materiali in tehnologije / Materials and technology 44 (2010) 2, 67–72 I. PAULIN ET AL.: AES AND XPS CHARACTERIZATION OF TITANIUM HYDRIDE POWDER AES AND XPS CHARACTERIZATION OF TITANIUM HYDRIDE POWDER PREISKAVE PRA[KA TITANOVEGA HIDRIDA S SPEKTROSKOPIJO AUGERJEVIH ELEKTRONOV IN RENTGENSKO FOTOELEKTRONSKO SPEKTROSKOPIJO Irena Paulin, Djordje Mandrino, ^rtomir Donik, Monika Jenko Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia irena.paulin@imt.si Prejem rokopisa – received: 2009-12-01; sprejem za objavo – accepted for publication: 2009-12-24 Titanium hydride powder, manufactured by ball milling titanium in a hydrogen atmosphere down to micron-sized particles, was analyzed by X-ray Photoelectron Spectroscopy (XPS). A strong titanium oxide signal was also measured, which decreased somewhat after intense sputtering of the sample, but was impossible to get rid of completely. This was probably due to the high surface/volume ratio of each TiHx particle, which contributes to a substantial titanium oxide/TiHx ratio, and due to the surface morphology of the powder sample, which leaves a considerable part of the oxide layer shaded during the sputtering. Auger Electron Spectroscopy (AES) was then employed and some characteristic differences in the shape of the Ti LMM spectra between the TiHx and Ti were observed; however, they can be ascribed to the TiHx only after a comparison with the same types of spectra measured on titanium oxide. An additional XPS measurement was performed with TiHx powder and powdered Ti. The peaks were fitted with Ti oxide and metallic Ti components and the shifts of the metallic component between the Ti and TiHx (a shift of 0.4 eV was expected) were checked for. Keywords: titanium hydride, ball milling, AES, XPS Pra{ek titanovega hidrida, sintetiziran z mletjem titana v krogli~nem mlin~ku v vodikovi atmosferi do mikrometrske velikosti delcev, smo analizirali z rentgensko fotoelektronsko spektroskopijo (XPS). Pri tem smo izmerili mo~an signal titanovega oksida, ki je sicer upadel po intenzivnem ionskem jedkanju povr{ine vzorca, ni pa izginil. Verjeten razlog je visoko razmerje povr{ina/volumen posameznih delcev TiHx, ki prispeva k znatnemu razmerju titanov oksid/TiHx, kakor tudi morfologija povr{ine pra{kastega vzorca, zaradi katere je del oksida zasen~en med ionskim jedkanjem. S spektroskopijo Augerjevih elektronov (AES) smo opazili nekatere karakteristi~ne razlike v obliki Ti LMM-spektrov med TiHx in Ti, vendar jih je bilo mogo~e pripisati TiHx {ele po primerjavi z enakim tipom spektrov, izmerjenih pri titanovem oksidu. Nadaljnjo meritev XPS smo opravili pri pra{ku TiHx in pra{ku Ti. Dolo~ili smo oksidne in kovinske komponente vrhov in preverjali premike kovinskih komponent pri TiHx glede na Ti (premik naj bi bil 0,4 eV). Klju~ne besede: titanov hidrid, krogli~ni mlin~ek, AES, XPS 1 INTRODUCTION Titanium hydride can be used as a catalyst in the reversible dehydrogenation of other hydrides and carbon nanotubes.1,2,3 It is also used as a catalyst in the prepa- ration of titanium compounds,1,4,5 as a source of pure hydrogen,1 in the manufacturing of ceramic and glass seals from a mixture of active metal titanium or titanium hydride in powder form 1,6 and titanium coatings.1 It is also well known as a blowing agent in the production of aluminum foams and some other foam-like structures produced by powder metallurgy.7,8 For these purposes extensive studies of titanium hydride as well as of dehydrogenation and their effect on the formation of alloy foams have been performed.9,10 In this study an attempt was made to apply surface-analysis techniques (AES and XPS) to titanium hydride powder, manu- factured by ball milling titanium in a hydrogen atmo- sphere, that is used for the commercial production of aluminum foam. It was shown previously that the characteristic signatures from ultra-high-vacuum (UHV) deposited thin-film titanium hydride can be obtained using these techniques.11,12 In our case the titanium hydride was in the shape of micron-sized powder particles. Therefore, in the XPS a strong titanium oxide signal was also measured, which decreased somewhat after the intense sputtering of the sample, but was impossible to remove completely due to the high surface/volume ratio of each TiHx particle, which contributes to the substantial titanium oxide/TiHx ratio, and due to the surface mor- phology of the powder sample, which leaves a consi- derable part of the oxide layer shaded during the sputtering. This oxide component also influences the shape of the Ti LMM peaks. Thus, while relying on the same type of titanium hydride, the characteristic features in AES and XPS spectra, as observed by Lisowski et al.,11 had to be resolved from rather complex data (due to the complexity of the system in this study compared to the system as described by Lisowski et al. 11). 2 EXPERIMENTAL Titanium hydride powder, milled titanium, polished titanium plate and titanium oxide powder samples were fixed onto sample holders for the SEM/AES/XPS inve- Materiali in tehnologije / Materials and technology 44 (2010) 2, 73–76 73 UDK 543.428.3:620.1/.2 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(2)73(2010) stigations by means of UHV-compatible double-sided sticky tape. Th esputter cleaning of the samples was performed under UHV conditions inside the main vacuum chamber of the SEM/AES/XPS apparatus, where the samples were introduced via a fast-entry air-lock. The SEM imaging as well as the AES and XPS depth profiling of the samples were performed using a VG-Scientific Microlab 310F SEM/AES/XPS. For all the XPS measurements, Mg K radiation at 1253.6 eV with an anode voltage × emission current = 12.5 kV × 16 mA = 200 W was used. For the XPS profiling measure- ments an Ar+ energy of 3 keV at 1 µA ion current over a 6 × 6 mm2 area was used. Similar ion-beam parameters with a 4 × 4 mm2 area were used for the AES profiling. A rough estimate for the sputtering rate during the XPS profiling parameters is of the order of 1 nm/min, and 2 nm/min for the AES. The AES and XPS spectra were acquired using the Avantage 3.41v data-acquisition & data-processing soft- ware supplied by the SEM/AES/XPS equipment manu- facturer. Casa XPS software 13 was also used for detailed data processing. 3 RESULTS AND DISCUSSION A SEM image of the particles of as-received Ti hydride powder manufactured by the ball milling of titanium in a hydrogen atmosphere is shown in Figure 1. It can be seen that powder grains are of irregular shape, ranging over more than two orders of magnitude in size: from below 0.5 µm to well over 50 µm. In Figure 2, high-resolution XPS spectra of the Ti 2p from pure Ti (Figure 2a, line 1) and Ti hydride (Figure 2a, line2), as measured by Lisowski et al. 11 on UHV- deposited thin films, are shown. The Ti 2p of Ti hydride at 1200 s (top), 8400 s and 36000 s sputtering times are shown in Figure 2b. The Ti 2p of the milled Ti at 1200 s (top), 4800 s and 8400 s sputtering times are shown in Figure 2c. While in the spectra of Lisowski et al. 11 the metallic-type Ti seems to be predominant, with an appro- ximately 0.4 eV shift towards higher binding energies for Ti 2p3/2, in the case of the Ti hydride, the spectra measured on the ball-milled Ti and the Ti hydride are much more complex. Ti 2p peaks obtained were fitted with the Ti 2p3/2 and Ti 2p1/2 components of metallic Ti, TiO2 and TiOx.14,15 The metallic Ti 2p3/2 is the right-most component. The Ti hydride powder appears to be covered by a TiO2 layer, which is partially removed or/and reduced only after prolonged sputtering (Figure I. PAULIN ET AL.: AES AND XPS CHARACTERIZATION OF TITANIUM HYDRIDE POWDER 74 Materiali in tehnologije / Materials and technology 44 (2010) 2, 73–76 Figure 2: High-resolution XPS spectra of Ti 2p from pure Ti (dotted line 1) and Ti hydride (line 2) as measured by Lisowski et al. 11 on UHV-deposited thin films; Ti 2p of Ti hydride at 1200 s (top), 8400 s and 36000 s sputtering times (b); Ti 2p of milled Ti at 1200 s (top), 4800 s and 8400 s sputtering times (c). Slika 2: Visokolo~ljivi spektri XPS-prehodov Ti 2p s ~istega Ti (~rtkana ~rta 1) in Ti- hidrida (~rta 2), kot so jih izmerili Lisowski in sodelavci 11 na tankih plasteh, nanesenih v UHV; Ti 2p s Ti-hidrida po 1200 s (zgoraj), 8400 s in 36 000 s ionskega jedkanja (b); Ti 2p z mletega Ti po 1200 s (zgoraj), 4800 s in 8400 s ionskega jedkanja (c). Figure 1: SEM image of the Ti hydride powder Slika 1: SEM slika pra{ka Ti hidrida 2b). The changes to the milled Ti with sputtering are much less pronounced, with TiOx 16 and metallic Ti gaining slightly versus TiO2 (Figure 2c). The average binding- energy value for the metallic Ti 2p3/2 component for the milled Ti can be determined from the XPS measurements after all sputtering cycles as (454.4 ± 0.1) eV. In Figure 3 the deviations, EB, of the metallic Ti 2p3/2 binding energy with sputtering time for Ti hydride from the average value of 454.4 eV determined for the milled Ti are shown. Also shown are the linear deviation trend and the average deviation, which is close to 0.4 eV. This can be interpreted as the Ti hydride characteristic shift in the metallic Ti 2p3/2 binding energy observed in Ti hydride thin films.11 The highly scattered data are probably an artifact of the measurement as well as the fitting procedure. It is, nevertheless, possible to observe a declining deviation trend or even hypothesize about a possible switch from a high valued deviation (0.5–0.6 eV) to a low valued one (0.1–0.2 eV) in the 20 000–30 000 s sputtering time range. Whatever the precise form of this decline, it suggests that the Ti hydride powder I. PAULIN ET AL.: AES AND XPS CHARACTERIZATION OF TITANIUM HYDRIDE POWDER Materiali in tehnologije / Materials and technology 44 (2010) 2, 73–76 75 Fig 4: Ti LMM Auger transitions from pure Ti (a1) and Ti hydride (a2) as measured by Lisowski et al. 11 on UHV-deposited thin films; Ti LMM Auger transitions from pure Ti (b1, c1), Ti hydride (b2, c2) and Ti oxide (b3, c3) after 1200 s (b1, b2, b3) and 8400 s (c1, c2, c3) of sputtering. Slika 4: Augerjevi prehodi Ti LMM s ~istega Ti (a1) in Ti-hidrida (a2), kot so jih izmerili Lisowski in sodelavci11 na tankih plasteh, nanesenih v UHV; Augerjevi prehodi Ti LMM s ~istega Ti (b1, c1), Ti-hidrida (b2, c2) in Ti-oksida (b3, c3) po 1200 s (b1, b2, b3) in 8400 s (c1, c2, c3) ionskega jedkanja. Figure 3: Deviation of the metallic Ti 2p3/2 binding energy with sputtering time for Ti hydride from the average value of 454.4 eV obtained for the milled Ti. The linear deviation trend (the thick solid line) and the average deviation (the thin horizontal line), which is close to 0.4 eV, are also shown. Slika 3: Odvisnost razlike med vezavno energijo za kovinski Ti 2p3/2 Ti-hidrida in povpre~no vrednostjo 454,4 eV, izmerjeno pri mletem Ti, od jedkalnega ~asa. Prikazana sta tudi linearen trend razlike (debela neprekinjena ~rta) in povpre~na vrednost razlike (tanka vodoravna ~rta), ki je blizu 0,4 eV. grains may not be of homogeneous consistency, with (more) hydrogen being closer to the grain surface. In Figure 4a are the Ti LMM Auger transitions from pure Ti (Figure 4a1) and Ti hydride (Figure 4a2) as measured by Lisowski et al. 11 on UHV-deposited thin films. The Ti LMM Auger transitions from the polished Ti plate (Figure 4b1, 4c1), the Ti hydride powder (Figure 4b2, 4c2) and the TiO2 powder (Figure 4b3, 4c3) after 1200 s (Figure 4b1, 4b2, 4b3) and 8400 s (Figure 4c1, Figure 4c2, Figure 4c3) of sputtering are also shown. The differences in the peak shapes between all three samples can be observed and they become more pronounced with sputtering. The differences between the Ti hydride and the Ti do not completely agree with the differences found by Lisovski et al. 11; however, it can be seen from Figure 4 that they are not of the same nature as the differences between the Ti oxide and the Ti. The most useful features are the less pointed shape of the maximum at approximately 375 eV 11, nearly dissappeared forking of TiH minimum between 385–390 eV and the small hydrogen-induced peak around 443 eV 11 in the spectra measured on the Ti hydride. Some other features also characteristic of Ti hydride may appear in these samples due to the Ti oxide. 4 CONCLUSIONS XPS and AES characterizations of ball-milled Ti hydride powder were attempted using the same characte- ristic signatures in both methods as for Ti hydride thin films manufactured in UHV. Due to the manufacturing process the Ti hydride in this study was found to be covered with a layer of oxide that could not be com- pletely removed, so the characteristic signal of Ti hydride could barely be extracted by peak fitting in the case of the XPS and had to be verified by comparison with metallic Ti and Ti oxide spectra in the case of AES. An additional benefit from the fitting of the Ti 2p is the suggestion that the titanium hydride grain may not have a homogeneous composition. To verify this, a cross-sec- tional AES study of individual grains by immersing titanium hydride powder in a low-melting-point alloy, and polishing the cross-section, a technique already used for study of soft-magnetic powders 17, is planned. 