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Metallurgia Italiana

International Journal of the Italian Association for Metallurgy

n. 1 Gennaio 2018 Organo ufficiale dellâ&#x20AC;&#x2122;Associazione Italiana di Metallurgia. Rivista fondata nel 1909

Corso itinerante

Solidificazione e colata continua 8-9-15-16-22-23 marzo 2018 Organizzato da

CENTRO DI STUDIO ACCIAIERIA

Con il supporto di

L’Associazione Italiana di Metallurgia organizza una nuova edizione del Corso sulla solidificazione e colata continua degli acciai per continuare a sostenere le imprese nell’azione di formazione del proprio personale. Il Corso abbraccerà i temi della solidificazione, le problematiche relative alla struttura della macchina, i componenti refrattari, le polveri di copertura, la difettologia, il colaggio di billette, blumi e bramme etc.. Data l’importanza dei semilavorati destinati alle operazioni di forgia, durante il Corso si affronteranno anche le problematiche relative al colaggio dei lingotti e dei blumi di grande dimensione. Secondo la formula tradizione, la formazione verrà erogata con due modalità: lezioni di tipo teorico - volte a fornire i concetti di base relativi agli aspetti metallurgici e al funzionamento degli impianti - e visite tecniche presso gli impianti produttivi. Per rispondere a queste esigenze il Corso è itinerante e le lezioni si svolgeranno presso alcune interessanti realtà produttive: AST, Acciaieria di Calvisano, Alfa Acciai, NMLK e A.C.P. In occasione delle visite i tecnici delle società ospitanti presenteranno gli impianti con una particolare attenzione agli aspetti caratteristici di ogni sistema di colata. Queste attività saranno affiancate ed integrate da interventi didattici tenuti da docenti universitari, nonché da esperti di società di ingegneria di riconosciuta esperienza. Gli interventi didattici e le visite tecniche si articolano su un arco di sei giorni e sono organizzati con cadenza tale da evitare ai partecipanti un’assenza eccessivamente prolungata dalle proprie aziende. Coordinatori del Corso: Silvia Barella, Francesco Magni, Carlo Mapelli Per informazioni ed iscrizioni: AIM · Associazione Italiana di Metallurgia Tel. 02-76021132 / 02-76397770 · E-mail: info@aimnet.it · www.aimnet.it

#corso #formazione #impianti #solidificazione #colatacontinua

La Metallurgia Italiana International Journal of the Italian Association for Metallurgy Organo ufficiale dell’Associazione Italiana di Metallurgia. House organ of AIM Italian Association for Metallurgy. Rivista fondata nel 1909

Direttore responsabile/Chief editor: Mario Cusolito Direttore vicario/Deputy director: Gianangelo Camona Comitato scientifico/Editorial panel: Livio Battezzati, Christian Bernhard, Massimiliano Bestetti, Wolfgang Bleck, Franco Bonollo, Bruno Buchmayr, Enrique Mariano Castrodeza, Emanuela Cerri, Lorella Ceschini, Mario Conserva, Vladislav Deev, Augusto Di Gianfrancesco, Bernd Kleimt, Carlo Mapelli, Jean Denis Mithieux, Marco Ormellese, Massimo Pellizzari, Giorgio Poli, Pedro Dolabella Portella, Barbara Previtali, Evgeny S. Prusov, Emilio Ramous, Roberto Roberti, Dieter Senk, Du Sichen, Karl-Hermann Tacke, Stefano Trasatti Segreteria di redazione/Editorial secretary: Valeria Scarano Comitato di redazione/Editorial committee: Federica Bassani, Gianangelo Camona, Mario Cusolito, Ottavio Lecis, Carlo Mapelli, Valeria Scarano Direzione e redazione/Editorial and executive office: AIM - Via F. Turati 8 - 20121 Milano tel. 02 76 02 11 32 - fax 02 76 02 05 51 met@aimnet.it - www.aimnet.it

n. 1 Gennaio 2018 Anno 110 - ISSN 0026-0843

Acciai inossidabili e acciai duplex / Stainless steel & duplex Laser welding of plastically deformed lean duplex stainless steel I. Calliari, C. Gennari, E. Hurtado Delgado, A. F. Miranda Pérez, B. R. Rodriguez Vargas 5 Martensite quantification, mechanical properties and cold rolling in AISI 301 Austenitic Stainless Steel P. Piccardo, R. Spotorno, D. Lanteri, F. Canepa, I. Citi 11 Mechanical properties evolution on heat treated severe cold rolled UNS S32760 Super Duplex Stainless Steel C.M. Tromellini, A.F. Ciuffini, A. Gruttadauria, S. Barella, C. Di Cecca, C. Mapelli 20 Different effects of carbon and nitrogen on precipitation behavior and mechanical properties in austenitic stainless steels Kyung-Shik Kim, Jee-Hyun Kang, Sung-Joon Kim 28 Attualità industriale / Industry news New light corrosion resistant steel without chromium a cura di: C. Mapelli, A. F. Ciuffini, S. Barella, A. Gruttadauria, D. Mombelli

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Gestione editoriale e pubblicità Publisher and marketing office: Siderweb spa Via Don Milani, 5 - 25020 Flero (BS) tel. 030 25 400 06 - fax 030 25 400 41 commerciale@siderweb.com - www.siderweb.com La riproduzione degli articoli e delle illustrazioni è permessa solo citando la fonte e previa autorizzazione della Direzione della rivista. Reproduction in whole or in part of articles and images is permitted only upon receipt of required permission and provided that the source is cited. Reg. Trib. Milano n. 499 del 18/9/1948. Sped. in abb. Post. - D.L.353/2003 (conv. L. 27/02/2004 n. 46) art. 1, comma 1, DCB UD Siderweb spa è iscritta al Roc con il num. 26116 Stampa/Printed by: Poligrafiche San Marco sas - Cormòns (GO)

Caratterizzazione microstrutturale ed elettrica di Cu1.2Mn1.8O4 applicato come rivestimento protettivo in pile a combustibile ad ossidi solidi (SOFC) a cura di: R. Spotorno, P. Piccardo, S. Barison, S. Fasolin 44 Comparative investigation of deep drawing formability in austenitic (AISI 321) and in ferritic (DIN 1.4509) stainless steel sheets a cura di: C. de Paula Camargo Pisano, H. J. B. Alves, T. Reis de Oliveira, C. G. Schön 50 Strain induced martensite evolution in a rolling contact of SS AISI 304 a cura di: M. Werschler, P. Gümpel, K. Werner 62 Scenari / Experts’ Corner Le norme EN 10088 per gli acciai inossidabili a cura di: Mario Cusolito

Atti e notizie / Aim news Calendario degli eventi internazionali / International events calendar 74 AIM – UNSIDER Norme pubblicate e progetti in inchiesta (aggiornamento 29 dicembre 2017) 75

l’editoriale La Metallurgia Italiana Cari conSoci,

Prof. Gianangelo Camona

AIM e lettori de La Metallurgia Italiana, giunto al termine dell’incarico di direttore responsabile della Rivista – durato vent’anni, dal 1997 – voglio accommiatarmi innanzitutto augurandovi un buon anno, sia personale che di lavoro, e poi facendo insieme un rapido esame di questo ventennio durante il quale anche paesi con una illustre tradizione metallurgica hanno dovuto rinunciare alle loro riviste nazionali. Ripensandoci ora, per La Metallurgia Italiana e la sua sopravvivenza, si è trattato di un periodo ricco di cambiamenti graduali, ma significativi, che hanno culminato nella decisione che ha portato – a metà del 2016 - al passaggio dal supporto cartaceo a quello informatico e alla sostanziale monotematicità di ogni numero della Rivista. Naturalmente la molla di tutto ciò non è stata un amore del nuovo in sé ma la duplice esigenza di tener d’occhio gli aspetti economici e di cercare di essere sempre più efficaci nel supportare la cultura metallurgica nel nostro Paese, secondo il mandato statutario dell’AIM. Per il primo aspetto, una volta individuata una veste formale della Rivista (numero di pagine, policromia, tipo di carta, cadenza delle pubblicazioni, ecc. ), si sono applicate soluzioni che pesassero il meno possibile sul bilancio dell’Associazione, soprattutto incentivando la raccolta di inserzioni pubblicitarie. Questo compito era stato svolto, decenni fa, direttamente dall’AIM quale editore della Rivista, ma poi venne affidato ad un editore esterno. Comunque sia, nell’ultimo ventennio gli anni floridi per il comparto metallurgico sono stati pochi e questo ha spinto a cercare di volta in volta le soluzioni che risultavano interessanti per i costi di stampa e distribuzione e per la raccolta pubblicitaria. E ciò ha comportato, via via nel ventennio, al passaggio dalla Franco Angeli all’ Edimet, alla Honegger, a CONSEDIT e ora alla Siderweb. Per l’aspetto relativo al supporto ai lettori - in termini di cultura metallurgica – si è provveduto progressivamente ad introdurre nuove rubriche - come quella mirata a rendere noti i lavori dei Centri di Studio tramite estratti dai verbali dei Comitati Tecnici o quella che aggiorna sulla normativa - ma soprattutto immettendo nella sezione delle memorie i lavori risultati più significativi nelle varie manifestazioni dell’AIM, in modo da renderli accessibili anche a chi non avesse potuto partecipare all’evento o volesse semplicemente essere aggiornato su problematiche non necessariamente strettamente attinenti alle proprie competenze. Ovviamente la novità più significativa risulta la già citata impostazione generale del 2016, realizzata su impulso del Presidente Mapelli, che va ben al di là della semplice informatizzazione perché finalizzata a favorire l’avvicinamento delle aziende all’Associazione, con evidenti vantaggi per la cultura metallurgica nazionale. Infatti, fascicoli monotematici e, poi, rubriche come: “Attività Industriale”, “Scenari”, “Economia e produzione” permettono di entrare in modo approfondito ed esteso nei problemi delle varie problematiche metallurgiche, soprattutto calandoli maggiormente nelle realtà produttive. Termino richiamando alla memoria di chi c’era, e informando le nuove leve, che nel bel mezzo del ventennio di cui sto parlando – nel 2009, come appare nella riproduzione della copertina (nell'immagine a sinistra) - la Rivista ha raggiunto il secolo di vita, festeggiato con una importante manifestazione presso il Museo della Scienza e della Tecnologia “Leonardo da Vinci”. Ancora buon anno e buon lavoro.