5 REFERENCES 1 R. S. Venilla, A. Durygin, M. Merlini, Z. Wang, International Journal of Hydrogen Energy, 33 (2008), 6667–6671 2 R. S Kumar, A. L. Cornelius, M. G. Pravica, M. F. Nicol, M. Y. Hu, P. C. Chow, Diam. & Relat. & Mater., 16 (2007), 1136–1139 3 I. O. Bashkin, A. I. Koleshnikov, M. A. Adams, M. F. Nicol, M. Y. Hu, P.C. Chow, J. Phys. Condens. Matter., 12 (2000) 4757–4765 4 M. N. Ozyagcilar, M.W. Davis; Catalysts for synthesis of ammonia, United States Patent 4623532, 1986 5 V. E. Antonov, I. O. Bashkin, V. K. Fedotov, S. S. Khasanov, T. Hansen, A. S. Ivanov, Phys Rev B, 73 (2006)5, 054107–054112 6 I. Paulin, C. Donik, M. Jenko, Mechanisms of HF Bonding in Dry Scrubber in Aluminium Electrolysis, Materiali in Tehnologije, 43 (2009) 4, 189 7 PM Proa-Flores, RAL Drew, Advanced Engineering Materials, 10 (2008)9, 830–834 8 A. Ibrahim, C. Koerner, R. F. Singer, Advanced Engineering Materials, 10 (2008), 845–848 9 Y. W. Gu, M. S. Yong, B. Y. Tay, C. S. Lim, Materials Science & Engineering C, 29 (2009), 1515–1520 10 M. Vijay, V. Selvarajan, K. P. Sreekumar, J. Yu, S. Liu, P.V. Anantha- padmanabhan, Solar Energy Materials & Solar Cells, 93 (2009), 1540–1549 11 W. Lisowski, A. H. J. van den Berg, D. Leonard, H. J. Mathieu, Sur- face and Interface Analysis, 29 (2000), 292–297 12 A. Kocijan, C. Donik, M. Jenko, The Corrosion Behaviour of duplex stainless steel in chloride solutions studied by XPS, Materiali in Tehnologije, 43 (2009) 4, 195 13 Fairley N, CasaXPS VAMAS Processing Software. Available from World Wide Web: http://www.casaxps.com/ 14 Chastain J. (ed.), Handbook of X-ray Photoelectron Spectroscopy, Physical Electronics, Eden Prairie 1995 15 C. D. Wagner, A.V. Naumkin, A. Kraut-Vass, J. W. Allison, C. J. Powell, J. R. Rumble (eds.), NIST X-ray Photoelectron Spectroscopy Database. Available from World Wide Web: http://srdata.nist.gov/xps 16 C. Donik, A. Kocijan, I. Paulin, M. Jenko, The Oxidation of Duplex Stainless Steel at Moderately Elevated Temperatures, Materiali in Tehnologije, 43 (2009) 3, 137 17 M. Godec, Dj. Mandrino, B. [u{tar{i~, M. Jenko, Surface and Inter- face Analysis, 34 (2002), 346–351 I. PAULIN ET AL.: AES AND XPS CHARACTERIZATION OF TITANIUM HYDRIDE POWDER 76 Materiali in tehnologije / Materials and technology 44 (2010) 2, 73–76 P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 PRELOMNE ZNA^ILNOSTI LEDEBURITNEGA Cr-V-JEKLA VANADIS 6 Peter Jur~i1, Borivoj [u{tar{i~2, Vojteh Leskov{ek2 1 Czech Technical University in Prague, Faculty of Mechanical Engineering, Karlovo nám. 13, 121 35 Prague 2, Czech Republic 2Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia p.jurci@seznam.cz Prejem rokopisa – received: 2009-11-16; sprejem za objavo – accepted for publication: 2009-12-08 The P/M Vanadis 6 is a very popular tool steel used for different cold work applications. The specimens for the three point bending tests and fracture toughness test were austenitized at two different temperatures, quenched with or without the sub-zero treatment and then double tempered. The bending strength was lower by higher austenitizing temperature and also with the use of sub-zero treatment. The average fracture toughness was lower by higher austenitizing temperature also and increased slightly with the sub-zero treatment. The lowering of bending strength can be considered as logical result because of the austenitic grain growth, as well as increased internal stresses in the case of sub-zero processed samples. For fracture toughness, the situation seems to be more complex. On one side, the effect of austenitizing temperature on fracture toughness is similar as for bending strength, on the other side, the sub-zero treatment induces a rather slight increase of fracture toughness. The paper presents some details on the experiments performed including possible explanation of the material’s behaviour considering microstructure investigations. Key words: ledeburitic Cr-V PM steel, heat treatment, sub-zeto treatment, bending strength, fracture toughness, fracture surface PM-jeklo Vanadis 6 je zelo popularno orodno jeklo za razli~no uporabo v hladnem. Preizku{ance za tri to~kovni upogib in `ilavost loma smo avstenitizirali pri dveh razli~nih temperaturah, kalili z obdelavo in brez nje pod temperaturno ni~lo in dvakrat popustili. Ugotovili smo, da je upogibna trdnost ni`ja po kaljenju z vi{je temperature in po obdelavi pod ni~lo. @ilavost loma je bila ni`ja po kaljenju z vi{je temperature in je bila nekoliko ve~ja po obdelavi pod ni~lo. Zmanj{ana upogibna trdnost je logi~en rezultat pove~anja velikosti avstenitnih zrn in pove~anja notranjih napetosti po obdelavi pod ni~lo. Primer `ilavosti loma je bolj kompleksen, kajti na eni strani temperatura avstenitizacije vpliva podobno na upogibno trdnost in `ilavost loma, po drugi pa obdelava pod ni~lo pove~a `ilavost loma.V ~lanku so predstavljeni detajli o opravljenih preizkusih in predstavljena je razlaga vedenja materiala z upo{tevanjem mikrostrukture. Klju~ne besede: ledeburitno PM CrV-jeklo, toplotna obdelava, obdelava pod ni~lo, upogibna trdnost, `ilavost loma, povr{ina preloma 1 GENERAL REMARKS Ledeburitic steels have a high wear resistance and hardness and are used in many industrial operations like: metal cutting, wood cutting, fine blanking, bending etc. In these operations, the material must meet various requirements. It has to withstand compressive stresses, abrasive and/or adhesive wear, but also chipping and total tool collapse. To meet these demands, the ledeburitic steels must have an optimal chemistry, as well as phase constitution. Moreover, a proper heat treatment must be performed before the use. The ledeburitic steels contain a high content of carbon and alloying elements that form carbides responsible not only for high hardness and wear resistance but also for the capability of the material to achieve a high strength and acceptable fracture tough- ness after heat treatment. For the resistance to chipping and total collapse, two basic characteristics are import- ant: three-point bending strength as a measure for the resistance against the crack initiation and the fracture toughness that characterises the resistance of the material against crack propagation. The three-point bending strength is sensitive to various factors. First of all it is important how the mate- rial was made. It is known that the large carbide net- works, clusters or bands are responsible for a significant lowering of three-point bending strength and its anisotropy 1,2. This is also the main reason why the steels made via the powder metallurgy (P/M) of rapidly solidified particles have higher bending strength than the materials of the same chemistry made by conventional metallurgy. The three-point bending strength is also influenced by the material cleanless. It was reported that oxides, sulphides and other impurities can decrease the three-point bending strength significantly even at relatively low quantity 3. The quality of the surface should also not be neglected. Some investigations have confirmed that the three-point bending strength decreases as the surface roughness increases 4. The three point-bending strength decreases with the increase of the austenitizing temperature because of the coarsening of the austenitic grains during austenitization 5. On the other hand, the fracture toughness depends mainly on the matrix hardness and its ductility. Gene- rally, it also depends to its relation to the carbide network Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 77 UDK 669.14.018.252:539.42 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(2)77(2010) (cohesive strength at carbide-matrix boundary i.e. shape, size and size distribution of carbide particles, carbide clusters, distance among the particles etc). Ledeburitic steels have to be heat treated up to approximately HRc 60 (or more in some cases); and their fracture toughness is relatively low. For the conventionally manufactured chromium ledeburitic steels, as well as high speed steels, it decreases with the increasing austenitizing temperature and also with increasing tempering temperature, but only up to the maximum of secondary hardness 6. The orien- tation of carbides and/or the state of the carbides (net- works or bands) play only minor role 4,7,8. On the other hand, the role of carbide distribution is highlighted in the soft state of the steels 7,8 the microstructure with carbide network had the lower fracture toughness, although the wrought material had better values. No data, except the early work of Olsson and Fischmeister 9, are known on the assessment of fracture toughness of P/M made lede- buritic steels. However, results published in the men- tioned paper are negatively influenced by the state of manufacturing technique, porosity of the consolidated material, and therefore they did not have a sufficient value for today’s manufactured and used materials. The ideas about the sub-zero treatment on the mecha- nical properties of ledeburitic steels was developed over last few decades 10,11. This technique has early gained scientific interest, but the attempts of its application did not produce doubtless results. Kulmburg et al. 11, for instance, reported for the M2-type steel a reduction of the amount of retained austenite after sub-zero treatment at –196 °C for 1 h and slight lowering of three point bending strength. Berns 10, on the other hand, reported a significant hardness increase for the sub-zero processed X290Cr12 ledeburitic steel. Last decade has brought new accurate investigations on the nature of possible impro- vement of material characteristics after the sub-zero processing. Stratton 12 pointed out, that especially for the D2 ledeburitic steel, the improvement of wear resistance due to the sub-zero processing could be increased to extremely high values. According to investigations of Collins and Meng 13–15, it is assumed that deep cooling leads to an arrangement of the structure responsible for the hindering of the dislocation movement. This induces a strengthening of the material. In addition, Meng assumed 14 that martensite originated at very low tempe- ratures can differ from that transformed at higher temperatures in the lattice parameter. The martensitic transformation is probably superposed with the precipi- tation of nanosized carbides particles coherent with the matrix. However, up to now the principal explanation of the positive effect on essential properties (hardness, wear resistance) is not entirely clear. Moreover, also an optimal processing route of the cryogenic processing is not clearly determined and the opinions about it are different. The main goal of our experimental effort was a reliable investigation on the influence of various factors (heat treatment, sub-zero treatment, tempering) on fracture toughness and three point bending strength of the PM ledeburitic steels. The Vanadis 6 P/M steel produced by Uddeholm was selected as example because of its relatively simple chemistry. 2 EXPERIMENTAL The experimental material was the ledeburitic steel Vanadis 6 with nominally 2.1 % C, 1.0 % Si, 0.4 % Mn, 6.8 % Cr, 1.5% Mo, 5.4 % V and Fe as balance, manu- factured with P/M (HIP of rapidly solidified particles) and soft annealed to the hardness of HV10 = 284. Two types of specimens were prepared for the investigations. Samples for the three point bend testing (numbered 1 to 28) had a cross section of (10 × 10) mm and a length of 100 mm. They were ground to obtain a final surface roughness of 0.2–0.3 µm. The second type of specimens were the circumferentially notched and pre-cracked specimens (numbered 33 to 69) prepared according to the method described in 6,16 (see also Figure 1). The specimens were then submitted to the heat treatment procedure that included vacuum austenitization up to temperature 1000 °C or 1050 °C, nitrogen gas quenching at 5 bar pressure and double tempering, each cycle at 550 °C for 1 h. In selected cases, the sub-zero treatment in liquid nitrogen was inserted between the quenching and tempering. The parameters of heat-treat- ment, the average KIC and final hardness and are summarized in Table 1. The fracture toughness was determinated according to the method published elsewhere 6,16. At least eight samples were tested for a given parameter of the heat treatment. Three point bending tests have been carried out at following parameters: distance between supports of 88.9 mm, loading in the central region and loading rate of 5 mm/min up to fracture. The results of bend strength testing and hardness measurements are given in Table 2. P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 78 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 Figure 1: Precracked and fractured V notched cylindrical samples for KIC determination Slika 1: Prelomljeni valjasti preizku{anci z V zarezo za dolo~itev KIC z vnaprej{njo razpoko The microstructure of the material was investigated with light microscopy (LM) and scanning electron microscopy (SEM). The fractures surfaces were inve- stigated with SEM. Hardness measurements were made using the Vickers method at a load of 98.1 N (HV10) as well as using the Rockwell "C" method (HRC). 