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex Laser welding of plastically deformed lean duplex stainless steel I. Calliari, C. Gennari, E. Hurtado Delgado, A. F. Miranda Pérez, B. R. Rodriguez Vargas Duplex stainless steels are preferred in industries where a combination of mechanical properties with higher corrosion resistance is required. Their excellent properties are due to the biphasic microstructure, however, when thermal cycles are applied can be highly affected by the formation of hazardous intermetallic phases decreasing their performance. Welding as beneficial can be detriment for joining stainless steels; moreover, some welding processes are limited for narrow thicknesses or the electrode can promote intermetallic phases. Laser welding is an advanced joining process that is currently replacing traditional fusion processes, besides it can be employed to join thin plates. The aim of this study is to observe the plates of 2304 DSS plastically deformed joined by pulsed laser welding varying the peak power in order to observe the microstructural modifications and penetration of the joint. The experiments were performed by fixing some laser welding parameters in order to observe only the influence of the peak power. Microstructural observations by means of optical and electron microscope were performed, within EDS analysis to identify possible intermetallic phases. Microhardness evaluation was performed observing variations mainly at the surface.

KEYWORDS: LASER WELDING, DSS, SEM, COLD WORKED, CHARACTERIZATION

INTRODUCTION Industry demands materials with excellent performance, good mechanical properties and environmental resistive (1); Duplex stainless steels (DSS) fulfill these requirements, that is how are predominantly employed in pulp and paper, chemical tankers, pipelines, reactors plants and nuclear waste containers (2; 3). In corrosive conditions DSS proved to have higher resistance than conventional stainless steels, besides the approximately equal combination of both phases (ferrite and austenite) provides them of higher mechanical strength (4; 5). The yield strength achieves 450 MPa, and is excellent for applications were stress corrosion resistive materials are required. Additionally, referring to Lean Duplex Stainless (LDSS), it is almost Mo free and the nickel content is lower than 4% that is traduced on low-price stainless steel (6). DSS are in valuable demand of grades with excellent performances that were customized by more than 20 years of successful service applications (3). In manufacturing, assembly for DSS became a challenge since some fusion welding processes significantly affects the phase balance, and in certain duplex grades strength the formation of deleterious phases (7). The processes more often employed are submerged arc welding (SAW), gas tungsten arc welding (GTAW), shielded metal arc welding (SMAW) and laser beam welding (LBW). Laser is related as products with high quality, precision and greater speed, which is traduced in high productivity at several industrials applications. This process is eight times faster than Gas Tungsten Arc Welding (GTAW) (8). Many have reported investigations regarding LBW, fiber laser can La Metallurgia Italiana - n. 1 2018

be employed with proper parameters and shielding N2 gas as reported by R. Lai et al. (9). However, Neodymium-doped Yttrium Aluminum Garnet (Nd:YAG) industrial lasers confirmed reductions in the heat affected zone, improved weldability and narrow fusion zone (10). In these lasers, the energy is supplied by a laser diodes, the wavelength is shorter than fibers lasers,

I. Calliari, C. Gennari Industrial Engineering Department (DII), University of Padua, Via Marzolo 9, Padua, Italy

E. Hurtado Delgado, A. F. Miranda Pérez, B. R. Rodriguez Vargas Corporación Mexicana de Investigación en Materiales, Ciencia y Tecnología 790. Saltillo, México

Stainless steel & duplex which supports the laser for robotic purposes. Several studies agree with the increment of ferrite phase during joining (11), which can affect severely the corrosion resistance (12). In this study, cold rolled lean duplex stainless steel was joined by Nd: YAG laser welding in order to observe the microstructural effects after welding, with the purpose to increase the mechanical strength of the material.

cold rolled until achieve 40% of thickness reduction (Fig. 1), Then, welded by Nd:YAG micro-laser of 160 W model HTS LS P-160 using a 1064-nm wavelength, no clamping to the plates was carried out. Prior to laser welding, no polished with grit paper was needed since the surface was adequate for avoiding the reflectivity. However, the samples were cleaned with acetone to remove any grease and other contaminants.

EXPERIMENT Lean duplex 2304 plates (254 Hv) of 5 mm thickness were

Fig. 1 - Schematic view of the performed laser welds The micro-laser process was pulsed type, employing argon atmosphere (flow: 0.5 bar) in order to reduce metal oxidation and contamination. The welding speed was fixed at 0.2 mm/s, pulse width of 7 ms with a repetition rate of 5.5. The peak power was

then modified as shown on Tab. 1, acting as the varied process parameter. The chemical composition of the base metal is presented on Tab. 2.

Tab. 1 - Pulsed laser parameter.

Peak Power (W)

Heat input (J/mm)

Sample 1

136

680

Sample 2

128

640

Sample 3

120

600

Sample 4

112

560

Sample 5

144

720

Sample 6

152

760

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex Tab. 2 - Chemical composition (wt. %) of the LDSS 2304. C

0.02

0.48

0.29

1.37

0.36

23.5

5.02

0.014

0.031

0.10

Balance

Selected samples were cross-sectioned and prepared with the traditional metallographic procedure (grinded and polished), and then etched with Villella’s (picric acid) for 15 seconds. For macroscopic examinations, a stereoscope was employed, and for microscopic analysis, the optical and scanning electron (TESCAN MIRA 3) microscopes were used. Semi-quantitative chemical analysis of the samples was obtained using a Quantax EDS Bruker detector. Hardness profile was fixed a 500 kgf using a Vickers indenter. RESULTS AND DISCUSSION Microstructural characterization The as-received material presented elongated grains that af-

ter deformation were slightly flattened. Macrographs of three representative samples are presented in Fig. 2. The joint with less peak power presents lower penetration compared to higher one, still when the heat input is increased the generated energy displace the plates and distortion is observed as in Fig. 2c. Grains size development varied in the selected samples, in pulsed laser welding the melting has two principal stages: at first the interaction of heat source which melts on one-step, and re-melting in order to achieve a hermetic seal produce by heat. Stress and distortion are produced due to a higher cooling rate which produce a rapidly solidification with solid-state transformation (10).

Fig. 2 – Weld cross sections of a) 128 W, b) 144 W) and c) 152 W of peak power samples. Deformation by shrinkage is observed in Fig. 3, the heat generated in the process provokes a non-uniform expansion of the base metal within contraction for the heating and cooling cycle. Since the pulse in this process re-melts the material, physical properties of the base metal decreases while thermal expan-

sion increases. In this case, the achieved mechanical strength due to cold rolled decreases since the heat input generated by the process softens the materials leading to crack initiation. Despite the low heat input values it generates distortion.

Fig. 3 – Initiation crack at joint 1 with 680 J/mm of heat input La Metallurgia Italiana - n. 1 2018

Stainless steel & duplex During solidification, rapid cooling rate impinges the ferriteaustenite balance as observed in Fig. 4. Austenite is then decreased in volume, but is clearly grown as Grain Boundary Austenite (GBA); the transformation of this phase involves the formation along the ferrite grain boundaries of coar-

se ferrite grains that leads into a higher density austenite (7,13). As can be seen the heat affected zone is almost invisible since the cooling rate occurs rapidly.

Fig. 4 â&#x20AC;&#x201C; Welding interface in joint 6 The variation in the overlap factor produced by pulsed laser welding is reported in Fig. 5, layers were formed and epitaxial ferrite growth is presented at the first pulsed, se-

quentially at the second pulse the grain is dissolved, in some cases by the half, and grain texture is modified.

Fig. 5 â&#x20AC;&#x201C; Scanning electron micrograph of the joint 5 at the fusion zone Fusion line was selected for EDS analysis, in the macro view Cr and Ni distribution seems to be not affected by the process, however, the accumulation of Mn in the solid is confirmed by the map in Fig. 6. The Cr/Ni redistribution is evident

in the fusion zone and the typical microstructure of banding disappears. In the electron micrograph, the heat-affected zone is undetectable.

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex

Fig. 6 – SEM micrograph and EDS maps at the interface (fusion zone and base metal) Microhardness analysis Fig. 7 presents the hardness profiles; initially the hardness value for the base metal (without plastic deformation) is 254 Hv, while the cold rolled material is about 370 Hv. The higher mechanical strength was marked even after welding, which means that LDSS maintain its mechanical properties

in the weld. On the other hand, some fusion zones were affected by recovery since the material softens due to high heat input. Those grains in the fusion zone were not recrystallized, but were hardened by recovery during plastic deformation.

Fig. 7 – Hardness evaluation of 680 J/mm (sample 1), 640 J/mm (sample 2) and 720 J/mm (sample 5) heat input CONCLUSIONS In this investigation the microstructural features produced by plastic deformation and welding in LDSS were studied. By cold rolling, 40% of thickness reduction was achieved within approximately 370 Hv of hardened material. The main conclusions are summarized as follows: • Increment of mechanical strength due to plastic deformation was obtained. • Higher penetration by the laser welding process was achieved with a 720 J/mm heat input. Despite of low heat input, distortion was presented in some samples. • The laser welding presents some advantages, such as, lack La Metallurgia Italiana - n. 1 2018

of intermetallics formation during weld, welding quality and absence of the heat affected zone maintaining their mechanical properties. Even though, the higher cooling rates softens the materials leading to shrinkage which ends as a crack. • Ferrite phase was higher in the fusion zone compared to austenite, which is presented as a grain boundary austenite. • During plastic deformation some grains were constrained of recrystallization, but presented recovery. After laser welding, which occurs so rapidly at high temperature, those grains were softened decreasing the hardness nearly the base metal-fusion zone interface. 9

Stainless steel & duplex REFERENCES [1]

A. VINOTH JEBARAJ, L. AJAYKUMAR, C.R. DEEPAK and K.V.V. ADITYA, J. Advan. Resear. 8, (2017), p. 183.

[2]

M. BOLUT, C.Y. KONG, J. BLACKBURN, K.A. CASHELL and P.R. HOBSON. Phys. Proced. 83, (2016), p. 417.

[3]

M.A. ELSAADY, W. KHALIFA, M.A. NABIL and I.S. EL-MAHALLAWI. Ain. Sham. Eng. J. (2016), p. 1.

[4]

S. SRIKANTH, P. SARAVANAN, P. GOVINDARAJAN, S. SISODIA and K. RAVI. Advan. Mater. Resear. 794, (2013), p. 714.

[5]

J. VERMA and R. VASANTRAO TAIWADE. J. Manuf. Proc. 25, (2017), p. 134.

[6]

C. HERRERA, D. PONGE and D. RAABE, Mater. Scie. Foru. 715, (2012), p. 550.

[7]

L. KARLSSON. Weld. in the World (2004) , 26, p. 65.

[8]

K. WEMAN. Welding Processes Handbook. Woodhead (2012), p. 205.

[9]

R. LAI, Y. CAI, Y. WU, F. LI and X. HUA. J. Mater. Process. Techn. 231, (2016), p. 397.

[10]

G. SIVAKUMAR, S. SARAVANAN and K. RAGHUKANDAN. Optik. 131, (2017), p. 1.

[11]

F. MIRAKHORLI, F. MALEK GHAINI and M.J. TORKAMANY, JMEPEG. 21, (2012), p. 2173.

[12]

J.S. KU, N.J. HO and S.C. TJONG. J. Mater. Process. Techn. 63, (1997), p. 770.

[13]

I.CALLIARI, M.ZANESCO, E.RAMOUS AND P.BASSANI: JMEP 16,1 (2007), P. 109.

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex Martensite quantification, mechanical properties and cold rolling in AISI 301 Austenitic Stainless Steel P. Piccardo, R. Spotorno, D. Lanteri, F. Canepa I. Citi Austenitic Stainless steel known as AISI 301 are often used as cold rolled foil with very thin thickness (down to 0.2 mm). Such materials are further shaped by bending or stamping and need to show a suitable deformability even at the cold hardened stage as is the case of the commercialized product. The quantification of the martensite, precipitated as main by product of the cold deformation of the metastable austenitic microstructure, is considered of dramatic importance to predict the toughness of this alloy. In this paper an original method to quantify the martensite by using the magnetic susceptibility (SQUID) is presented and the combination of the so collected data with stress-strain curves on differently deformed samples are discussed. The conclusions introduce the role of cold hardening in the mechanical properties of cold deformed AISI 301 beside the effect of the martensite volume fraction.

KEYWORDS: AUSTENITIC STAINLESS STEEL, MARTENSITE, COLD HARDENING, SQUID, STRESS-STRAIN CURVES, MECHANICAL PROPERTIES

INTRODUCTION Austenitic stainless steels are interesting engineering materials, due to their high corrosion resistance and versatile mechanical properties. They have excellent ductility and toughness in annealed condition. Furthermore their tensile strength can be greatly increased by cold working due to their metastable nature. Such a skill is mainly related to the formation of martensite microphases dispersed into the austenitic matrix as a consequence of the energy absorbed during the cold deformation or the mechanical stress in general. The martensite is stronger and harder than the austenitic structure, causing a sort of precipitation strengthening effect and thus a high strain effect. This known phenomenon [1] is on one side explained by the natural evolution of the metastable austenite into a more stable lattice at room temperature but on the other side there are still kinetic and thermodynamic aspects to be investigated and understood in order to model the martensite transformation and to make it profitable for technical applications [2]. The thermal treatment between the cold rolling stages seems also to play a role by its efficacy in recovering deformability and, at the same time, dissolving the martensite precipitates. The final state is often the cold rolled one and thus cold hardened with an amount of martensite which needs to be quantified in order to predict the basic mechanical properties of the metal sheet. Due to their expected excellent ductility austenitic stainless steels find applications in those fields where severe forming operations are required [3]. They are also gaining more interest for their combination of formability and high strength after forming. An annealed sheet can easily be formed to complex shapes but this La Metallurgia Italiana - n. 1 2018

skill is also required finished cold rolled sheets. The combination of cold rolling steps and recovery thermal treatments become thus strategic in order to offer the lowest risk of failure. At the same time a reliable prediction of the mechanical properties under stress would make such steels even more attractive for specific applications. To reach such a goal a reliable tool to quantify martensite is therefore demanded as well as a deeper understanding of its formation beside other transformation observed

P. Piccardo DCCI, University of Genoa, via Dodecaneso 31, 16146, Genoa, Italy - paolo.piccardo@unige.it - ICMATE, CNR, Via De Marini 6,16149 Genoa, Italy

R. Spotorno, D. Lanteri, F. Canepa DCCI, University of Genoa, via Dodecaneso 31, 16146, Genoa, Italy

I. Citi ARINOX S.p.A., Via Gramsci 41/A, 16039 SESTRI LEVANTE (Genoa), Italy

Stainless steel & duplex during the manufacturing process. The deformation induced martensitic transformation is a complex transformation and the theory on martensitic transformation remains vague. This issue depends on many factors both practical and analytical. On one hand there is the lack of profound understanding on how parameters such as chemical composition, temperature, strain rate, grain size, deformation mode affect the transformation. On the other hand the is the need to measure the amount of martensite even at very low concentration and to relate its volume fraction to the manufacturing parameters and to the mechanical properties. It is known that the presence of martensite is changing the magnetic properties of the austenitic stainless steel [4] and thus this can be used as an alternative to the classical XRD method which is not enough sensitive to low

martensite volume fraction and to the expensive and too local TEM investigation. In this paper the SQUID measuring system, based on the quantification of the martensite volume fraction by its magnetic properties, is used on industrial products obtained by standard manufacturing. The same materials are also mechanically characterized by stress-strain experiments and the mechanical properties are related to the microstructural features. MATERIALS AND METHODS The compositional data reported in table 1 are referred to the rolled stainless steel AISI 301 that has been studied in this research.

Tab.1 - Steel composition (AISI 301) Element

Wt.%

0.10

1.11

1.20

0.024

<0.01

0.13

6.4

16.9

0.68

0.059

To reach the final thickness of the metal sheet several sessions of cold rolling and recovery bright annealing were applied

using industrial parameters on an unidirectional mill. Table 2 reports the thermomechanical treatment details.

Tab. 2 - Thermomechanical parameters Cold Rolling Cycle

initial - final thickness (mm)

No. of steps

% reduction

Bright Annealing

0.90 - 0.36

1150Â°C in N2

0.36 - 0.25

30.5

With the purpose to quantify the martensite volume fraction and to determine the mechanical properties a sample from

each step of the second cold rolling cycle was acquired and investigated, table 3 reports the complete list of samples.

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex Tab. 3 - List of samples Sample Name

Step

Thickness (mm)

Def. rate (%)

∆(def.rate)

0.360

ANNEALED

0.36

0.330

0.33

8.3

0.296

0.295

17.8

9.5

0.285

20.8

2.7

0.280

0.28

22.2

1.4

0.276

23.3

1.1

0.261

27.5

4.2

0.254

29.4

1.9

0.252

0.6

0.250

0.25

30.5

0.5

The Δ(def.rate) was calculated by subtracting the deformation rate of previous sample to the deformation rate of the considered

sample (as shown in equation 1): is a worth information to easily evaluate the strain entities to single step.

Δ(def.rate) = [Deformation Rate(sample x)] – [Deformation Rate(sample x-1)] The investigation methods applied to all samples are: 1. Tensile test 2. Magnetometer SQUID M vs H measure 3. Optical Microscope (LOM) metallography 4. Scanning Electron Microscope (SEM) fractographic analysis Only results statistically representative of a group of samples and meaningful for the discussion will be presented in the following paragraphs. Investigations at points 1 and 2 were made on the sampled specimen without further preparation but simply sizing the fragment according to the analytical method while samples for characterization methods 3 and 4 were hot mounted in resin and metallographically polished up to 250 nm diamond grain size in order to show a longitudinal cross section. The microstructural features were highlighted by etching the polished sample with a water diluted aqua regia (H2O, HCl,

La Metallurgia Italiana - n. 1 2018

(1)

HNO3, vol:vol 35:35:1). The fracture of the samples generated during the tensile test where characterized by SEM (point 4). RESULTS AND DISCUSSION OPTICAL MICROSCOPY The optical metallography helps to define the matrix microstructure and to detect features directly related to the manufacturing process such as mechanical twins and slip bands. The concentration of a deformed matrix observable in the center of the most rolled samples was also observed. It was decided to not quantitatively asses the grain size due to strong microstuctural distortion. Figure 1 shows the most significant microstructures according to the rolling step.

Stainless steel & duplex

Fig. 1 - Metallography, magnifying 500x. Sample: 0.360annealed (a), 0.250 (b), 0.296 (c left), 0.285 (c right), 0.276 (d left), 0.261 (d right)

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Acciai inossidabili e acciai duplex The pictures (a) and (b) relatively show the microstructure of the as bright annealed (a, 0.360mm) and the final cold rolled product (b, 0.250mm). The metallography couples (c) and (d) are the most significant, since they correspond to steps 2 and 6, which have a larger value of Î&#x201D;(def.rate), which entails a greater structural deformation. Especially sample 0.285 compared to sample 0.296 shows a greater amount of alteration with an higher concentration of slip bands and distorted grains. Sample 0.261 compared to sample 0.276 presents areas looking like a coalescence of darker phases (possibly martensite) and much more grains elongated in the rolling direction.