3 RESULTS AND DISCUSSION The microstructure of the as-received material is shown in Figure 2a and 2b. The LM micrograph (Figure 2a) gives us the overview. One can see that the material consists of the matrix and carbides particles that are very fine and uniformly distributed throughout the material. The SEM micrograph, (Figure 2b) shows that carbide particles of different size with the maximum size of particles is around 3 µm and the smallest have a size less than 0.5 µm. The steel microstructure after heat processing is shown in Figure 3a–d. It consists of the martensitic matrix and of eutectic and part of secondary carbide particles. These particles are very fine and uniformly distributed in the matrix. The amount of carbides is slightly smaller for the specimens processed at higher austenitizing temperature. This is natural as higher temperature normally leads to the dissolution of large amount of secondary carbides. Heat treatment leads to a substantial hardness increase. With increasing austenitizing temperature the hardness increases. It is natural because higher auste- nitizing temperature induces a better dissolution of secondary carbides and higher supersaturation level of martensite after the quenching. On the first sight, it is rather surprising that the sub-zero processing reduces the hardness only slightly. Several authors, for instance Berns 10, Kulmburg et al. 11, reported either slight or P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 79 Table 1: KIC for V notched cylindrical samples Tabela 1: KIC valjastih preizku{ancev z V-zarezo Specimens Testing conditions KIC Notch strength Rockwellhardness Preheating and quenching Subzero treat- ment in liqu. N2 Tempering Mpa √m MPa HRc * 33−40 650/850 °C → 1000 °C/30 → N2 (5 bar) – 2-times 550 °C/1 h 12.9 ± 0.9/1.3 (11.2 at x0) 283.3 ± 19.8/29.0 57.5 ± 0.3/0.7 41−50 650/850 °C → 1000 °C/30 → N2 (5 bar) −196 °C/24 h 13.2 ± 2.4/1.9 (9.5 at x0)) 289.4 ± 51.6/42.0 56.4 ± 0.6/0.5 51−60 650/850 °C → 1050 °C/30 → N2 (5 bar) – 10.9 ± 1.5/1.2 (8.1 at x0)) 238.0 ± 33.8/26.2 61.3 ± 0.4/0.8 61−69 650/850 °C → 1050 °C/30 → N2 (5 bar) −196 °C/24 h 12.5 ± 1.4/1.9 (9.1 at x0)) 273.1 ± 31.2/40.9 58.8 ± 0.3/0.3 * One (1) Rockwell is added as a correction for rounded (cylindrical) samples Table 2: Bend strength of samples with square section Tabela 2: Upogibna trdnost preizku{ancev s kvadratnim prerezom Specimens Testing conditions Bend strength Rockwell hardness Preheating and quenching Subzero treatment in liqu. N2 Tempering MPa HRc 1−7 650/850 °C → 1000 °C/30 → N2 (5 bar) – 2-times 550 °C/1 h 3626 ± 186/259 56.3 ± 0.7/1.2 8−14 650/850 °C → 1000 °C/30 → N2 (5 bar) −196 oC/24 h 3587 ± 143/169 55.3 ± 0.8/1.0 15−21 650/850 °C → 1050 °C/30 → N2 5 bar – 3474 ± 357/514 59.9 ± 0.8/1.1 22−28 650/850 °C → 1050 °C/30 → N2 (5 bar) −196 °C/24 h 3352 ± 385/307 58.1 ± 0.5/0.4 more significant hardness increase of hardness after sub-zero treatment. However, in the mentioned papers, other ledeburitic steels were investigated, the sub-zero processing was carried out at different parameters and the tempering was shorter. It is assumed also, that two processes were happening during the deep freezing related to high stresses and strains; i.e. retained austenite to martensite transformation (hardness increase) and mechanical annealing (hardness decrease). Anyway, the reasons for the reduction in hardness requires more detailed TEM investigation. The three point bending strength (Table 2 and Figure 4a), decreases with the increasing austenitizing temperature. This was expeced because at increased austenitizing temperature grains coarse structural constituents are changed and the bending strength is decreased. Also after the sub-zero processing, the three point bending strength is decreased. Similar results were published also by other authors 10,11,14 that did not suggest a reliable explanation. The bending strength is lower because of the increased internal stresses after quenching and deep freezing. These stresses are relieved by the P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 80 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 Figure 3: Microstructure of the Vanadis 6 steel after the heat treatment: a) austenitizing 1000 °C + quenching + 2x tempering, b) austenitizing 1000 °C + quenching + sub-zero + 2x tempering, c) austenitizing 1050 °C + quenching + 2x tempering, d) austenitizing 1050 °C + quenching + sub-zero + 2x tempering, Slika 3: Mikrostruktura jekla po toplotni obdelavi: a) avstenitizacija pri 1000 °C + kaljenje + dvakratno popu{~anje, b) avstenitizacija pri 1000 °C + kaljenje + obdelava pod ni~lo + dvakratno popu{~anje, c) avstenitizacija pri 1050 °C + kaljenje + dvakratno popu{~anje, d) avstenitizacija pri 1050 °C + kaljenje + obdelava pod ni~lo + dvakratno popu{~anje. Figure 2: Microstructure of the investigated steel in as-received condition: a) LM micrograph and b) SEM micrograph Slika 2: Mikrostruktura jekla v dobavljenem stanju; a) opti~ni posnetek, b) SEM-posnetek tempering and this supports the conclusion that the sub-zero treatment decreases the material’s resistivity to crack initiation. A possible explanations could be in the increased surface roughness because of martensite transformation (lattice reorientation and deformation) and mechanical annealing (dislocation glide and annihilation). Further investigations should prove this explanation. On the other hand the results of fracture toughness (Table 1 and Figure 4b) show slightly opposite ten- dency. Normally, it can be expected that higher austenitizing temperature would lead to a lower fracture toughness, since it is function of the matrix hardness. The principal explanation would the higher hardness of the matrix after higher austenitizing temperature. Never- theless, the average fracture toughness after heat pro- cessing of Vanadis 6 steel was found to be slightly higher for sub-zero processed material. This is rather contra- dictory to the generally expected behaviour, as it is normally assumed that not only the resistance against initiation, but also the resistance against propagation of the crack should be lowered. However, it seems that the situation can be considered to be more complex, the more complete transformation of austenite into the martensite would be in favour of lower fracture toughness. On the other hand, the so-called effect of mechanical annealing (dislocation glide to the grain boundaries and a decrease of dislocation concentration inside the grains) can induce a resistance to the crack propagation resulting in the improvement of fracture toughness. It is noticed that former investigations have shown a linear correlation between so called weak spot (x0) and fracture toughness KIC. In the case of tool and high speed steels this linear correlation shows a relatively large statistical reliability (r2 > 0.7) and as the relevant mate- rial’s KIC value is used, the KIC value at x0 = 0 18–20 are very conservative. In the case of materials with larger KIC (for example hot work tool steels) this correlation is not so distinct and the mean value of all the determined KIC values is usually used. In Table 1 both approaches of KIC results are given. For the KIC = f(x0) approach a relatively good statistical reliability of results is obtained for three cases (r2 = 0.7–0.9), except for the first set of experi- ments (specimens 33–40, Table 1). The obtained KIC values at x0 = 0 seems also to be relatively too conser- vative, therefore mean values (arithmetic average) of KIC are used as relevant. The fracture surface of not sub-zero processed sample after three point bending test is in Figure 5a. The fracture was initiated on the tensile strained side of the specimen (left margin of the micrograph) and propagated downwards the material. Some propagation lines are visible in the micrograph. In the direction from the tensile side to the core of the sample the fracture surface becomes a slightly more marked topography. As reported 5, such fracture surface (typical for hard steels) is known as low energetic ductile because some energy is also spent for the plastic deformation. However, the plastic energy is relatively low and corresponds to a relatively flat surface relief of the surface. More detailed micro- graph, Figure 5b, shows that the fracture was propagated by two main mechanisms. The first is the cracking of coarser carbide particles (see Figure 5c) and the second is the de-cohesion at the carbide/matrix interfaces which can be attributed to different plasticity of these two microstructural constituents. In some places dimples are (tracks of extracted carbides) visible with in vicinity a plastically deformed matrix, Figure 5d. The fracture surface of sub-zero processed material is shown in Figure 6. On the overview, Figure 6a, also propagation lines on the tensile strained side are visible. Detail micrograph, Figure 6b, confirms that the mechanism of fracture propagation is very similar to that of the sample processed without sub-zero period. Both, cracking of carbide particles and de-cohesion on the matrix/particle interfaces are clearly visible in the micrograph. To explain the fracture propagation more precisely, the fracture surface of one specimen processed with a sub-zero period was examined to more in detail. Figure 7a shows a SEM micrograph at high magnification. P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 81 Figure 4: Graphical presentation of obtained experimental results for a) three-point bend strength and b) average fracture toughnes for different hardness Slika 4: Eksperimentalni rezultati za a) tri to~kovni preizkus upogibne trdnosti in b) `ilavosti loma pri razli~ni trdoti P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 82 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 Figure 6: Fracture surface of the steel Vanadis 6 processed with a sub-zero period by different magnification Slika 6: Povr{ina preloma jekla po obdelavi pod ni~lo; SEM-posnetka pri razli~ni pove~avi Figure 5: Fracture surface of the steel Vanadis 6 without a sub-zero treatment by different magnification Slika 5: Povr{ina preloma jekla brez obdelave pod ni~lo; EM-posnetki pri razli~nih pove~avah Three characteristic areas are marked. The first and second are carbide particles. The particle "1" was broken at fracture propagation. The particle "2" did not break and at the interface with the particle "1" decohesion is visible. The place "3" is a typical cleavage area in the matrix. EDX-mapping shows that the particle 1 is a chromium rich phase, probably of the M7C3-type. Its chemistry was estimated to 39 % Cr, 40 % Fe, 12 % V. The particle "2" contains more vanadium and less other elements – EDX fixed following chemistry: 13 % Fe, 61 % V, 9 % Cr. The cleavage region "3" can be referred as matrix with the chemical composition 89 % Fe, 6 % Cr, 1.7 % V. 4 CONCLUSIONS The microstructure of the as-received material consisted of the matrix and carbide particles uniformly distributed in the material. The maximal size of particles is around 3 µm, but there are also particles with a size below of 0.5 µm (pearlitic carbides) are found. After the heat-treatment the microstructure consists of the martensitic matrix, eutectic and part of secondary carbide particles. The amount of particles is smaller than in the as-received condition. For the heat processed material, the portion of carbide is slightly smaller for the specimens processed at higher austenitizing temperature. The heat-treatment leads to a substantial hardness increase. With increasing austenitizing temperature the hardness increases, while, the sub-zero processing leads to slight decrease in hardness. This is rather surprising and inconsistent with earler observations. In this moment it is not clear why the sub-zero processing lowers the hardness of the steel Vanadis 6 steel and the explanation requires a more detailed investigation. The three point bending strength decreases with the increasing austenitizing temperature. Also, with the P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 83 Figure 7: Fracture surface of the steel Vanadis 6 processed with a sub-zero period, fracture surface and EDS mapping for main alloying elements Slika 7: Povr{ina preloma jekla po obdelavi pod ni~lo; prelomna povr{ina in rasterski EDS-posnetek za osnovne legirne elemente application of sub-zero processing the three point ben- ding strength is lowered. The fracture toughness is lowered with increased austenitizing temperature. It is rather surprising that the sub-zero treatment increases the average fracture tough- ness slightly. Similarly to the hardness behaviour, this is not entirely clear and needs detailed TEM examination. The fracture surfaces of the sub-zero- and the not sub-zero processed samples have similar basic characte- ristics. Both surfaces shown clearly dimple morphology. In some areas details of crack propagation like breaking of carbide particles and decohesion at the phase inter- faces occur, also. ACKNOWLEDGEMENTS This work was done within the project ICDAM (INNOVATIVE CENTRE OF DIAGNOSTICS AND APPLICATION OF MATERIALS) and as a result of long-time scientific co-operation between IMT in Ljub- ljana and CTU in Prague. 5 REFERENCES 1 Geller, J. A.: Instrumenta¾nyje stali, 5. vyd., Metallurgija, Moskva, 1983 2 Grga~, P.: Neue Hütte, 12 (1989), 459 3 Fremunt, P., Krej~ík, J., Podrábský, T.: Nástrojové oceli, Dum techniky Brno, 1994 4 Spies, H.-J., Riese, A., Hoffmann, W.: Neue Hütte, 3 (1990) 3, 96 5 Hnilica, F., ^makal, J., Jur~i, P.: Materiali in Tehnologije, 38 (2004) 5, 263 6 Leskov{ek, V., Ule, B.: J. Mater. Proc. Technol. 82 (1998), 89 7 Berns, H., Fischer, A., Hönsch, W.: Härterei-Tech. Mitt., 45 (1990) 4, 217 8 Berns, H., Bröckmann, C., Weichert, D.: Engineering Fracture Mechanics, 58 (1997), 311 9 Olsson, L. R., Fischmeister, H. F. Powder Metalurgy, 1 (1978), 13 10 Berns, H.: HTM, 29 (1974) 4, 236 11 Kulmburg, A. et al.: HTM 47 (1992) 5, 318 12 Stratton, P. F.: In.: Proc. of the 1st Int. Conf. on Heat Treatment and Surf. Eng. of Tools and Dies, Pula, Croatia, 8.–11. 6. 2005, 11–19 13 Collins, D. N.: Heat Treatment of Metals 23 (1996), 2 14 Meng, F. et al.: ISIJ International, 34 (1994) 2, 205–210 15 Collins, D. N., Dormer, J.: Heat Treatment of Metals, 24 (1997) 3, 71–74. 16 [u{tar{i~, B., Leskov{ek, V., Jutri{a, G.: Materiali in Tehnologije, 37(2003), 369 17 Z. W. Shan et al: Mechanical annealing and source-limited defor- mation in sub-micrometre-diameter Ni crystals, Nature Materials, 7 (2008), 115 18 B. Ule, V. Leskov{ek, B. Tuma: Estimation of plain strain fracture toughness of AISI M2 steel from precracked round-bar specimens, Eng. fract. mech., 65(2000) 5, 559–572 19 V. Leskov{ek, B. Ule, B. Li{~i}: Relations between fracture tough- ness, hardness and microstructure of vacuum heat-treated high-speed steel, J. mater. process. technol., 127 (2002), 298–308 20 V. Leskov{ek: Modelling of high-speed steels fracture toughness, Mater. manuf. process., 24 (2009) 6, 603–609 P. JUR^I ET AL.: FRACTURE CHARACTERISTICS OF THE Cr-V LEDEBURITIC STEEL VANADIS 6 84 Materiali in tehnologije / Materials and technology 44 (2010) 2, 77–84 F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... SURFACE MODIFICATIONS OF MARAGING STEELS USED IN THE MANUFACTURE OF MOULDS AND DIES MODIFIKACIJA POVR[INE JEKLA MARAGING IN UPORABA PRI IZDELAVI KOKIL IN UTOPOV Franjo Cajner1, Darko Landek1, Vojteh Leskov{ek2 1University of Zagreb, Faculty of Mechanical Engineering and Naval Architecture, Ivana Lu~i}a 1, 10000 Zagreb, Croatia 2Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia franjo.cajner@sb.hr Prejem rokopisa – received: 2009-11-02; sprejem za objavo – accepted for publication: 2009-12-09 Low-carbon, high-alloy, precipitation hardening MARAGING steels have been developed in the early sixties of the last century as a high strength structural materials for application in aeronautical and missile engineering. Due to their excellent properties, such as: high hardenability, good toughness, and high resistance to thermal fatigue, and due to simple heat treatment (without protective atmosphere) with very small distortions, MARAGING steels are successfully used for the fabrication of moulds and dies. The main drawback of these steels is their relatively low wear resistance, particularly if the die is subjected to extensive wear during service. The paper presents an overview of existing MARAGING steels used in mould manufacturing and gives their general properties. Also, the results of the applied modification and coating processes are presented with a special focus on the increase in wear resistance. Key words: MARAGING steels, thermo-chemical treatment, PVD coating, wear resistance Malooglji~na, visokolegirana in izlo~evalno utrjena jekla MARAGING so bila razvita v zgodnjih {estdesetih letih prej{njega stoletja kot zlitine z visoko trdnostjo za uporabo v letalski in raketni tehniki. Zaradi izvrstnih lastnosti, kot so velika kaljivost, dobra `ilavost, velika odpornost proti toplotni utrujenosti in enostavna toplotna obdelava (brez varovalne atmosfere) z majhnim krivljenjem, omogo~ajo u~inkovito uporabo teh jekel tudi za izdelavo kokil in utopov. Njihova pomanjkljivost je majhna odpornost proti obrabi, {e posebej, ~e je utopno orodje izpostavljeno veliki obrabi. V ~lanku je pregled jekel vrst MARAGING, ki se uporabljajo pri izdelavi utopov, in njihovih lastnosti. Predstavljeni so tudi rezultati uporabljenih metod za obdelavo in prekritje povr{ine s poudarkom na pove~ani odpornosti proti obrabi. Klju~ne besede: jekla MARAGING, termo-kemi~na oddelava, PVD-prekritja, odpornost proti obrabi 1 INTRODUCTION Low-carbon, high alloy MARAGING steels belong to the Fe-Ni-Co alloying system, using Mo, Ti, and Al as alloying elements. These steels were developed in the early sixties of the last century as high strength structural materials intended for application in aeronautical and aerospace engineering (e.g. for propeller drive shafts, pilot seat frames, liquid fuel tanks, armour plates, etc.). By modifying their composition with the addition of Cr (9–3 %), Fe-Ni-Cr or Fe-Ni-Co-Cr alloying systems were created, MARAGING steels belonging to these systems are suitable for the application in highly corrosive environments 2. Later on, MARAGING steels started to be used for the manufacturing of tools, where are superior to other tool materials due to their superior properties 1–3: • High toughness and high fracture toughness com- bined with very high strength; • High resistance to thermal fatigue; • Protective atmosphere in heat treatment is not required, i.e. in solution annealing and precipitation hardening – aging, as there is no risk of decarburiza- tion and oxidation of the surface. • Hardenability, even with the largest tool sizes, is achieved by slow cooling from the temperature of solution annealing; therefore, the risks of distortions and the occurrence of cracks caused by the difference in temperatures in the tool cross-section are significantly reduced. • Good weldability; • Good electro-erosion machinability; • Good chip machinability and/or cold deformability after quenching enable the tool manufacturing to the end dimensions before the final heat treatment (aging) is completed. • In solution annealing and in cooling, the expected shortening is approximately 0.1 %, and in the final heat treatment (aging), there are practically no defor- mations; therefore, these heat treatment processes can be considered as processes causing no distortion. • Compared to the conventional heat treatment of tool steels, the heat treatment of MARAGING steels (solution annealing and aging) is much simpler. The main disadvantage of MARAGING steels compared to high alloy tool steels is their relatively low hardness (HRC = 50–57 at the most), and consequently the insufficient resistance to wear. Due to their lower hardness, MARAGING steels are not suitable for the manufacture of cutting tools, but are suitable for the manufacture of moulds and dies (for die-casting and for Materiali in tehnologije / Materials and technology 44 (2010) 1, 85–91 85 UDK 669.14.018.298:539.4 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(2)85(2010) polymer processing), of forming tools, etc. 1,2. A relati- vely high price of MARAGING steels (they are several times more expensive than the high alloy tool steels produced by standard methods) cannot be taken as a major drawback in modern machine tool industry since tool steels produced by powder metal forming have a similar price. In references 3–5, the application of nitriding and/or nitrocarburising is usually recommended to improve the wear resistance. The investigation of the application of other thermo-chemical treatments (carburizing, boriding) and of vapour deposition coating processes 6–10 seems to be a natural choice. Results found in literature and results of our research on the possibility of improving the wear resistance of MARAGING steels are presented in this work. 2 PROPERTIES OF MARAGING TOOL STEELS Iron-nickel phase diagram and the influence of the content of nickel on the hardness in the solution- annealed steels (Figure 1) can contribute to the under- standing of the hardening mechanism of MARAGING steels. Figure 1 a, shows that the alloy with approxi- mately 18 % Ni will be in the single-phase austenite area above the temperature of 650 °C (generally the temperature of austenization-homogenization is approxi- mately of 820 °C). With this content of nickel, the alloy exhibits the highest strength after cooling (Figure 1b). As the solution annealed MARAGING steel is cooled from the temperature of annealing to approximately 250 °C (slow cooling in the air is sufficient), the low-carbon nickel-martensite, which is actually an oversaturated solution of Co, Mo, Ti and Al in the alloy Fe-Ni-(Co), starts to form. The transformation of austenite into nickel-martensite is completed at a temperature of approximately 200 °C; and at this temperature, the microstructure should be without residual austenite. Nickel-martensite (with a hardness of approx. HV 300) can be easily machined by the chips removal and by deformation machining and can even be welded, if necessary. The solution annealing of MARAGING steels is carried out in steel mills, and the manufacturer of a machine part or a tool carries out only the final heat treatment (aging) after the machining to the specified dimensions. The process of aging is performed in furnaces without protective atmosphere and at relatively low temperatures (approx. 500 °C). In the slow cooling process that follows, small distortions occur. In the age hardening process, a large number of finely dispersed precipitates of intermetallic phases occur (e.g. Ni3Al, Ni3Ti, Fe2Mo, FeCr, and Fe7Mo6) in nickel martensite, which results in an increase in steel hardness from HV = 280–380 to HV = 500–650. Also, the tensile strength of the steel is increased to 4200 N/mm2, and deformability and fracture toughness are rather good 1–4. Figure 2 gives a flow diagram of the heat treatment process of MARAGING steels. Table 1 lists several grades of MARAGING tool steels divided into three groups according to their resistance to elevated temperatures and corrosion 2: a) steels suitable for hot working at temperatures of up to 425 °C: X 3 NiCoMo 18 8 5, X 3 NiCoMo 18 9 5, X 2 NiCoMoTi 18 12 4 b) steel suitable for hot working at temperatures of up to 600 °C: X 2 NiCoMo 12 8 8 c) steels with superior corrosion stability: X 1 CrNiCoMo 9 10 3, X 1 CrNiCoMo 13 8 5, X 2 CrNiCoMo 12 8 5 F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... 86 Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 Figure 2: A flow diagram of the heat treatment process of MARA- GING steels 2 Slika 2: Diagram poteka procesa toplotne obdelave jekel MARA- GING jekel 2 Figure 1: a) Real phase diagram Fe-Ni 11, b) Quenching hardness of Fe-Ni alloys in dependence of the percentage of nickel (with the addition of Co, Mo, Ti, and Al) 12 Slika 1: a) Realni fazni diagram Fe-Ni 11, b) Kalilna trdota zlitin Fe-Ni v odvisnosti od vsebnosti niklja (z dodatkom Co, Mo, Ti in Al) 12 MARAGING steels suitable for working at tem- peratures of up to 425 °C have a rather low resistance to high temperatures since austenite, which decreases the hardness and strength in operating conditions, is recre- ated already at the temperature of 500 °C 13. Therefore, these steels are used for tools used in cold and lower temperature operating conditions, such as. 2: • moulds for polymer processing • dies for the die casting of silumin or zinc alloys • punches for the extrusion of lead cable sheaths • tools for compression-extrusion of Al-alloys • matrices and punches for the cold forging of bodies and heads of bolts. The MARAGING steel suitable for hot working at temperatures of up to 600 °C (X 2 NiCoMo 12 8 8) has a lower content of Ni than the steels with 18 % Ni. Therefore, it has to be solution annealed at a tempera- tures of approximately 900 °C (Figure 1) and also precipitation hardened (aged) at temperatures of up to 625 °C. The increased content of Mo (from 5 % to 8 %) has also a beneficial influence on the increase in hard- ness during the aging process by additional precipitation of Ni3Mo. Due to the improved resistance to thermal fatigue, this steel grade can be recommended for the manufacture of mould components for the die casting of aluminium alloys 2. MARAGING steels with improved corrosion sta- bility have a lower content of nickel (from 8 % to 10 %) and cobalt Co (from 3 % to 5%) and contain between 9 % to 13 % of chromium (as an active corrosion resistant ingredient, chrome contributes to the improved corrosion resistance). The addition of chromium decreases the hardness, as well as the strength (to approximately 1600 N/mm2) and the yield stress (to approx. 1500 N/mm2). It decreases also, the temperatures of the start (Ms) and the finish (Mf) of the formation of nickel-martensite. There- fore, with these steels, retained austenite may occur after quenching and a need for deep cooling may arise. MARAGING steels high-alloyed by chrome can be recommended for the manufacture of mould components for the processing of highly corrosive polymers. 3 IMPROVED WEAR RESISTANCE In reference, results of research into the feasibility of the application of various surface treatment procedures (nitriding, carburizing, boriding, and PVD coating) in order to improve the wear resistance of MARAGING steels are found 3,6,7,9,10,11,12,13. Our investigations 6,9 on the MARAGING steel 14 10 5 (Table 2) included the following surface treatment procedures: nitriding in Tenifer salt bath, nitriding in plasma gases, as well as carburizing and boriding according to the parameters listed in Table III. In addition, a possible application of the duplex treatment of plasma nitriding and of physical vapour deposition (PVD) procedure has been investi- gated also. In Table 3 the performed researches are summarised and all heat treatment parameters applied are listed. The effect of processes of modification and coating on the properties of MARAGING steel 14 10 5 has been estimated with microstructure, hardness distribution on F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 87 Table 1: MARAGING tool steels 2 Tabela 1: Jekla MARAGING 2 R es is ta nc e to te m - pe ra tu re / co rr os io n Steel grade Heat treatment Mechanical properties H om og en iz at io n te m pe ra tu re , o C A gi ng te m pe ra tu re , o C / /A gi ng ti m e, h after solution annealing (homogenization) after aging R m /M Pa R p0 .2 /M Pa m ax . H V m in . A 5/ % m in . K V 20 °C /J R m /M Pa R p0 .2 /M Pa m ax . H V m in . A 5/ % m in . K V 20 °C <425 °C X3 NiCoMo 18 8 5 820 480 / 3 1130 830 350 15 60 1920 1720 500 8 20 X3 NiCoMo 18 9 5 820 480 / 3 1130 830 350 15 1960 1910 570 7 30 X2 NiCoMoTi 18 12 4 820 500 / 6 1130 830 350 15 2350 2260 590 6 10 <600 °C X2 NiCoMoTi 12 8 8 900 550/ 2 1980 1800 560 Corro- sion resi- stance X1 CrNiCoMo 9 10 3 835 480 / 6 1500 1400 12 60 X1 CrNiCoMo 13 8 5 820 480 / 6 1700 1500 11 50 X2 CrNiCoMo 12 8 5 880 480 / 6 1000 700 1870 1650 10 40 Table 2: Chemical composition of MARAGING steel 14 10 5 Tabela 2: Kemi~na sestava jekla MARAGING 14 10 5 w/% C Si Mn P S Cr Mo Ni V Al Cu Ti Nb N B Co 0.01 0.88 0.087 0.003 0.002 0.11 4.81 13.60 0.024 0.14 0.08 0.17 0.10 0.013 0.002 9.50 the cross-section and tribological tests (abrasive, erosive and adhesive wear). Figures 3 to 8 present the results of these tests. The testing of resistance to abrasive wear was carried out by the dry sand/rubber wheel tests (Ottawa 50/70 quartz sand) at a compressive force of 45 N (Figu- re 4). The same sand type was used in the tests of erosive wear with sand grains colliding with the test sample surface at an angle of incidence of 90° (Figure 6). Tests of abrasive wear were carried out by a friction ring made of hardened steel at the compressive force of 100 N. During testing, values of friction coefficient were determined (Figure 8). Details on the conducted tests are given in 9,10. 3.1 Carburization of MARAGING steels The carburization of MARAGING steels was carried out at the temperatures of solution annealing. After the cooling process, high-carbon martensite with increased hardness (and a probable presence of residual/retained austenite) is obtained in the boundary layer. The total depth of carburized layer on the MARAGING steel 14 10 5 was approx. 0.8 mm. The subsequent aging improved the core hardness, and reduced the hardness of the carburized layer (Figure 3). The carburization of MARAGING steels improved their resistance to abrasive wear compared with the solution annealed and aged condition or with that of solution annealed and nitrided (Figure 4). The carburization of MARAGING steels also has some adverse effects. The formation of a hard martensite layer (with hardness of approx. HV 700) is accompanied with additional distortions and by a more difficult chip machining. Thus, one of major advantages of MARAGING steels, i.e. the heat treatment without distortions, is lost. 3.2 Boriding of MARAGING steels The boriding of MARAGING steel 14 10 5 by using the standard means and parameters results in a boride F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... 