SCANNING ANALYSIS

ELECTRON

MICROSCOPE

FRACTOGRAPHIC

Fractures produced by tensile tests was observed by Scanning Electron Microscope (SEM). The goal of this analysis is to observe the possible presence of fracture surface alteration and austenite according to the different fracture type. Figure 2 shows representative examples of fractography showing evidences of alteration of the classical micro-dimples ductile fracture surface which is typical of the austenitic stainless steels.

Fig. 2 - Frattographies, SEM-SE. Sample: 0.360annealed (a), 0.250 (b)

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Stainless steel & duplex As shwon in the pictures of figure 2, all cold rolled samples froom the annealed to the final product show a mostly ductile fracture surface regardless of the deformation rate. TENSILE TEST This test represents the routine qualification essay to be made on the final product in order to define its mechanical-strength properties. In the specific case of this research tensile tests were

carried out on all intermediate cold rolling steps in order to appreciate the variation of resistance properties at each rolling step. All acquired profiles are shown in Figure 3. The curves are clearly showing a decrease of toughness coherently with the cold rolling progression. The UTS is more and more corresponding to the breaking position and the YS (measured at 0.2 of permanent deformation) is also increasing with a limited variation of the Young modulus.

Fig. 3 - Tensile curves of each rolling step of the second cycle

MAGNETOMETER SQUID M VS H MEASURE Magnetic measurements with dc-SQUID instrument were carried out in order to measure amount (volume fraction) of the magnetic phase (i.e. martensite) in the austenitic matrix. The ferromagnetic phase quantification was performed using a theoretical Msaturation value of a standard AISI 301 steel, theoretically containing 100% martensite. To quantify martensite is necessary to assess the sample saturation magnetization.

However, the small contribution of austenitic phase magnetic response (the paramagnetic phase) is to be taken into account: the value of austenitic magnetization must be eliminated. For this purpose the linear portion of the sample chart with lower Msaturation is plotted, i.e. the sample containing the greater austenite amount (paramagnetic phase). This value was subtracted from each sample, so that the obtained profile is solely referred to martensite magnetization. The martensite value % is calculated by Msaturation, as shown in table 4.

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Acciai inossidabili e acciai duplex Tab. 4 - Msaturazion graphic interpolation; infinite magnetic field applied.

Sample

Msat.(A/m)

Q.tymart.(%)

Msat. 100%martensite (A/m)

0.360

3.07E+04

2.3

1.31E+06

0.330

3.51E+04

2.7

0.295

1.28E+05

9.7

0.285

2.02E+05

15.4

0.280

2.54E+05

19.4

0.276

3.18E+05

24.2

0.261

2.96E+05

22.6

0.254

3.71E+05

28.3

0.252

4.17E+05

31.8

0.250

4.68E+05

35.6

Comparing the estimated martensite volume fraction with the deformation rate (fig. 4) and then with the data acquired by the tensile tests (fig. 5) it is possible to draw some interesting sug-

gestions for the discussion. Figures 4 and 5 graphically describe a tentative combination of data.

Fig. 4 - Correlation between martensite amount and strain rate; red dots = high Î&#x201D;(def.rate); blue dots = low Î&#x201D;(def.rate).

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Stainless steel & duplex

Fig. 5 - YS (Rp0.2) and martensite amount correlation; punti red dots = high Δ(def.rate); blue dots = low Δ(def.rate). During the cold deformation the austenitic lattice has two main microstructural modifications: the formation of martensite and the hardening process (e.g. stacking faults, mechanical twins, slip bands). The importance, in terms of how the metal is af-

fected, of the two phenomena is strongly depending from the way the stress is applied on the metal, in other words, from the experimental conditions (eq. 2)

EDeformation = EQ + EHardening + EMartensite + Deformation where: EDeformation = Energy released during cold rolling; EQ = Energy lost as heat; EHardening = Energy absorbed as work hardening; EMartensite = Energy used to produce martensite; Deformation = plastic deformation. The martensite amount exponentially increases as a deformation rate function. In particular high Δ(def.rate) steps are part of the exponential curve, while low Δ(def.rate) steps deviate from the aforementioned trend. This leads to the hypothesis that small Δ(def.rate) involves a deformation mechanism other than the one dominating when there is a large Δ(def.rate). The usage of the quantitative analysis results gathered by the SQUID magnetometer it was possible to express the YS can be expressed as martensite amount function, as shown in Figure 5, and thus to estimate if a deviation similar to that of figure 4 was reported. It results that YS, which roughly indicates the elastic limit for these steels, does not vary linearly as a martensite function, as should be expected, but is affected indeed by by a combination of martensite and a secondary effect related to the cold rolling process. In this case it is coherent to suggest that the energy accumulated in the lattice as defects and usually defined under the generic term of work hardening contributes to the definition of the mechanical properties of the steel. This leads to the formulation of the hypothesis that would explain the different transformation involving small Δ(def. 18

(2)

rather than high in the martensitic contribution to the YS: at low Δ(def.rate) corresponds YS increase. rate)

Equation (2) defines how the energy applied on the metal by cold rolling is handled by the microstructure, the graphs of figures 4 and 5 how the microstructural and mechanical properties are affected by it. The fact that the strain density (amount of deformation applied per rolling step) has to be seriously taken into account becomes thus a key aspect for further interpretation. As one of the main issues of this research work there is the evidence that martensite and work-hardening results from a distribution of the absorbed energy strictly related to the application rate. In particular it has been experimentally noticed that: - Low Δ(def.rate) are prone to generate a greater martensite amount, which is formed on energy costs, accumulated during hardening. So the martensite increment is balanced by an energy decrease from work hardening, resulting in only a modest increase in YS. - High Δ(deformation rate) produce both martensite and hardening, such that the martensite formation is not able to balance the work hardening, producing a steady increase of YS.

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Acciai inossidabili e acciai duplex Conclusion The experimental sessions and data gathered during the present research generated the awareness that not enough attention was previously paid to the work hardening of the austenite by accumulation of lattice defects during the manufacturing process. The formation of martensite confirmed to be the most important and evident phenomenon triggered by mechanical stress of the metastable austenitic matrix however the tensile curves measured on several samples issued by the cold rolling

process have shown that the martensite volume fraction has to be considered beside other matrix strengthening processes in order to justify the changes in YS. Further researches should be made in this direction in order to offer a sound contribute to reach the mature knowledge needed to refine the predictive models of Austenitic stainless steels behaviour under stress.

REFERENCES [1]

Kohyama, M.L. Grossbeck, G. Piatti, Journal of Nuclear Materials, 191â&#x20AC;&#x201C;194, Part A, (1992), p.37

[2]

Parag M. Ahmedabadi, Vivekanand Kain, Ashika Agrawal, Materials & Design, 109, (2016), p.466

[3]

G. Fargas, J.J. Roa, A. Mateo, Wear, 364â&#x20AC;&#x201C;365, (2016), p.40

[4]

I.R. Souza Filho, M.J.R. Sandim, R. Cohen, L.C.C.M. Nagamine, J. Hoffmann, R.E. Bolmaro, H.R.Z. Sandim, Journal of Magnetism and Magnetic Materials, 419, (2016), p.156

AKNOWLEDGMENTS A special credit has to be paid to dr. C. Froso from AP Steel s.r.l., who have actively contributed to the research procedures.

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Stainless steel & duplex Mechanical properties evolution on heat treated severe cold rolled UNS S32760 Super Duplex Stainless Steel C.M. Tromellini, A.F. Ciuffini, A. Gruttadauria, S. Barella, C. Di Cecca, C. Mapelli Super Duplex Stainless Steels (SDSSs) are characterized by the compresence of austenitic and ferritic grains. They display very high corrosion resistance, especially in chloride-rich environments. They also have better mechanical properties than single-phase counterparts. In the recent years, many studies have been performed to enhance the mechanical properties of SDSS without detrimental effects on corrosion resistance. In this work, Super Duplex Stainless Steel AISI F55 – UNS S32760 is cold rolled and heat treated at different temperatures and holding times to obtain a serious enhancement of the mechanical properties of the material. Much attention is devoted to the optimization of the heat treatment performed on the plastically deformed material, in order to obtain a suitable compromise between tensile strength and ductility, without the formation of undesired embrittling phases. The heat treatment at 620° C was considered the best compromise, combining 1539 MPa of tensile strength with 8,25 % of fracture elongation after 12,5 minutes of holding time. A detailed microstructural analysis has been carried out to understand the strengthening mechanisms acting on the material and to explain the evolution of the microstructure that leads to superior mechanical properties of the treated steel. The corrosion resistance of the heat treated material has been checked for the most promising thermo-mechanical treatments.