88 Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 Figure 5: The microstructure and hardness distribution in the edge layer of borided test samples made of MARAGING steel 14-10-5 10 Slika 5: Mikrostruktura in porazdelitev trdote v robni plasti boriranih preizku{ancev iz jekla MARAGING 14 10 5 10 Table 3: Heat treatment of test samples made of MARAGING steel 14 10 5 Tabela 3: Toplotna obdelava preizku{ancev iz jekla MARAGING 14 10 5 Heat treatment parameters S ol ut io n an ne al ed at 82 0 °C /1 h/ ai r Aged at 500 °C/4h Nitrocarburized by the TENIFER procedure at 580 °C/4h Ion-nitrided at 500 °C/72h Ion-nitrided at 500 °C/20h Ion-nitrided at 500 °C/20h + PVD coated at 450 °C/70 min Carburized in the Degussa KG 6 granulate at 900 °C/4h + aged at 500 °C/4h Borided in the EKABOR 2 (Degussa) powder at 900 °C/4h + aged at 500 °C/4h Figure 3: Hardness distribution in the edge layer of carburized test samples made of MARAGING steel 14-10-5 Slika 3: Porazdelitev trdote ob povr{ini naooglji~ene plasti pri preizku{ancu iz jekla MARAGING 14 10 5 Figure 4: The abrasive wear mass loss of test samples made of MARAGING steels 14-10-5 with and without surface modifications compared with the wear of samples made of high alloy tool steel X155CrVMo12 1 9,10 Slika 4: Izguba mase zaradi abrazivne obrabe preizku{ancev iz jekla MARAGING 14 10 5 z modifikacijo povr{ine in brez nje v primerjavi s preizku{anci iz mo~no legiranega orodnega jekla X155CrVMo 12 1 9,10 layer with a thickness of approx. 50 µm and with high hardness (above HV0.5 1000) (Figure 5). This layer showed a high resistance to erosive wear and an improved resistance to abrasive wear, even higher than that of highly wear resistant ledeburite steel for cold working (Figure 4 and 6). Along with these obvious advantages, one should also notice the problems in the application of the boriding of MARAGING steels, such as the occurrence of a soft zone below the boride layer, the presence of transverse cracks in the boride layer, and the risk of exceeding dimensional tolerances. Due to the achieved high hardness after boriding (prior to aging), it is not possible to correct dimensions and thus the property of "distortionless" heat treatment is lost. 3.3 Nitriding of MARAGING steels with and without a subsequent PVD coating procedure An increased wear resistance of MARAGING steels can be successfully obtained by plasma nitriding 6,9,13,14 at temperatures of approximately 500 °C, which are in the range of aging temperatures of these steels. The application of solution annealing and plasma nitriding results in an improved resistance to abrasive and erosive wear compared with the application of standard solution annealing and aging (Figure 4 and 6). Also, all the other advantages of the application of MARAGING steels pointed out before are retained. Research results reported in 13 show that the corrosion resistance of MARAGING steels exposed to salt mist (salt-spray test with exposure of the test sample to the aqueous solution of sodium chloride at a concentration of 5 % NaCl) is increased by three times if plasma nitriding is applied compared to the solution annealed and aged condition. This result is expected in all the cases in which a zone of compounds (Fe nitrides) is formed on the MARAGING steel surface. Here, the formation of a single-phase layer of Fe4N nitrides is recommended since, according to 14, these layers exhibit a good combination of high hardness and ductility. The application of nitriding and nitrocarburizing in gases, as well as nitrocarburizing in the Tenifer salt bath results in the overaging of the core due to a higher tem- perature of the treatment (550–600 °C). This pheno- menon is reflected in hardness decrease of the base metal and in the hardness of the nitrided layer (Figure 7). The overaging of the core can have an influence on the decrease in the resistance to abrasive and erosive wear (Figure 6) 9,10. The duplex treatment of plasma nitriding and of the physical vapour deposition (PVD) procedure seems to be of a particular interest regarding the improvement in wear resistance. In the case of the tested MARAGING steel 14 10 5, the thickness of the PVD TiN coating was 1.8 µm (Figure 8). In the plasma nitriding treatment it is possible to vary the parameters of the procedure in order F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 89 Figure 8: Microstructure and hardness distribution in the edge layer of test samples made of MARAGING steel 14-10-5, nitrided in plasma and coated by a PVD TiN coating 10 Slika 8: Mikrostruktura in porazdelitev trdote v robni plasti preizku- {ancev iz jekla MARAGING 14 10 5, ki so bili nitrirani v plazmi in pokriti s plastjo PVD TiN 10 Figure 7: Microstructure and hardness distribution in the edge layer of test samples of MARAGING steel 14-10-5, nitrided in plasma and nitrocarburized in the TENIFER salt bath 10 Slika 7: Mikrostruktura in porazdelitev trdote v robni plasti preizku- {ancev iz jekla MARAGING 14 10 5, ki so bili nitrirani v plazmi in karbonitrirani v solni kopeli TENIFER 10 Figure 6: Mass loss in erosive wear tests carried out on test samples made of MARAGING steel 14-10-5 with and without surface modifications 10 Slika 6: Izguba mase pri preizkusih erozivne obrabe preizku{ancev iz jekla MARAGING 14 10 5 z modifikacijo povr{ine in brez nje 10 to obtain a boundary layer with or without a zone of compounds, and the PVD treatment produces a hard coating with the properties of a high resistance to abrasive, erosive and adhesive wear, together with an extremely low friction coefficient (Figure 4 and 9). 3.4 Repair of worn tools from MARAGING steels Low carbon MARAGING steels can be welded, which is a good basis for the repair of worn tool surfaces and those with thermal fatigue cracks. A description of a successful welding procedure and the closure of surface microcracks in the MARAGING steel X 2 NiCoMoTi 12 8 8 with laser remelting (using a diode laser) without using a filler metal is given in 15. Diode lasers with beam power of 1 kW to 2 kW are compact and easily con- trolled sources of light beams suitable for the repair and maintenance of tool surfaces, especially of large and expensive tools, such as moulds for die casting, dies, tools for pressing polymers, etc. For each repair case and type of MARAGING steel, one has to determine appro- priate operating parameters, such as the power output and shift of the laser beam and its focusing points, the depth of remelting, and other parameters. Diode lasers can find their application not only in tool repair procedures for remelting tool surfaces, but also in the manufacture of tools made of MARAGING steels in order to improve the wear resistance of the surface by applying coating and alloying procedures. 4 CONCLUSION Good mechanical and technological properties and the heat treatment with almost no distortions make MARAGING tool steels suitable for the manufacture of moulds used in polymer processing, of dies for die- casting, drop-hammer dies, punches, matrices, and other non-cutting tools. In addition to the basic group of MARAGING steels with 18 % Ni, two special groups of steels resistant to temperatures of up to 600 °C and those used for tools exposed to corrosion have been developed. In the investigation, it is confirmed that the rather poor wear resistance of MARAGING steels can be improved with thermo-chemical heat treatments, as nitriding, nitrocarburizing, boriding and carburizing. The carburizing process improve the resistance to abrasion, however it cannot be recommended as a heat treatment process for the surface modification of these steels because the risk of exceeding the required dimensional tolerances and of shape changes (similarly as in the case of boriding) and the chip machining to be conducted prior to aging is made more difficult (thus, a major advantage has been lost). The process of boriding significantly improves the resistance to abrasive and erosive wear, but transverse cracks occurring in the boride layer, a change in dimen- sions and the loss of "distortionless" heat treatment make this process, for the time being, unsuitable for the manufacture of tools. Nitrocarburizing in gases and in salt baths produces good-quality layers, but since the process is carried out at very high temperatures (550–600 °C), this causes a significant decrease in the core hardness of standard MARAGING steels. This fact should be taken into account when the application of these processes is considered. Plasma nitriding is the most suitable heat treatment process for the processing of MARAGING steels because the temperature of nitriding can be lower than (or equal to) the optimum temperature of ageing. Therefore, the process of ageing can be carried out during nitriding. As far as an improved resistance to abrasive, erosive and adhesive wear is concerned, the plasma nitriding of MARAGING steels proves to be more effective than solution annealing and aging, but still less effective than carburizing and boriding. It is to be expected that further improvements and the optimi- zation of plasma nitriding parameters will produce layers with even a higher resistance to wear. In addition to the already mentioned advantages of the procedure, plasma nitriding results in some other favourable properties, such as an improved resistance to thermal fatigue, improved corrosion stability, and improved resistance to adhesion of the processed molten mixture. All this contributes to a conclusion that this procedure can be recommended in the manufacture of moulds made of MARAGING steels used for the processing of polymers and metals Significant improvements in the resistance to erosion and adhesion of MARAGING steel 14 10 5 achieved by plasma nitriding followed by coating the metal by a TiN layer speak in favour of the application of this combi- nation of processes for the manufacture of moulds that F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... 90 Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 Figure 9: Friction coefficient determined in the adhesive wear testing of test samples made of MARAGING steel 14-10-5 with and without modifications of the surface 10 Slika 9: Koeficient trenja pri preizkusu adhezivne obrabe preizku{ancev iz jekla MARAGING 14 10 5 z modifikacijo povr{ine in brez nje 10 have to meet the requirement of an increased wear resistance. It is expected that laser beams used for the modification and coating of surfaces will find their application in the manufacture of tools and in the repair of worn tool surfaces made of MARAGING steels, particularly if compact and portable diode lasers are used. 5 LITERATURE 1 Novosel M., Cajner F., Krumes D.: Tool Materials, Faculty of Mechanical Engineering in Slavonski Brod, Slavonski Brod, Croatia, 1996 (book published on croatian) 2 Cajner F., Krumes D.: MARAGING steels – new tool materials, Proceedings of symposium Modern technologies of heat treatment of steels, Zagreb, Croatia, 11. 6. 1998, 69–80 3 ASM Handbook, Part 4, ASM International, Metals Park, Ohio, 1994 4 Heberling J. M.: MARAGING steel: A dependable alloy, Heat Treating, (1993), Sept., 22–24 5 Decker, R. F.: MARAGING steels: Getting better with age, AMP Metal Progress (1988) 6, 45–50 6 Novosel M:, Cajner F.: Applicability of thermochemical treatment for MARAGING steels and use for manufacturing of the moulds for polymer processing, Proceedings of the 13th days of Polymer and Rubber Society, Zagreb, Croatia, 1995, 3–12 7 Cajner F., Leskov{ek V.: Influence of thermochemical treatment of MARAGING steel on wear resistance, Proceeding of international conference MATRIB’99, Trogir, Croatia, 1999, 15–22 8 Kladari}, I., Cajner F., Krumes D.: Optimization of Parameters for Ageing of MARAGING Steel, Proceedings of the 8th Seminar of the International Federation for Heat Treatment and Surface Engi- neering, Dubrovnik-Cavtat, Croatia, 12–14. September 2001, 111–117 9 Cajner F.; Landek D. Kladari} I.: Improvement of wear resistance of MARAGING steel by applying thermochemical processes, Proceedings of the 1st International Conference on Material & Tribology 2002., Dublin, Ireland, 2002 10 Cajner F., Landek, D. [oli} S., Cajner H.: Effects of thermochemical treatments on properties of MARAGING steels, Surface Engi- neering, 22 (2006) 6, 468–471 11 Johs F. W., Pumphrey W. I.: Free energy and metastabile in the iron-nickel and iron-manganese systems, Journal Iron and Steel Institute, 163 (1949), 121–128 12 Gold{tejn M. J., Gra~ev S. V., Veksler Y. G.: Specijalnie stali, Metalurgija, Moskva, 1985 13 Shetty, K., Kumar, S., Raghothama, R.: Effect of ion nitriding on the microstructure and properties of Maraging steel (250 Grade), Surface & Coatings Technology 203 (2009), 1530–1536 14 Yan, M.F., Wu, J. Y., Liu, R., L.: Plasticity and ab initio characteri- zations on Fe4N produced on the surface of nanocrystallized 18Ni-maraging steel plasma nitrided at lower temperature, Applied Surface Science 255 (2009), 8902–8906 15 Grum, J., Slabe, J. M.: Effect of laser-remelting of surface cracks on microstructure and residual stresses in 12Ni maraging steel, Applied Surface Science 252 (2006), 4486–4492 F. CAJNER ET AL.: SURFACE MODIFICATIONS OF MARAGING STEELS ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 85–91 91 K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE OF MASSIVE DUCTILE CAST-IRON CASTING NUMERI^NA MODELA STRUKTURE STRJEVANJA MASIVNEGA DUKTILNEGA @ELEZOVEGA ULITKA Karel Stransky1, Jana Dobrovska2, Frantisek Kavicka1, Vasilij Gontarev3, Bohumil Sekanina1, Josef Stetina1 1Faculty of Mechanical Engineering, Brno University of Technology,Technicka 2, 616 69 Brno, Czech Republic 2VSB – Technical University of Ostrava, Tr. 17. listopadu, 708 33 Ostrava, Czech Republic 3University of Ljubljana, A{ker~eva 12, 1000 Ljubljana, Slovenia stransky@fme.vutbr.cz Prejem rokopisa – received: 2009-07-02; sprejem za objavo – accepted for publication: 2009-11-26 An original three-dimensional (3D) model of solidification is used to describe the process of solidification and cooling of massive (500 × 1000 × 500) mm cast-iron sand moulds castings. The calculated mode of the kinetics of the temperature field of the casting was verified during casting with temperature measurements in selected points. The sizes and positions (xi, yi, zi, where i = 1, 2, 3 is the number of samples taken) of the experimental samples are exactly defined and corresponding with the decreasing rate of solidification. The experimental samples – 15 mm in diameter and 12 mm high – were metallographically analysed and also in terms of heterogeneity of chemical composition. The coordinates xi, yi, zi characterise approximately – within an accuracy of ±5 mm – the centres of the samples. Successively, the local solidification time Q (i.e. the time the specified position of the casting, defined by the coordinates xi, yi, zi, remained within the temperature range between the liquid and solid curves) is also calculated using the 3 D model. The following dependences are later determined according to experimental and calculated data: the average size of graphite spheroids rg and of graphite cells Rb as well as the average distances among the particles of graphite Lg – always as function of the local solidification time Q [xi, yi, zi]. Furthermore, it was founded that the given basic characteristics of the structure of the cast iron rg, Rb and Lg are directly proportional to the logarithm of the local solidification time. The original spatial model of solidification can therefore be used therefore as first approximation for the assessment of the casting structure of massive cast iron parts. This paper creates the starting point for the estimation of the local mechanical properties and fracture behaviour of massive ductile cast iron castings. Keywords: ductile cast-iron, solidification time, segregation, structural characteristics, model Za opis procesa strjevanja in ohlajanja masivnih `elezovih ulitkov velikosti (500 × 1000 × 500) mm smo uporabili osnovni tridimenzionalni (3 D) model strjevanja. Ulitki so bili uliti v pe{~ene forme. Izra~unan kineti~ni model temperaturnega polja ulitka je bil preverjen med ulivanjem z meritvami temperature v izbranih to~kah. Velikosti in lege (xi, yi, zi, kjer je i = 1, 2, 3 {tevilo vzetih vzorcev) poskusnih vzorcev so bile to~no dolo~ene in so v skladu s padajo~o hitrostjo strjevanja. Poskusni vzorci premera 15 mm in vi{ine 12 mm so bili metalografsko analizirani in preverjena je bila tudi heterogenost kemijske sestave. Koordinate xi, yi, zi ozna~ujejo sredino vzorcev s pribli`no z natan~nostjo ±5 mm. Lokalni ~as strjevanja Q (t. j. ~as dolo~enega polo`aja v vzorcu, definiranega s koordinatami xi, yi, zi ostaja v temperaturnem obmo~ju med likvidusno in solidusno krivuljo) je bil tudi izra~unan z uporabo 3D-modela. Kasneje so bile dolo~ene naslednje odvisnosti glede na eksperimentalne in izra~unane podatke: srednja velikost kroglastega grafita rg, grafitnih celic Rb in srednje razdalje med delci grafita Lg – vedno kot funkcije lokalnega ~asa strjevanja Q [xi, yi, zi]. Nadalje je bilo ugotovljeno, da so dane zna~ilnosti strukture litega `eleza rg, Rb in Lg direktno sorazmerne z logaritmom lokalnega ~asa strjevanja. Osnovni prostorski model strjevanja lahko tako uporabimo v prvem pribli`ku za ugotovitev lite strukture masivnih `elezovih ulitkov. Ta ~lanek omogo~a za~etno stopnjo ocene lokalnih mehanskih lastnosti in vedenja pri zlomih masivnih duktilnih `elezovih ulitkov. Klju~ne besede: siva litina s kroglastim grafitom, ~as strjevanja, izcejanje, zna~ilnosti strukture, modeliranje 1 INTRODUCTION The problem of optimisation of properties and production technology for the casting of massive ductile cast-iron (spheroidal graphite) castings had been investigated into in the past few years and, besides an extensive number of publications both nationally as well as internationally, the results of the investigations have been published in the final report of this investigation 1. During the investigations, the centre of focus were not only the purely practical questions relating to metallurgy and foundry technology, but mainly the verification of the possibility of applying two original models – the 3 D model of transient solidification and cooling of a massive cast-iron casting and the model of chemical and struc- tural heterogeneity. Both models have only been applied to describing the temperature field, the control of solidi- fication and the cooling of continually cast steel slabs, to the descriptions of their chemical heterogeneity and to determining the basic characteristics of their micro- structure. The model of chemical and structural hetero- geneity seems to be a suitable partner of the 3D model of the transient temperature field 2. The original model and also an original application of the software ANSYS were one of the outcomes of this research. In this combination, it is possible to optimise the technology of casting cast-iron parts and successive Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 93 UDK 669.1.621.74 ISSN 1580-2949 Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 44(2)93(2010) cooling in order to achieve the most convenient micro- structure. This includes a microstructure of spheroids of graphite, preferably with as high a density as possible and spread throughout the casting as evenly as possible, with a minimal ratio of particles of graphite marked as degenerated shapes. The appropriate density of the spheroids of graphite is simultaneously one of the conditions of good mechani- cal properties of castings of ductile cast-iron, especially good ductility and contraction while maintaining good strength properties – yield point and tensile strength. The discussed 3D model establishes a system via which it is possible to pre-simulate the method of eutectic crystallisation of the melt of ductile cast-iron. With the liquidus and solidus known, it is also possible – in selected parts of the casting (that would be suitably discretised) – to determine the time the temperature of the relevant part of the metal remained between the liquidus and solidus. This region is characterised by the coexistence of solid and liquid states – the so-called šmushy zone’. In describing the solidification of steel, the time for which the metal remains at a temperature between the liquidus and solidus is called the local solidification time, and the volume of metal corresponding to this time determined by the sizes of the dendrites 3,4. When describing the solidification of cast-iron, including cast-iron with spheroidal graphite, the term šlocal solidification time’ has not been used, even despite the fact that the eutectic crystallisation of grey cast-iron always runs within a certain temperature interval and naturally even a time solidification interval dependent on time. 2 AIMS AND METHODOLOGY The massive experimental cast-iron castings, produ- ced within the research 5, had the following dimensions: width × length × height = (500 × 1000 × 500) mm. The verifying numerical calculation of the local solidification times /s – conducted according to the 3D model proved that, along the height, width and length of these massive castings, there are various points with differences in the solidification time of up to two orders. The aim was to verify the extent to which the revealed differences in the local solidification time affect the following parameters: a) The average size of spheroidal graphite particles; b) The average density of spheroidal graphite particles; c) The average dimensions of graphite cells, and d) The chemical heterogeneity of elements in the cross- sections of individual graphite cells. The relationships – among the given four parameters and the corresponding local solidification time – were determined in the series of samples that had been selec- ted from defined positions of the massive casting. 2.1 Experimental cast-iron casting and selection of samples The experimental casting was selected from a series of three castings and it was marked as casting No.1 5. The bottom part of its sand mould was lined with (a total number of) 15 cylindrical chills of a diameter of 150 mm and a height of 200 mm. The upper part of the mould was not lined with any chills. The average chemical composition of the cast-iron before casting is given in Table 1. A (500 × 500 × 40) mm plate had been mechanically cut out of the middle of the length by two parallel trans- versal cuts. Then, further samples were taken from exactly defined points and tested in terms of their struc- tural parameters and chemical heterogeneity. Samples in the form of testing test-samples for ductility testing, with threaded ends, were taken from the bottom part of the casting (A), from the middle part (C) and from the upper part (G). The 15 mm in diameter and 12 mm high cylindrical samples served the actual measurements in order to determine the structural parameters and che- mical heterogeneity. K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... 94 Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 Table 1: Chemical composition of ductile cast-iron Tabela 1: Kemijska sestava duktilne `elezove litine Element C Mn Si P S Ti Al Cr Ni Mg w/% 3.75 0.12 2.15 0.039 0.004 0.01 0.013 0.07 0.03 0.045 Table 2: Measured and calculated structural parameters and the coordinates x, y, z of the measured samples Tabela 2: Merjeni in izra~unani strukturni parametri ter koordinate x, y, z merjenih vzorcev Sample rg/ m Rb/ m Lg/ m rghm/ m Rbhm/ m Lghm/ m x/ mm y/ mm z/ mm ls / s Ihet I s A 27.6 ± 3.6 82.8 165.6 28 83 110 190 50 507.5 48 0.952 3.34 C 36.4 ± 10.4 103.9 207.7 36 104 136 190 210 507.5 2509 1.091 4.52 G 38.9 ± 12.5 109.3 218.6 39 109 140 190 450 507.5 4542 0.916 3.98 Note: rg, Rb, Lg – metallographic analysis (measured), rghm, Rbhm, Lghm – chemical micro-heterogeneity (selected for analysis), Lghm  2Rb – 2rg, ls – local solidification time, Ihet – arithmetic mean of heterogeneity index, Is – arithmetic mean of segregation index of ten measured elements In the points of the defined positions of the samples prepared in this way, the quantitative metallographic analysis was used to establish the structural parameters of cast-iron 6, the in-line point analysis to establish the chemical composition of elements 7 and numerical calculation using the 3D model to establish the local solidification time 1. 2.2 Quantitative metallographic analysis Quantitative analysis of the basic micro-structural parameters in the samples, i.e. the radius of the spheroids of graphite – rg, the distances between the particles of graphite – Lg and the radius of the graphite cells – Rb had been the subject of a special study 4. The measurement of the size parameters of the graphite had been conducted on the Olympus CUE4 image analyzer under standard conditions, i.e. with a magnification of 100-times and on each sample a total number of 49 views were evaluated. The measurement results are given in Table 2. 2.3 Chemical heterogeneity of samples The concentration of elements in each of the samples was measured between two particles of spheroidal graphite. The analyzed region in the sample structure had been selected in order for the structural parameters of graphite (rghm, Rbhm, Lghm within the analyzed region) to approach the average parameters of graphite within the sample (rg, Rb, Lg) measured using quantitative metallo- graphic analysis. The differences of the average values of the parameters in the structures of samples and of para- meters selected for analysis of their chemical micro- heterogeneity of elements are based on the comparison of values in Table 2. Then differences occur only between the values of Lg and Lghm, which is given by the fact that parameter Lg represents the average distance between the particles of spheroidal graphite, whereas parameter Lghm represents the measured length of the line between the edges of the graphite within the matrix. This distance was selected in order for the following relation to apply: Lghm  2Rb – 2rg. The actual measurements of concentrations of ten elements – Mg, Al, Si, P, S, Ti, Cr, Mn, Fe, Ni – was carried out on the JEOL – JSM 840/LINK AN 10/85S analytical complex with an energy dispersive X-ray analyzer, an acceleration voltage of the electron beam of 25 kV and exposition time of 50 s. On each of the samples, the concentrations of all ten elements had been measured in three intervals with each individual step being 3 µm. By means of the Neophot light microscope, the interval was documented within which the concentrations were measured. The method of selection of measurement points is illustrated in Figures 1 and 2 (from the same sample). The results of measure- ments of the chemical heterogeneity of elements in cells were evaluated also statistically with the aim to be able to predict the values of two parameters: the element heterogeneity index Ihet (which is defined as the quotient of its standard deviation and its arithmetic mean) and the element segregation index IS (which is defined as the quotient of maximal concentration of elements in the cell and its arithmetic mean). The results of measurements of the chemical hetero- geneity were evaluated statistically and entered into Table 3 according to the analysed samples (x is the arithmetic mean of the concentration of the element within the measured interval, IH is the element hetero- genity index defined as the quotient of its standard deviation and its arithmetic mean, xmax is the maximum concentration of the element within the measured inter- val and IH is the segregation index of the element defined by the quotient IS = xmax / x.). The macrostructure between the bottom and upper part of the massive casting shown in Figures 3 and 4 was very different. The structure of the bottom of a massive casting is practically without foundry defects (see Figure 3 and Table 4 – local solidification time 48 s). On the other hand, in the structure of the upper part of the same casting numerous foundry defects, for example shrink hole, cavities and so on were identified (see Figure 4 and Table 4 – local solidification time 4572 s). K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 95 Figure 1: A micro-heterogeneity measurement of ductile cast-iron (the distance between two analysed graphites is 165 µm). Etched by 2 % nital. Slika 1: Meritve mikroheterogenosti duktilne `elezove litine (razdalja med dvema analiziranima grafitnima zrnoma je 165 µm). Jedkano z 2-odstotnim nitalom. Figure 2: A micro-heterogeneity measurement of ductile cast-iron (the distance between two analysed graphites is 167 µm). Etched by 2 % nital. Slika 2: Meritve mikroheterogenosti duktilne `elezove litine (razdalja med dvema analiziranima grafitnima zrnoma je 167 µm). Jedkano z 2-odstotnim nitalom. 