KEYWORDS: SUPER-DUPLEX STAINLESS STEELS (SDSS), ANNEALING THERMAL TREATMENT, COLD ROLLING

INTRODUCTION Super Duplex Stainless Steels (SDSSs) are characterized by the compresence of austenitic and ferritic grains. They display very high corrosion resistance, especially in chloride-rich environments and have better mechanical properties than singlephase counterparts. Their unique combination of properties makes them the material of choice for many industries and the ever growing demand of better performing components for industrial and consumer applications ensure that the market for Super Duplex Stainless Steels will grow in the next future. Many studies have been performed to enhance the mechanical properties of SDSSs without detrimental effects on corrosion resistance. [0103] Cold plastic deformation with high reduction ratio followed by heat treatment is a well-known processing route to obtain a material with superior mechanical performances. [04] The main purpose of this work is to document the effects of different heat treatments on the mechanical properties of a cold rolled UNS S32760 steel. The solution annealed material was cold rolled up to 88% of reduction. The rolled material was heat treated at five different temperatures, 550° C, 620° C, 700° C, 800° C and 1080° C with different holding times, ranging from 1,5 minutes to 60 minutes. Much attention was devoted to the optimization of the heat treatment performed on the plastically deformed material, in order to obtain a suitable compromise 20

between yield strength and ductility, and avoid the formation of embrittling phases, deleterious both for mechanical and corrosion resistance. Tensile tests were performed to assess the mechanical resistance of the material. A detailed microstructural analysis has been carried out to understand the strengthening mechanisms acting on the material and to explain the evolution of the microstructure that leads to the superior mechanical properties of the treated steel. Beside the enhanced mechanical properties, the corrosion resistance of the material has to be carefully checked to ensure a very limited loss of performance, if any. For that reason, the corrosion resistance of the heat-treated material has been checked for the most promising thermo-mechanical treatments.

C.M. Tromellini, A.F. Ciuffini, A. Gruttadauria, S. Barella, C. Di Cecca, C. Mapelli Politecnico di Milano, Italy

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Acciai inossidabili e acciai duplex EXPERIMENTAL PROCEDURE The material used in this work is a F55 – S32760 super duplex stainless steel with chemical composition accordant to the ASTM A240/A240M – 04 standard. Samples were cut from the ingot in the transversal direction, in order to obtain 10 mm thick slices. Then, the specimens were cold rolled using a lab-scale rolling mill to a final thickness of 1,2 mm, through many rolling steps. The samples were cooled during rolling, in order to keep the temperature of the material well below 250° C, which can be considered the limit working temperature for this material, preventing the precipitation of embrittling phases. [05]

The samples were loaded in a preheated muffle furnace and water quenched after the required holding time. Because of the reduced thickness of the specimens, the heating and cooling rates are very fast. This is important to isolate the effect of temperature and holding time on the material. The thermal treatments temperatures were: 550° C, 620° C, 700° C, 800° C, 1080° C. The soaking times of the thermal treatments vary from a minimum holding time of 2 minutes up to 60 minutes. (Table 1)

Tab.1 - Compositions (wt%) of the samples and, determined with SEM-EDS for the major alloying elements, and with EPMA for nitrogen.

THERMAL TREATMENTS SUMMARY Temperature [° C] Holding time [min] 550 1,5 2 5_2 3 5 5_5 7,5 5_7,5 10 5_10 12,5 5_12,5 15 5_15 20 5_20 25 5_25 30 5_30 40 5_40 60 5_60

620 6_2 6_3 6_5 6_7,5 6_10 6_12,5 6_15 6_20 6_25 6_30

In order to assess the mechanical resistance of the investigated material, tensile tests have been executed following the ASTM E8/E8M – 16 standard. Potentiodynamic polarization testing is an electrochemical technique used to study the corrosion process of a metal alloy. In this work it has been performed following the ISO 17475 reference norm. A “three-electrode” Amel 2553 potentiostat was used. The working electrode was a sample of the material under examination. Two counter electrodes made of platinum were used to close the circuit, and a “saturated calomel electrode (S.C.E.)” was adopted as reference electrode. The electrode is included in a “Luggin capillary”. All the above-mentioned components were submerged in an electrically conductive solution, the electrolyte, which properly simulates the aggressive environment where the steel has to work (marine environment). The solution was made by 35 g/L (0.6 M) of pure NaCl in distilled water. The voltage range was bounded between -600 and 1400 mV and the scan rate was set to 0,04 mV/s. All the tests were performed at La Metallurgia Italiana - n. 1 2018

700 7_1,5 7_2 7_3 7_5 7_7,5 7_10 7_12,5 7_15

800 7_1,5 7_2 7_3 7_5 7_7,5

1080 10_1,5 10_2 10_3 10_5 10_7,5

room temperature. The surface of the samples was prepared according to the norm, in order to have an average roughness smaller than 1 micrometer. They were grinded by means of an abrasive paper up to 800 grit. The X-ray diffraction method was used to assess the amount of the austenite volume fraction in the cold rolled and heat treated samples. A Stresstech Xstresses 3000 G3R X-ray diffractometer was used. The samples were mechanically grinded and polished up to a 1 micron diamond paste, to have a clear surface and to avoid the plastic deformation of the surface, which can generate strain-induced martensite and produce misleading results. The Kα radiation of Chromium was used as X-ray source, and the calculation of the phase fraction of austenite were carried out according to the ASTM E975-03 Standard Practice for X-ray Determination of Retained Austenite in Steel with Near Random Crystallographic Orientation. In this work, the evaluation of the diffraction peaks of the austenite at 130° and 80° degrees was carried out by using the “parabolic” function to reduce the background noise and 21

Stainless steel & duplex the “Gauss” function to fit the curve. For the diffraction peaks of the ferrite at 156,4° and 106,1° the Pearson VII function was used to fit the peaks and a linear function for noise reduction. The measure of the austenite content is based on the calculation of the areas under the diffraction peaks of the ferritic and austenitic phases, as these areas are proportional to the content of each phase.

the basis of the temperature of the heat treatment. The asrolled material was tested to assess the strain-hardening of the alloy subjected to cold rolling, up to a high reduction of the cross section. Heat treated samples show a further increase in yield stress, tensile strength and also fracture elongation with respect to the as-rolled condition. This leads to impressive results for the mechanical resistance.

RESULTS Tensile tests results are reported in the following, grouped on

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Acciai inossidabili e acciai duplex

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Stainless steel & duplex

Fig. 1 - Tensile tests results

The corrosion resistance results, extrapolated from the potentiodynamic polarization tests, have been reviewed in Table 2. It can be observed the slight loss in corrosion resistance related to the thermal treatments. The data related

to the samples 7_12,5 and 10_7,5 are not detailed since a passive behavior has not been detected.

Tab. 2 - Potentiodynamic polarization results

Sample As-received As-rolled 5_5 5_7,5 5_12,5 6_5 6_7,5 6_12,5 7_5 7_7,5 7_12,5 10_7,5

CORROSION RESISTANCE Current density [A/m2] 1,3 * 10^-2 4,6 * 10^-2 1,62 * 10^-1 1,81 * 10^-1 2,01 * 10^-1 3,41 * 10^-1 3,75 * 10^-1 2,72 * 10^-1 1,18 2,86 ---

Pitting potential [V vs SCE] 0,95 0,94 0,93 0,89 0,92 0,9 0,92 0,92 0,91 0,38 ---

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Acciai inossidabili e acciai duplex The X-ray analysis results, revealing the austenite content of the specimens after the thermos-mechanical treatments,

have been reported in Table 3.

Tab. 2 - Potentiodynamic polarization results

Sample As-received As-rolled 5_5 5_7,5 5_12,5 6_5 6_7,5 6_12,5 7_5 7_7,5 7_12,5 10_7,5

AUSTENITE CONTENT γ phase [%] 35,7 22,6 20,2 18,7 18,6 19,7 26 21,5 26,8 27 30,6 37,9

DISCUSSION When subjected to severe plastic deformation the material undergoes serious changes in the mechanical properties. Two main processes occur in the duplex structure under these conditions: • strong grain fragmentation, that leads to a strong microstructure refinement. • nucleation of strain-induced martensite, by the γ → α’ reaction. Both these processes are responsible for a high work-hardening and loss of fracture elongation, but, as no diffusion of chemical species takes place during cold deformation, corrosion resistance should not be greatly influenced. This is perfectly adherent to the results obtained on as-rolled sample. To assess the amount of austenite transformed into straininduced martensite was used the direct comparison method described by Cullity et al. [06]. The result is around 13 % of strain-induced martensite, that is coherent with other results in literature. During a heat treatment at around 550° C many changes in the microstructure of the material takes place: • precipitation of α’Cr phase either by nucleation and growth or by spinodal decomposition. • precipitation of G-phase. • precipitation of Cr2N. Precipitation of α’Cr phase is known to occur up to 550° C and plastic deformation can change the nature of the reaction δ → α + α’Cr from nucleation and growth to spinodal decomposition. The decomposition of ferrite is also responsible for enhanLa Metallurgia Italiana - n. 1 2018

Error [%] 0,2 0,4 0,2 0,1 0,3 0,3 0,5 0,2 0,2 0,3 0,5 0,3

cing the precipitation of Cr2N, which is frequently reported in case of spinodal decomposition of ferrite. [07] G-phase does normally precipitates after ferrite decomposition, because is rich in alloying elements that segregates during the decomposition process. Preferential sites for G-phase nucleation are dislocations, probably due to the enhanced diffusion through the dislocation length. The X-ray analysis reveal that the amount of austenite after 5 minutes of holding time is already lower with respect to the as-rolled material, and keeps lowering up to 7,5 minutes. After 7,5 minutes of holding time the amount of retained austenite stays constant up to 12,5 minutes (Tab. 3). These results, combined with the results of the tensile tests, suggest that most of the precipitation processes take place in the first 10 minutes of heat treatment or less. The changes in corrosion resistance are not dramatic, and the material still displays a satisfying stability against Cl-rich environments. This is in agreement with the hypothesis that the main microstructural change occurring in the material is the precipitation of α’Cr phase, which is known to have a weak effect on corrosion resistance. This also may discourage the hypothesis of Cr2N precipitation, as they are known to have small impact on mechanical properties but can seriously decrease the corrosion resistance. [08] During heat treatment at 620° C there are four main reactions which can take place in super duplex stainless steels, especially if they are plastically deformed: • strain-induced martensite reversion to austenite. • static recovery. 25