2.4 Local solidification time The local solidification times of the selected samples of known coordinates within the casting were calculated using an original in-house 3D model 2 and are given in Table 2. The calculation of the liquidus and solidus tem- peratures for a melt with a composition according to the data in Table 1, was performed using special software with the temperature values: 1130 oC (liquidus) and 1110 °C (solidus). The values of the local solidification time ls given in Table 2 therefore relate to the temperature diffe- rence between the liquidus and solidus (Tls = 20 °C). If the local solidification time is known, then it is possible to determine the average rate of cooling of the mushy zone as a quotient of the temperature interval and the local solidification time wls = Tls /ls /(°C/s). 3 EVALUATION OF RESULTS It is obvious from the results in Tables 2 and 3 that in vertical direction from the bottom of the massive casting (sample A: y = 50 mm) to the top (gradually samples C: y = 210 mm and G: y = 450 mm) the charac- teristic and significant relations are the following: The average size of the spheroids of graphite, the average size of the cells of graphite and also the average distance between the individual particles of the graphite are all increasing. This relation was confirmed by quantitative metallographic analysis 6. The chemical heterogeneity within the individual graphite cells is also increasing. The increase in the chemical heterogeneity is reflected most significantly in the increase in the indexes of segregation IS for magnesium and for titanium, which are increasing in the direction from the bottom of the massive casting to the top in the following order: magnesium ISMg = 3.08-to- K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... 96 Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 Figure 4: In the macrostructure of the upper part of the same casting is possible to find numerous foundry defects. Etched by 2 % nital. Slika 4: Makrostruktura gornjega dela istega ulitka, kjer je lahko videti {tevilne livne napake. Jedkano z 2-odstotnim nitalom. Figure 3: Macrostructure of the massive ductile iron casting in the bottom part is practically without foundry defects. Etched by 2 % nital. Slika 3: Makrostruktura masovnega duktilnega `elezovega ulitka na spodnjem delu je prakti~no brez livnih napak. Jedkano z 2-odstotnim nitalom. Table 3: Results of the chemical heterogeneity measurements – x, xmax (w/%), IH, IS Tabela 3: Rezultati merjenj kemijske heterogenosti – x, xmax (w/%) IH, IS Sample Element Mg Al Si P S Ti Cr Mn Fe Ni rghm Rbhm Lghm x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xmax IS x IH xa IS x A 28 83 110 0.0652 0.984 0.201 3.083 0.1002 0.597 0.235 2.345 1.500 0.071 1.676 1.117 0.0151 1.805 0.120 7.947 0.0187 1.377 0.103 5.508 0.0118 1.718 0.068 5.763 0.0615 0.571 0.143 2.325 0.101 0.657 0.242 2.393 97.974 0.002 98.634 1.007 0.156 0.538 0.312 2.000 C 36 104 136 0.0387 1.550 0.260 6.718 0.0620 1.048 0.200 3.226 1.562 0.068 1.841 1.179 0.0151 2.306 0.164 10.861 0.0210 1.539 0.107 5.065 0.0065 2.382 0.061 9.385 0.0615 0.751 0.191 3.106 0.0695 0.815 0.193 2.777 97.981 0.002 98.297 1.003 0.184 0.453 0.349 1.897 G 39 109 140 0.0872 1.138 0.363 4.163 0.0761 0.815 0.216 2.838 1.396 0.076 1.650 1.182 0.0119 1.910 0.085 7.143 0.0282 1.264 0.124 4.397 0.0068 2.314 0.079 11.618 0.0959 0.4867 0.2222 2.315 0.107 0.637 0.283 2.645 98.025 0.002 98.491 1.005 0.166 0.513 0.417 2.512 -6.72-to-4.16; titanium ISTi = 5.76-to-9.39-to-11.62 (Table 3). The local solidification time, which increases from the bottom of the casting to the top – from the value of 48 s more than 50-times (near the centre of the casting) and 95-times (at the top of the massive casting), increases very significantly. The relationships between the structural characteri- stics of graphite in the casting 2L and the local solidification time were expressed quantitatively using a semi-logarithmic dependence. Despite the fact that, for the structural characteristics of graphite rghm, Rbhm and Lghm, there are only three pairs of measured values, i.e. (rghm,ls), (Rbhm,ls) and (Lghm,ls), the given dependences can be considered significant. As obvious from the research report 6, the quantitative metallographic analysis covers 49 measured views (with a magnification of 100-times) on each of the three 3D samples. This research can therefore be considered as statistically significant. The relationship between the radius of the graphite spheroids rghm and the local solidification time ls rghm /µm = 19.08 + 2.274 ln (ls/s) (1) had been found using the least-squares method. The correlation coefficient r = 0.99. Similarly, the relation Rbhm /µm = 61.33 + 5.567 ln (ls/s) (2) was established between the radius of the graphite cells and the local solidification time – with a correlation coefficient of r = 1.00, and also between the average distance of graphite particles and the local solidification time there is the relation Lghm /µm = 84.50 + 6.586 ln ( ls/s) (3) As far as chemical heterogenity of the measured elements is concerned, an analogous relation was established only for the dependence of the segregation index of titanium on the local solidification time, which has a steadily increasing course from the bottom of the casting (sample A) all the way up to the top (sample G). The relevant relation was expressed in the form of a logarithmic equation: ln IS Ti /µm = 1.201 + 0.1410 ln (ls/s) (4) where r = 0.96. 4 DISCUSSION The local solidification time  naturally affects the mechanical properties of cast-iron, however with regard to the dimensions of the test pieces, it is not possible to assign the entire body a single local solidification time. The samples for the testing of tensile strength were taken from the test-sample of the experimental casting in such a way that one had been taken from under the metallo- graphic sample and the second was taken from above. In this way, the test-samples from along the entire height of the massive casting had been taken and marked: (1A2), (3C4) and (6G7). For example, according to the marking (3C4) test-sample 3 was to be found in experimental casting beneath the metallographic sample C and test- sample 4 above it. The mechanical values determined on these test-samples are arranged in Table 4. The last column contains the local solidification times relating to samples A, C and G. Table 4: Mechanical properties of the samples from experimental casting Tabela 4: Mehanske lastnosti eksperimentalno ulitih vzorcev Test-sample Yield point Rp0.2/MPa Tensile strength Rm/MPa Ductility A5 /% Local solidifica- tion time* ls/s (1A2) 1 262 388 21.4 48 2 260 392 24.6 (3C4) 3 261 394 20.6 2509 4 266 390 19.4 (6G7) 6 260 391 14.0 4572 Note: *) metallographic samples A, C, G Table 4 indicates that in this case the local solidi- fication time ls does not affect the yield point Rp0.2 and the tensile strength Rm of the ductile cast-iron, however, it has significant influence on the ductility A5. In terms of analytical approximation, the following relationship between the ductility and the local solidification time applies: A5/% = 23.399 – 8.1703 (ls/h) (5) where r = 0.91. Simultaneously, it could be stated that for four degrees of freedom, which, according to Table 4, characterise the six statistically processed pairs (ductility, local solidification time), the correlation coefficient for this level of reliability is 0.917 8. In Eq. (5) the ductility is expressed in percentage and the local solidification time in hours. The equation indicates that the reduction in ductility of cast-iron in the state immediately after pouring is – in the first approximation – directly proportional to the square of the local solidi- fication time. 5 CONCLUSION It can be seen from previous experimentation and the evaluations of the results that led to equations (1) to (5) that – in the general case of the solidification of ductile cast-iron – there could be a dependence of the size of the spheroids of graphite, the size of the graphite cells and therefore even the distance among the graphite particles on the local solidification time, i.e. on the solidification time in which the considered point remains within the mushy zone. The described connection with the 3D model of a transient temperature field, which makes it possible to determine the local solidification time, seems to be the means via which it is possible to estimate the K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 97 differences in structural characteristics of graphite in cast-iron and also the effect of the local solidification time on ductility in the poured casting. It is known that, for example, the density of spheroids of graphite also significantly influences the mechanical properties of cast-iron, especially contraction and ductility, where the influence of the local solidification time on ductility has been verified (Equation 5). The application of the 3D model of the temperature field, together with the known and experimentally and quantitatively verified relation of microstructural charac- teristics of cast-iron, could become an effective tool for verifying the above relations in cases comprising complex and massive cast-iron castings. 6 REFERENCES 1 Kavicka, F. et al.: Optimization of the properties and the technology of the production of massive cast-iron castings. Final Research Report for the GA^R Project No. 106/01/1164. VUT-FSI, Institute of Power Engineering, Brno 2003, 68 pages 2 Dobrovska J., Kavicka F., Stetina J., Stransky K., Heger J.: Numerical models of the temperature field and chemical heterogenity of a concast steel slab. Proceedings of the 21st Canadian Congress of Applied mechanics CANCAM 2007, Toronto, Ontario, Canada, June 2007, 384–385 3 Smrha, L.: Solidification and crystallisation of steel ingots. SNTL, Prague 1983, 305 pages 4 Chvorinov, N.: Crystallisation and heterogeneity of steels. N^SAV, Prague 1954, 381 pages 5 Kavicka F., Sekanina B., Stetina J., Stransky K., Gontarev V., Dobrovska J.: Numerical optimization of the method of cooling of a massive casting of ductile cast-iron. Materiali in tehnologie/Materials and technology 43 (2009) 2, 73–78 6 Belko, J., Stransky, K.: Analysing Graphite in Cast-iron. Research Report (611–57, 811–28) VTUO Brno, Brno, November 2002 7 Winkler, Z., Stransky, K.: Heterogeneity of the Compositions of Ele- ments in Ductile Cast-iron Castings. Research Report (811-11-02) VTÚO Brno, Brno, December 2002 8 Murdoch, J., Barnes, J. A. Statistical Tables for science, engineering, management and business studies. Macmillan, Cranfield 1970, 40 pages K. STRANSKY ET AL.: TWO NUMERICAL MODELS OF THE SOLIDIFICATION STRUCTURE ... 98 Materiali in tehnologije / Materials and technology 44 (2010) 2, 93–98 A. BABUTSKY ET AL.: CORRELATION BETWEEN THE CORROSION RESISTANCE ... CORRELATION BETWEEN THE CORROSION RESISTANCE AND THE HARDNESS SCATTERING OF STRUCTURAL METALS TREATED WITH A PULSED ELECTRIC CURRENT KORELACIJA MED ODPORNOSTJO PROTI KOROZIJI IN RAZTROSOM TRDOTE PRI KONSTRUKCIJSKIH MATERIALIH, OBDELANIH S PULZIRANJEM ELEKTRI^NEGA TOKA Anatoly Babutsky1, A. Chrysanthou2, J. Ioannou2, Ilija Mamuzic3 1Institute for Problems of Strength, National Academy of Sciences of Ukraine, Timiyazevska 2 str., 01014 Kiev, Ukraine 2School of Aerospace, Automotive and Design Engineering, University of Hertfordshire, College Lane, Hatfield, AL10 9AB Herts, UK 3University of Zagreb Sisak, Croatia bai@ipp.kiev.ua Prejem rokopisa – received: 2009-09-25; sprejem za objavo – accepted for publication: 2009-10-05 The results of corrosion tests of two kinds of structural metals initially treated by a pulsed electric current of high density are presented. According to the data obtained, the treatment influences the corrosion of metals. There is a correlation between the corrosion test results and the hardness scattering of the metals. Key words: corrosion, hardness scattering, pulse electric current Predstavljeni so rezultati korozijskih preizkusov za dve vrsti kovin, ki sta bili obdelani s pulziranjem elektri~nega toka z veliko gostoto. Dobljeni rezultati ka`ejo, da obdelava s tokom vpliva na korozijo kovin. Opredeljena je bila korelacija med rezultati preizkusov korozije in raztrosom meritev trdote. Klju~ne besede: korozija, raztros trdote, pulzirajo~i elektri~ni tok 1 INTRODUCTION The interest relating to corrosion in structural metals remains high because it is related to considerable financial costs. Corrosion leads not only to a loss of metal, but also to the degradation of its mechanical and physical properties, which decreases the lifetime of components and sometimes leads to catastrophic failure. For example, a majority of failures in oil-field pipes occurs because of corrosion damage, and similar damage may occur in ships and other structures 1,2. Existing methods of metal and alloy protection from corrosion include the use of coatings with deposition, electrochemical methods of protection (e.g., cathodic protection) and others. However, there is considerable interest in developing new and more effective corrosion- protection methods. One of the common types of metal and alloy corro- sion is electrochemical corrosion, which leads to damage in conductive environments (electrolytic solutions). In general, the electrochemical corrosion mechanism involves the appearance of short-circuited micro-galva- nic elements on the metal surface with different values of e.m.f. as a result of the formation of anodic (with low electrode potential) and cathodic (with high electrode