Stainless steel & duplex • precipitation of secondary austenite through martensiticshear process. • precipitation of R-phase. • precipitation of δ phase. The first and the second processes may explain the softening of the material by the nucleation of austenite. The X-ray analysis of the austenite content may confirm the trend. Especially between 5 and 7,5 minutes, the amount of austenite in the material has a noticeable increase, weather after 12,5 minutes of holding time is almost the same amount as in the as-rolled material. This oscillation in the amount of retained austenite can be explained by the local phase transformations and the diffusive processes controlling the kinetics of the process (Tab. 3). [09] Static recovery is thought to take place in the ferritic phase, but has a very marginal effect on fracture elongation and tensile stress. This fits the experimental results that show a very high tensile stress and only moderate increase in fracture elongation, with respect to the “as-received” condition. The precipitation of R phase is known to occur at maximum rate around 600° C, and cold working significantly favors the precipitation of this intermetallic phase. R phase has a limited impact on tensile strength, whether is known to reduce the fracture elongation. [08,10] Analyzing the corrosion resistance results, they may suggest that a microstructural transformation happens already at 5 minutes of T. T. holding time, as there is a change in both passive current and shape of the overall potentiodynamic curve. This change may be also related to the R phase precipitation, as shown by Kim et al. [08]. Combining the two results it may be assumed that the variation of tensile properties is due to the combined effect of γ2 and R phase nucleation, which lead to a best combination around 12,5 minutes of holding times. For longer holding time the coarsening of R phase leads to an overall decrease of the mechanical properties of the material. During heat treatment at 700° C the corrosion resistance sharply decreases, where the mechanical properties do not change so drastically, even though a certain decreasing trend in Rm and Rp0,2 may be present. This behavior is explained with the precipitation of Chromium carbides and nitrides, because they have a small impact on mechanical properties but can seriously compromise the corrosion resistance of the alloy. They are known to precipitate around 700° C, but because of the very low amount of Carbon present in the analyzed alloy, the reduction of corrosion properties may be related to nitrides precipitation only. The shape of the stress strain curve on the other hand shows quite a change as the heat treatment goes on, and for longer holding times the strain hardening is much more evident. This behavior may be coupled with the results from the X-ray evaluation of the austenite content, which shows a continuous increase in austenite content with holding time. It may be assumed that the main reaction responsible for the increasing amount of austenitic phase should be σ → γ2 + σ, as the eutectoid reaction is known to be favored by cold working. The precipitation of σ phase can be also responsible for 26

losses in corrosion resistance, but cannot be claimed to be the only one, as the mechanical properties of the material do not follow such a fast degradation, which is expected since both properties are strongly affected by this precipitation process. [08] During aging at 800° C there are many precipitation processes that may occur. The first is the precipitation of Chromium nitrides, at phase boundary between austenite and ferrite. After that, the precipitation of χ phase and σ phase, which will form through the eutectoid reaction δ → γ2 + σ, takes place, and seriously affect the mechanical properties of the material. It is also known that cold working shortens the incubation time for σ phase precipitation and increases the precipitation kinetics. In effect, it has been reported dislocation can act as nucleation site for σ phase. [11] The results of the mechanical testing show a fast and serious degradation of the material, especially the fracture elongation, which may be related to the ductility of the material. This implies that the precipitation inside the material must be relevant, and the only reaction able to compromise the properties of the material to such an extent is the eutectoid reaction δ → γ2 + σ. At 1080° C, the only stable phases are ferrite and austenite, which implies that the precipitation of any other phase may not takes place during the heat treatment. Moreover, at this temperature recovery, recrystallization and coarsening of ferritic and austenitic phase occur. This would bring back the material to the solution annealed condition. A sufficiently intense plastic deformation can induce grain refinement, which is known to enhance the mechanical properties of the material. [12] This is in agreement with the results of the tensile test, that show a sudden decrease in yield stress and tensile strength, together with a noticeable increase in fracture elongation, typical of recrystallization phenomena. Around two minutes of holding time there is also a small peak in the mechanical resistance, that can be related to grain refinement. After 7,5 minutes of holding time, the amount of austenite inside the microstructure of the material is almost the same as in the as-received case, that suggests no other phases are left after annealing at this temperature, and the material is approaching the equilibrium microstructure (Tab. 3). The results for corrosion resistance show a serious deterioration with respect to the as-received material. This can happen due to the high cooling rate imposed to the material when water-quenched. At temperatures above 1000° C, if the cooling rate is too high, nitrogen cannot diffuse towards γ phase, and a supersaturated ferritic phase give rise to precipitation of very fine Chromium nitrides inside the ferritic grains. These nitrides have almost no influence on the mechanical properties, but do affect the corrosion resistance because they create a small zone around them depleted both in Cr and N, and for this reason susceptible to localized corrosion attack. [13] This can explain the loss of corrosion resistance in the samples heat treated at 1080° C.

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Acciai inossidabili e acciai duplex CONCLUSION In this work, the mechanical properties of super duplex stainless steels have been enhanced, preventing losses in corrosion resistance. The material was cold rolled and heat treated at different temperatures and for different holding times. The most important results can be summarized as follows: • The as-rolled material shows a great enhancement of tensile strength, as result of the formation of strain-induced martensite during the rolling process. The corrosion resistance is almost the same as the solution annealed material. • The heat treatment at 550° C gives the best results for yield strength (Rp0,2=1645 MPa) and tensile strength (Rm=1726 MPa) at around 12,5 minutes of holding times, but the material shows very limited strain-hardening. The corrosion resistance of the alloy is still very good, even though there is a limited loss with respect to the as-rolled condition. The decomposition of the ferritic matrix along with the precipitation of G phase are the responsibles for this change in material’s behavior. • The heat treatment at 620° C leads to the best compromise in terms of mechanical properties and corrosion resistan-

ce. Around 12,5 minutes of holding time the material shows excellent yield stress (Rp0,2=1407 MPa) and tensile stress (Rp=1539 MPa), a reasonable fracture elongation (Ax=8,25 %) and a certain strain-hardening behavior. The corrosion resistance is lower with respect to the as-rolled material but still shows a clear passive behavior. The reversion of strain-induced martensite and a certain recovery, alongside with precipitation of R phase and secondary austenite are thought to be the main processes acting during treatments at this temperature. • At 700° C and 800° C the corrosion resistance is negatively affected by the precipitation of Chromium nitrides and/ or σ phase. Further, the mechanical properties obtained do not compensate such losses, precluding any possible application. • The 1080° C treatment restores the mechanical behavior of the solution annealed material by promoting recrystallization. There is an increase of the fracture elongation (Ax=35,58 %) and a certain increase in tensile stress (Rm=954 MPa). However, the material shows poor corrosion resistance due to the precipitation of Chromium nitrides, generated by an excessive cooling rate.

REFERENCES [01]

R.N. GUNN, Duplex Stainless Steels, Microstructure, properties and applications, Woodhead Publishing (1997).

[02]

I. ALVAREZ-ARMAS and S. DEGALLAIX-MOREUIL, Duplex stainless steels, ISTE (2009).

[03]

G. FARGAS et Al Journal of Materials Processing Technology 209 (2009), p.1770.

[04]

S.-H. KIM et Al., Nature 518 (2015) p. 77.

[05]

F. TEHOVNIK et Al., Materials and technologies/Materiali in tehnologije 45 (2011).

[06]

B. D. CULLITY et Al., American Journal of Physics 25 (1957) p. 394.

[07]

A. WEIDNER et Al. Metallurgical and Materials Transactions A 47 (2016) p. 2112.

[08]

J.-H. KIM et Al., Journal of Materials Engineering and Performance 25 (2015) p. 9.

[09]

M. MARTINS et Al., Materials Characterization 59 (2008) p. 162.

[10]

B.-C. LEE et Al., Korean Journal of Materials Research 24 (2014) p.

[11]

T. BERECZ et Al., Journal of Materials Engineering and Performance 24 (2015) p. 4777.

[12]

L. JINLONG et Al., Materials Science and Engineering: C 62 (2016) p. 558.

[13]

N. PETTERSSON et Al., Metallurgical and Materials Transactions A 46 (2015), p. 1062.

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Stainless steel & duplex Different effects of carbon and nitrogen on precipitation behavior and mechanical properties in austenitic stainless steels Kyung-Shik Kim, Jee-Hyun Kang, Sung-Joon Kim Austenitic stainless steels, which are the most common type of stainless steels have good weldability and formability, and can be used in a wide range of temperatures. However, precipitation of carbides and nitrides may have detrimental effects on their mechanical properties. Fe-15Cr-15Mn-4Ni based steels were investigated in this research with different contents of 0.2C, 0.2N and 0.2C+0.2N. Precipitation behavior was observed by aging the specimens in temperature range from 600 OC to 1000Â OC for up to 240 hours. Precipitates of M23C6 and Cr2N were formed. The specimens containing 0.2% of carbon formed M23C6 along grain boundaries and within grains depending on the aging condition. Cr2N, however, was formed only at grain boundaries in the specimens containing 0.2% of nitrogen. The specimens containing both carbon and nitrogen formed both carbide and nitride, the fractions of which are different. TTP (Time-Temperature-Precipitation) diagrams were constructed based on the observation. Moreover, tensile test was conducted to identify the effect of precipitation on mechanical properties. Specimens with carbon showed more decrease in elongation and increase in ultimate tensile strength (UTS), whereas ones with nitrides showed no distinct change in elongation and UTS. The different effects of nitrogen and carbon are discussed in detail.

KEYWORDS: AUSTENITIC STAINLESS STEEL, PRECIPITATION, CARBIDE, NITRIDE, MICROSTRUCTURE, MECHANICAL PROPERTY

INTRODUCTION Austenitic stainless steels (ASSs) are widely used due to their excellent corrosion resistance, toughness and formability. 70% of stainless steel market demand is satisfied with AISI 300 series austenitic stainless steels which contain 18 wt.% of chromium and 8-12 wt.% of nickel. However, the price of nickel is high and unstable. Therefore, replacing nickel with manganese has been actively investigated. Furthermore, carbon and nitrogen are added, since both nitrogen and carbon are strong austenite stabilizers and effectively enhance the strength and corrosion resistance of material [1-2]. However, the addition of these interstitials induces the precipitation of carbides and nitrides at high temperatures [3]. Precipitates are formed at grain boundaries in most cases, which can degrade the mechanical properties of materials [4-5]. As a consequence, research on the precipitate kinetics is important for the austenitic stainless steels containing high interstitials. Carbon and nitrogen tend to distribute differently in interstitial sites of a steel; carbon prefers to cluster, i.e. shortrange decomposition, while nitrogen never locates as a first neighbor to each other, i.e. short-range ordering, when they are in solid solution. Moreover, the effect of simultaneous alloying of carbon and nitrogen is not additive, and the au28

stenitic steel with C+N are even more stabilized than that with either carbon or nitrogen of the equivalent concentration [6-8]. Such different atomic distribution results in the stability variation of carbon and nitrogen in solid solution, which would affect the thermodynamics and kinetics of carbide / nitride precipitation. In the present study, the precipitation behavior of manganese containing austenitic stainless steels with either carbon or nitrogen is investigated at a 600-1000 OC temperature range. The effect of the precipitates on mechanical properties is also studied by tensile test at room temperature.

Kyung-Shik Kim, Jee-Hyun Kang, Sung-Joon Kim Graduate Institute of Ferrous Technology (GIFT), POSTECH, Republic of Korea

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex EXPERIMENTAL PROCEDURE

Tab.1 - Compositions (wt%) of the samples

CHEMICAL COMPOSITION (wt.%) Alloy

Base

15.4

14.8

4.1

0.33

0.002

0.008

14.9

15.0

4.0

0.29

0.208

0.004

15.0

15.2

4.1

0.26

0.015

0.192

CN20

15.1

15.3

4.0

0.33

0.210

0.195

The compositions of four alloys, in this study, Base, C2, N2 and CN20 are listed in Table 1. A 15Cr-15Mn-4Ni based steel can contain carbon and nitrogen up to 0.3 wt.% in solid solution and retain a full austenitic microstructure at room temperature. Thermodynamic calculations were carried out by MatCalc version 6.00 with mc_fe database to assess the equilibrium phases and the driving force for precipitates. 25 kg ingots of four alloys were heat treated at 1200 OC for two hours and hot rolled to 12 mm. Plates were then cut and annealed at 1100 OC for 30 minutes to obtain a full austenite microstructure. The grain size of austenite matrix for each alloy was 52 ± 6 μm for Base, 89 ± 5 μm for C2, 62 ± 5 μm for N2, and 56 ± 4 μm for CN20. After cutting the samples to 1.2 × 0.5 × 1.5 mm3, they were aged at temperatures from 600 to 1000 OC with a 100 OC interval. The aging times were 10, 30, 60, 180, 600, 1440, and 14400 minutes. Precipitates were analyzed with FE-SEM, EDS and TEM. SEM samples were prepared by mechanical polishing which was finished with colloidal silica. Back scattered electron (BSE) images were analyzed to produce time-temperature-precipitation (TTP) diagrams and to calculate the distribution of precipitate. FE-SEM EDS and TEM were employed to identify the precipitates. By detecting the different EDS signals of carbon and nitrogen in the precipitates, carbides and nitrides could be qualitatively distinguished. Detailed classification of the precipitates was performed by analyzing diffraction patterns taken by TEM. TEM specimens were prepared by jet polishing at 20.5 V with 10% perchloric acid mixed with 90% acetic acid as an electrolyte after mechanical polishing to a thickness of 100 μm.

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Based on the observed precipitation behavior, tensile tests were carried out to obtain the mechanical properties. The samples aged at 800 OC were used to clearly see the effect of the precipitates on the mechanical properties. Subsized tensile specimens with a thickness of 1.5 mm were machined according to ASTM E8M-4, and tested with constant displacement speed of 1 mm min-1 which corresponds to an initial strain rate of 6.67 × 10-4 s-1. Deformed samples were investigated by EBSD to identify their deformation mechanism. The distributions of phases and twin boundaries were observed in the samples without aging and after aging at 800 OC for 14400 min. The EBSD samples were prepared in the same way with the XRD samples. RESULTS AND DISCUSSION 1. Precipitation behavior Based on the observation of precipitates, TTP diagrams were constructed and shown in Fig. 1. According to the diagram, the generated precipitates largely depend on the alloy composition. Sigma phase and δ-ferrite were formed in Base, intergranular and intragranular M23C6 in C2, intergranular Cr2N in N2, and both M23C6 and Cr2N in CN20.

Stainless steel & duplex

Fig. 1 - TTP diagrams of (a) Base, (b) C2, (c) N2, and (d) CN20.

General morphologies of the precipitates after aging at 700 OC for 14400 min are shown in Fig. 2, where the dark particles in the images are the precipitates. Light elements give dark contrast in the BSE mode, and therefore, all M23C6, Cr2N, sigma phase which are enriched with lighter elements than Fe appear darker than the austenitic matrix. These precipitates are distinguished by their selected area diffraction patterns taken by TEM. Some examples are shown in Fig. 3. Both intergranular and intragranular M 23C6 showed

orientation relationship with the austenite matrix; however, sigma and nitride particles did not show any orientation relationship. The observed existences of different precipitates agree with thermodynamic calculation results in Fig. 4 where phase fractions of austenite, M23C6, Cr2N, sigma, and Î´-ferrite are shown according to temperature.

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Acciai inossidabili e acciai duplex

Fig. 2 - BSE images of (a) Base, (b) C2, (c) N2, (d) CN20, aged at 700 OC for 14400 min.

Fig. 3 - TEM dark field image and diffraction patterns of (a) sigma phase in Base after 700OC aging, (b) grain boundary M23C6 and (c) intragranuluar M23C6 in C2 after 900OC aging, and (d) Cr2N in N2 after 700OC aging. All photos are taken after aging for 14400 min.

La Metallurgia Italiana - n. 1 2018

Stainless steel & duplex

Fig. 4 - Phase fraction of (a) Base, (b) C2, (c) N2, and (d) CN20. Note that the phase fractions are given in log scale. In Base, δ-ferrite was observed at all temperature range (Fig. 1a). It seems that the δ-ferrite has formed during heat treatment before hot rolling or during hot rolling process as δ-ferrite is one of the equilibrium phases above 1000 OC (Fig. 2a). δ-ferrite is known to transform into sigma after heat treatment, producing a mixture of ferrite and sigma over the austenitic matrix [9]. Therefore, intergranular sigma phase was formed during the aging process at 600-700 O C. The sigma phase, either from austenite or δ-ferrite, can be easily distinguished from the preexisting delta by their location. Carbides and nitrides were not observed in this alloy because their estimated phase fractions are less than 0.1% (Fig. 1a). Due to a complicated matrix containing both austenite and δ-ferrite, the mechanical properties of Base are not studied. In C2, M23C6 formed at all temperatures between 600 and 1000 OC (Fig. 1b). This is consistent with the estimated phase fraction in Fig. 4b. M 23C6 always began to precipitate at grain boundaries, and spread into inside grains afterwards. Intragranular carbides were formed only at the temperatures over 700 OC. Nitrides did not form because the C2 alloy contains a negligible amount of nitrogen.

N2 showed Cr 2N at grain boundaries after aging under 800 O C, which corresponds to the equilibrium phase fraction (Fig. 4c). However, unlike other alloys, the precipitates were barely observed due to the low density. Carbides were not observed in this alloy because, as indicated in the Fig. 4c, the phase fraction of M23C6 is negligible (< 1%). Both carbide and nitride are detected in CN20 as in the calculation results in Fig. 4d. The carbides were formed along grain boundaries, and later inside grains. In comparison with C2, the temperature range for the intergranular carbides decreased from 1000 OC to 900 OC, and the formation was delayed under 700 OC. Intragranular carbides were formed only at 800-900 OC after 1440 min of aging. It is clear that the carbide formation is hindered by the introduction of nitrogen, which agrees with the previous studies [10]. Nitrides, on the other hand, were formed faster at 600 OC, and similar in 700-800 OC when compared with N2. Therefore, it seems that carbon promote the nitride formation. Sigma phase is expected to form more at 600 OC than M23C6, and Cr2N in the thermodynamic calculation (Fig. 4). In real observation however, they were not found in C2, N2 and CN20. It seems that sigma phase is suppressed by the for-

La Metallurgia Italiana - n. 1 2018

Acciai inossidabili e acciai duplex mation of carbide and nitride. Precipitate distributions have been obtained for the alloys aged at 800 OC where most precipitates appear (Fig. 1). The size evolutions of precipitates are similar for all C2, N2, and CN20 regardless of precipitate types and locations, as

shown in Fig.5. The size of carbides and nitrides increase to about 0.35 Îźm after 14400 min aging. The carbides formed inside grains also showed similar size as other intergranular precipitates although intragranular carbides always appear later than intergranular ones.

Fig. 5 - Size change of the precipitates in (a) C2, (b) N2, (c) CN20, with respect to the aging time at 800 OC.

2. Tensile behavior Engineering stress-strain curves of the alloys aged at 800 O C are plotted in Fig. 6, and their yield strength, ultimate tensile strength, and uniform elongation are given in Fig. 7 as a function of the aging time. It is observed that the

tensile properties are differently affected by aging at 800 O C depending on the addition of either carbon or nitrogen. Such phenomenon is closely related to the precipitation behavior.

Fig. 6 - Stress-strain curves of (a) C2, (b) N2, and (c) CN20 aged at 800 OC for different time periods.