potential) zones 3. These zones are generated by differen- ces in the metal structure, surface roughness, the existence of protective films and other factors. The diffe- rence in the metal microstructure (due to a difference of the grain size, composition, crystal anisotropy, the emergence of dislocations on the surface, the presence of impurities, inclusions, non-uniform mechanical stresses) can activate the corrosion processes. The role of mechanical stresses is important because in presence of tensile stresses in the metal, anode zones can appear. These zones evolve to become corrosion centers 4. It is known that a pulsed electric current (PEC) treatment causes the relaxation of the mechanical stresses in metals 5,6, and also the homogenization of their microstructure 7. Based upon common conside- rations, these data can serve as a basis for the creation of a new technology for the corrosion protection of metals. The structural homogeneity of a metal can be estimated by a measurement of its indentation hardness 8,9. The specifics of the indentation hardness are based on the existence of a quantity correlation with other mecha- nical properties (e.g., tensile strength, fatigue limit). Consequently, if a representative array of hardness measurements is obtained, it is possible to assess the variation of the mechanical properties of the metal and its microstructural homogeneity. For example, the authors 9 have shown that the lowest level of hardness scattering corresponded to the metal with the most uniform microstructure in the initial state. On the other Materiali in tehnologije / Materials and technology 44 (2010) 2, 99–102 99 UDK 620.197; 537.39; 539.8 ISSN 1580-2949 Professional article/Strokovni ~lanek MTAEC9, 44(2)99(2010) hand, the highest hardness scattering was registered for the metal with accumulated damage after long-term performance. The results of the investigation presented here aim to clarify the possibilities of a PEC treatment to improve the corrosion resistance of steel and an aluminum alloy and also to define a correlation between the corrosion resistance and the hardness scattering in the metals after the treatment. 2 METHOD OF EXPERIMENTAL RESEARCH Metallic sheet specimens of high-strength low-alloy steel HSLA of 100 mm length, 13 mm width and 1.1 mm thickness and an aluminum alloy 5182 of 100 mm length, 13 mm width and 1.4 mm thickness were used for the investigation (the compositions of the metals are in Table 1). Table 1: Composition of the metals used (mass fraction, w/%) Tabela 1: Sestava uporabljenih kovin (masni dele`, w/%) Specimens w(Mn)/% w(Mg) /% w(Ti) /% w(C) /% w(S) /% w(P) /% Steel HSLA 0.25 – 0.30 0.02 0.02 0.02 Al-alloy 5182 0.3 4.5 – – – – The corrosion was performed at 35 °C (± 1.5 °C) in a salt-spray chamber using a 5 % NaCl solution (in distilled water) according to the ASTM B117-97. The total time of exposure was 1000 h. The treatment was carried out in hourly cycles that involved sequences of spraying the samples with a fog spray (salt spray) for 10 minutes at a flow rate of 0.7 L/h and 50 min of hot air (drying). The sequence was repeated every hour. The samples were initially hand cleaned in hot water and further cleaned for 12 min in an ultra-sonic chamber with water at 48 °C. This process took place twice because a change of water was necessary. Then the specimens were washed with propanol and dried for 30 min at 70 °C, weighed and placed in the corrosion salt-spray chamber. At the end of the corrosion test the samples were cleaned in hot water to remove the deposits of salt and then washed with propanol and dried at 70 °C and weighed. The PEC treatment of the samples was undertaken using a pulsed current generator 10. Three electric current pulses were passed through each sample with the maximum current density as indicated in Table 2. To obtain reliable values of the metal hardness and exclude the operator’s mistakes, a computerized hardness tester COMPUTEST SC (ERNST, Switzerland) was used. The hardness HRB was measured before the PEC treatment and after the treatment under an indentation load of 49 N. 3 TEST RESULTS AND THEIR ANALYSIS The corrosion test results are presented in Table 3 and in figures 1 a, 1 b and 2. The different corrosion behavior of the steel and aluminum alloys must be consi- dered. In the test result analysis on steel, the corrosion products (Fe2O3 oxides) flake off the base material causing the reduction of the specimen weight after the corrosion tests, while, in the case of the aluminum alloy, an increase in the specimen weight is observed after the corrosion tests because Al2O3 oxides have high adhesion to the base material. The estimation of the influence of the PEC treatment on specimens corrosion was carried out in two steps. In the first step, the relative change, m, of the weight of each specimen after the corrosion test was calculated, m m m m / % = − ⋅c 0 0 100 where m0 and mc are the weights of the samples before and after the corrosion test, respectively. In the second step, the influence factor of the PEC treatment on the specimen corrosion, kc, was determined as: A. BABUTSKY ET AL.: CORRELATION BETWEEN THE CORROSION RESISTANCE ... 100 Materiali in tehnologije / Materials and technology 44 (2010) 2, 99–102 Table 2: Regimes of PEC treatment Tabela 2: Re`imi obdelave z elektri~nim tokom No. Steel HSLA Aluminum alloy5182 Current density, j(kA/mm2) 1 1.3 1.3 2 1.3 1.3 3 2 2 4 2 2 5 3 3 6 3 3 Remark: specimens No. 1, 3, 5 were treated using three pulses with a 1 min. interval, specimens No. 2, 4, 6 were treated using three pulses with a 3 min. interval Figure 1: Change of the effect of PEC treatment on metal corrosion kc (a, b) and hardness scattering kh (c, d) against the metal (a, c – HSLA steel; b, d – aluminium alloy 5182), density of PEC, and interval between electric current pulses (1, 3 – 1 min interval; 2, 4 – 3 min interval). Slika 1: Sprememba u~inka obdelave z elektri~nim tokom na korozijo kovine kc (a, b) in raztros trdote kh (c, d) v odvisnosti od kovine (a, c – konstrukcijsko jeklo; b,d – zlitina aluminija), gostota toka in interval med elektri~nimi pulzi (1, 3 – 1 min interval; 2, 4 – 3 min interval) k m m m c untreat treat untreat / % = − ⋅    100 where muntreat is the relative change of the weight of the untreated specimens (the mean value of the 6 specimens in the bottom row in Table 3) and mtreat is the relative change of the weight in percent of the six specimens treated by PEC (Table 3). Table 3: Results of the corrosion tests Tabela 3: Rezultati korozijskih preizkusov No. (corresponds to table 2) Steel HSLA Aluminum alloy 5182 After PEC treatment m/% kc/% m/% kc/% 1 –0.5224 +26 +0.1217 +15 2 –0.5310 +25 +0.1571 –9 3 –0.2706 +62 +0.1188 +17 4 –0.60 10 +15 +0.0614 +57 5 –0.6890 +2 +0.1640 –14 6 –1.2729 –81 +0.1426 +1 Without PEC treatment Mean for 6 specimens m/% m/% –0.7050 +0.1437 The corrosion test results suggest that the PEC treatment with the above regimes substantially influences the behavior of the investigated metals. The treatment causes considerable deceleration of the corrosion processes for both the HSLA and aluminum alloy 5182. Here, the results presented show that the regimes of the PEC treatment with maximum effect exist. Table 4: Results of indentation tests Tabela 4: Rezultati meritev trdote No. (corresponds to table 2) Steel HSLA Aluminum alloy 5182 µ kh/% µ kh/% 1 – untreated 0.0152 +11 0.0615 –5 1 – treated 0.0168 0.0584 2 – untreated 0.0120 +20 0.0501 +14 2 – treated 0.0144 0.0572 3 – untreated 0.0153 +39 0.0751 –22 3 – treated 0.0212 0.0587 4 – untreated 0.0171 +267 0.0847 –33 4 – treated 0.0627 0.0566 5 – untreated 0.0186 +431 0.0531 +19 5 – treated 0.0986 0.0630 6 – untreated 0.0150 +271 0.0512 +2 6 – treated 0.0557 0.0525 The indentation test results are presented in Table 4 and in Figures 1 c, 1 d. On average, 120 measurements of hardness per condition (treated and untreated) were fulfilled for each specimen. The results were processed and a relative value µ, with  as the standard deviation and µ as the hardness average were calculated for each specimen and condition. Then the influence factor of the PEC treatment on the metal hardness scattering, kh, was determined, as presented below: kh = − ⋅ µ µ µ    treat untreat untreat 100% The indentation test results demonstrate the sensiti- vity of the hardness scattering to the regimes of PEC treatment and the corrosion resistance. The PEC-treated HSLA sample sets 4 and 5 showed a lower increase of the corrosion resistance than the sample sets 1, 2 and 3, and the set 6 showed an increase of corrosion (Table 3 and Figure 1 a). At the same time, the samples sets 4, 5 and 6 exhibited a substantial increase in the hardness scattering (Table 4 and Figure 1 c). In the case of the aluminum alloy 5182, the effect of the PEC treatment on the corrosion resistance is mixed as shown in Table 3 and Figure 1 b with the PEC-treated sample sets 1, 3 and 4 with a lower corrosion intensity, sets 2 and 5 with a stronger corrosion, while the values for set 6 are practically similar. At the same time, the PEC-treated sample sets that corroded less (1, 3 and 4), also exhibited a lower level of scattering of the hardness value, indi- cating a more homogeneous composition and micro- structure (Table 4 and Figure 1 d). On the other hand, sets 2 and 5 corroded more, but showed a higher degree of scattering of the hardness data, implying a less uniform structure and composition. Macrographs of the surface appearance of the investi- gated HSLA specimens after the corrosion tests are Materiali in tehnologije / Materials and technology 44 (2010) 2, 99–102 101 A. BABUTSKY ET AL.: CORRELATION BETWEEN THE CORROSION RESISTANCE ... Figure 2: Surface appearances of HSLA steel specimens after corrosion tests: a, b (top photo) – metal in initial condition; b (bottom photo), c – after PEC treatment Slika 2: Videz povr{ine konstrukcijskega jekla po korozijskem preizkusu: a, b (zgornja slika) – kovina pred preizlkusom; b ( slika spodaj), c – po obdelavi s tokom presented in Figure 2 (the difference in the appearance of the treated and untreated specimens in the case of the aluminum alloy specimens is visibly insignificant). A substantial difference in the corrosion damage of the alloy surface took place for the HSLA specimens; in the case of the specimens without the PEC treatment, zones with selective, localized corrosion are visible in Figure 2 b (top) and more obviously it is visible in the areas bordered by the dashed lines in Figure 2 a. The PEC treatment was observed to cause weaker, more uniform and homogeneous corrosion on the whole surface of specimen, see Figure 2 b (bottom). Only small pittings are visible there, see the areas bordered by the dashed lines in Figure 2 c. The obtained results may be related to the different microstructure and the surface homogeneity of the treated and untreated specimens. The differences may arise during the production of the sheet with the cold-working that builds up non-uniform residual stresses. On the other hand, the PEC treatment may relieve these stresses, providing a more homogeneous material. 4 CONCLUSIONS The results of the investigation have shown that the PEC treatment influences the corrosion behavior of structural metals. Based on these observations, the following conclusions can be drawn: The regimes of the PEC treatment causing the deceleration of the corrosion with maximum effect exist for both the HSLA and aluminum alloy 5182. A correlation between the corrosion resistance and the hardness scattering of the PEC-treated metals exists. The specimens with the lower hardness scattering show the better corrosion resistance. As a result of the PEC treatment, the corrosion on the surface of HSLA steel samples became smaller and more homogenous and the zones of localized corrosion observed on the untreated samples are absent on the surface of PEC-treated HSLA. The investigation was based on one sample per variant of PEC treatment. Further research will be carried out using more samples, aiming to accumulate experimental data on the influence of the PEC treatment on the characteristics of the corrosion of structural metals, as well as to investigate the reasons and physical mechanisms of such effects. 5 REFERENCES 1 Emi, H.; Kumano, A.; Baba, N.; et al.: A study on life assessment of ships and offshore structures (Part 1: basic study), J. Soc. Nav. Archit. Jpn., 169 (1991), 443–454 2 Emi, H.; Yuasa, M.; Kumano, A.; et al.: A study on life assessment of ships and offshore structures (3rd report: corrosion control and condition evaluation for a long life service of the ship), J. Soc. Nav. Archit. Jpn., 174 (1993), 735–744 3 Isaev, N. I.: Theory of corrosion processes, Publishing house Metallurgy, Moscow 1997 (in Russian) 4 http://www.cathedral.ru/cathedra/num2/demidov (in Russian) 5 Baranov, Y. V., Troitskiy, O. A., et al.: Physical bases of electro- pulse and electro-plastic treatments and new materials, Publishing house MGIU, Moscow 2001 (in Russian) 6 Stepanov, G.; Babutsky, A.; Krushka, L.: Metal behavior under passage of impulse electric current, J.Phys. IV France, 110 (2003), 577–582 7 Stepanov, G. V.; Babutsky, A. I.; Chyzhyk G. V.: Estimation of impulse electric current influence on strength of metallic materials, Metals science and treatment (Metaloznavstvo ta obrobka metaliv), (2005) 2, 64–68 (in Ukrainian) 8 Lebedev, A. A.; Muzyka, N. R.; Volchek N. L.: Determination of damage accumulated in structural materials by the parameters of scatter of their hardness characteristics, Strength of Materials 34 (2002) 4, 312–321 9 Lebedev, A. A.; Muzyka, N. R.; Volchek N. L., et al.: Monitoring of current state of pipe metal in active gas pipelines. Experimental method and results, Strength of Materials, 35 (2003) 2, 122–127 10 Stepanov, G. V.; Babutskii, A. I.; Mameev, I. A.: High-density pulse current-induced unsteady stress-strain state in a long rod, Strength of Materials, 36 (2004) 4, 377–381 A. BABUTSKY ET AL.: CORRELATION BETWEEN THE CORROSION RESISTANCE ... 102 Materiali in tehnologije / Materials and technology 44 (2010) 2, 99–102