La Metallurgia Italiana - n. 1 2018

Stainless steel & duplex

Fig. 7 - Evolution of (a) yield strength, (b) ultimate tensile strength, and (c) uniform elongation of C2, N2, and CN20 after aging at 800 OC.

Before aging, all carbon and nitrogen atoms are in solid solution for C2, N2, and CN20. As yield strength increases with interstitial contents due to solid solution strengthening, highly alloyed CN20 exhibits the highest yield strength. However, about 50 MPa difference of yield strength exists between C2 and N2 although they have almost the same

concentration of interstitial elements ~0.21 wt.% (Table 1). Two factors are most likely to be responsible for such difference: (i) different contributions of carbon and nitrogen to solid solution strengthening, (ii) different grain sizes. Yield strength (YS) of austenitic stainless steels can be estimated by

(1)

where the elements denote their concentration in wt.%, δ is δ-ferrite content in vol.%, d is the grain size of austenite, and kHP is the Hall-Petch coefficient [11-12]. According to equation (1), the contribution of nitrogen to yield strength is about 1.39 times larger than that of carbon. Considering the carbon and nitrogen concentrations in C2 and N2 (Table 1), the estimated YS difference according to equation (1) is 28 MPa. Moreover, taking kHP = 14.723.2 MPa mm-1/2 [13], the difference in grain boundary strengthening is 10-14 MPa. By adding the strengthening effects of solid solution and grain boundary, estimated yield strength difference is 38-42 MPa, which explains the difference in the experimental values. Yield strengths of all three alloys slightly decrease with aging time. Since no change is detected for the grain size of austenite matrix during aging, the yield strength evolution is likely to be related to precipitation strengthening and solid solution strengthening. If precipitation of carbides or nitrides occur during aging, yield strength would increase with aging time due to precipitation strengthening, but it would be reduced by decreasing concentration of interstitials in solid solution, as shown in Fig. 7. Further study is needed to analyze the contribution from each component, and it is being done at the moment. 34

In order to identify the deformation mechanism, phase maps with twin boundaries of deformed samples without and with aging at 800 OC for 14400 min are shown in Fig. 8. Note that all samples maintained full austenitic microstructure before deformation. It is observed that the fractions of deformation induced ε and α’ martensites are higher in C2 after aging in comparison with C2 without aging. For the unaged samples on the other hand, more deformation twinning is activated (Fig. 8a). However, with aging, precipitates form and the deformation occur with martensite formation. As a consequence, uniform elongation decreases with aging. In the case of N2, the densities of Cr2N are much lower although the sizes are similar with carbides (Fig. 5). Thus, the total fraction of Cr2N formed during aging is too low, and the degree of deformation twinning and deformation induced martensite deformation are similar after aging (Fig. 8b, e). Therefore, both ultimate tensile strength and uniform elongation remains similar with aging.

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Acciai inossidabili e acciai duplex

Fig. 8 - EBSD image of deformed samples, unaged (a) C2, (b) N2, (c) CN20, and aged (d) C2, (e) N2, (f) CN20 at 800 OC for 14400 min. Red lines indicate the twin boundaries, green areas are ε-martensite, blue areas are α’-martensite. In CN20, the formation of carbides are inhibited by nitrogen as in Fig. 2d. Therefore, the deformation twinning activity and deformation induced martensite transformation are not affected as much by aging as those in C2. As a result, the ultimate tensile strength experiences a moderate change with aging (Fig. 7b). CONCLUSIONS 1. In the carbon and nitrogen induced 15Cr-15Mn-4Ni austenitic stainless steel, M23C6 carbide and Cr2N nitride and sigma were precipitated. Alloys containing carbon formed M23C6, ones having nitrogen formed Cr2N, and both carbide and nitride were precipitated together in alloy containing carbon and nitrogen. Carbides were formed with higher density than nitride with greater driving force.

2. Addition of nitrogen hindered the formation of carbide. The temperature range and speed of carbide precipitation were reduced with nitrogen. No distinct change in the kinetics of nitride were observed with adding carbon, although at 600OC, nitride formation was accelerated. 3. Tensile properties were differently affected by aging at 800 O C depending on the addition of either carbon or nitrogen. Most changes were observed in C2 with the highest density of precipitates, and N2 had no distinct change in tensile behavior due to low density.

REFERENCE [1] V.G. GAVRILJUK, B.D. SHANINA, and H. BERNS, Mater. Sci. Eng. A, 481-482, (2008), p.707. [2] H.Y. HA, T.H. LEE, C.S. OH, S.J. KIM, Scr. Mater., 61, (2009), p.121. [3] T.H. LEE, H.Y. HA, S.J. KIM, Metall. Mater. Trans., 42A, (2011), p.3543. [4] J.W. SIMMONS, Scr. Matall., 32, (1995), p.265. [5] H. CHANDRA HOLM, P.J. UGGOWITZER, and M.O. SPEIDEL, Scr. Metall., 21, (1987), p.513. [6] V.G. GAVRILJUK, B.D. SHANINA, and H. BERNS, Acta Mater., 48, (2000), p.3879. [7] H. BERNS, V.G. GAVRILJUK, S. RIEDNER, and A. TYSHECHENKO, Steel Research. Int., 78, (2007), p.659. [8] H. BERNS, V.G. GAVRILJUK, and S. RIEDNER, High interstitial stainless austenitic steels, Springer-Verlag Berlin Heidelberg (2013), p. 28 [9] C.C. HSIEH, and W. WU, ISRN Metall., 2012, (2012), 732471 [10] R. PRESSER, J.M. SILCOCK, Metal Sci., 17, (1983), p.241. [11] V.G. GAVRILJUK, and H. BERNS, High nitrogen steels, Springer-Verlag Berlin Heidelberg (1999), p.136. [12] K.J. IRVINE, T. GLADMANM, and F.B. PICKERING, JISI, 199, (1969), p.1017. [13] L.A. NORSTROM, Metal Science, 11, (1977), p.208

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Corso di base

Solidificazione Milano 28 febbraio - 1 marzo 2018 Organizzato da

CENTRO DI STUDIO METALLURGIA FISICA E SCIENZA DEI MATERIALI

La conoscenza dei meccanismi di solidificazione e il loro controllo assumono un ruolo determinante nella produzione metallurgica. Lo scopo del Corso base sulla Solidificazione è, dunque, fornire una panoramica sugli aspetti di base ed applicativi della solidificazione di leghe metalliche. Il corso si rivolge a tecnici di produzione operanti in svariati campi di tipo metallurgico, dalle acciaierie alle fonderie, sia di metalli ferrosi che non ferrosi. L’approccio di base di diverse relazioni, seppure non scollegato dal contesto dei casi reali, rende il corso di sicuro interesse anche per ricercatori in ambito industriale ed accademico. La solidificazione viene generalmente descritta con meccanismi di nucleazione e crescita, la cui combinazione determina la microstruttura dei prodotti di fonderia e di conseguenza le loro proprietà. Nella prima giornata, dopo una presentazione delle caratteristiche strutturali e termodinamiche del metallo liquido, verranno descritti in dettaglio gli elementi essenziali dei processi di nucleazione omogenea ed eterogenea. Nella stessa giornata, verranno presentate le potenzialità delle tecniche di modellazione per la previsione delle condizioni di solidificazione e della formazione dei difetti. Per questa lezione, i partecipanti sono invitati a portare con sé i propri computer portatili. Nella seconda giornata verranno descritti i processi di crescita dei cristalli e di generazione delle microstrutture di solidificazione. Anche in questo caso verranno trattati gli aspetti fondamentali che saranno in seguito applicati ad alcuni casi d’interesse pratico. Il corso terminerà con l’illustrazione di processi di solidificazione di interesse applicativo in campo industriale. Coordinatori del Corso: G. Angella, M. Baricco, R. Montanari Per informazioni ed iscrizioni: AIM · Associazione Italiana di Metallurgia Tel. 02-76021132 / 02-76397770 · E-mail: met@aimnet.it · www.aimnet.it

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La Metallurgia Italiana - n. 1 2018

Industry news New light corrosion resistant steel without chromium a cura di: C. Mapelli, A. F. Ciuffini, S. Barella, A. Gruttadauria, D. Mombelli Low density steels represent a topic of great interest within the scientific world, because of the great demand from the steel market of increasingly lighter materials with optimal mechanical properties. In this work, the interesting potential of the combined addition of Nickel and Chromium in low weight percentage to the Fe-Al-Mn-C class is analyzed. In particular, this study focus on wide assessment of the resistance to chemical degradation phenomena of the Fe-15%Mn-9.5%Al-6.5%Ni-1%Cr-0.43%C alloy, a new steel grade. The aim of this work is to compare this new steel grade with its natural rivals, commercial stainless steels, which have in the resistance to chemical degradation phenomena one of their predominant features. The iso-corrosion curves show excellent resistance of the 1300 °C annealed material, especially at low temperatures but also good at high temperatures. These alloys are promising for the exploitation of this steel grade in those applications where corrosion resistance in chlorine-rich environments or hot oxidation resistance are relevant.

KEYWORDS: LIGHT STEEL – Fe-Mn-Al-Ni STEEL – CORROSION – POTENTIODYNAMIC CURVES

Carlo Mapelli, Andrea Francesco Ciuffini, Silvia Barella, Andrea Gruttadauria, Davide Mombelli Politecnico di Milano - Dipartimento di Meccanica, Italy

INTRODUCTION New steel grades are continuously developed in order to optimize properties such as low density, good mechanical properties, corrosion resistance and high-temperature oxidation resistance. In particular, since 80’s have been investigated the Fe-Al-Mn-C class, which has a great potential as a possible substitute of the Fe-Ni-Cr-C class[1-6]. Nevertheless, up to nowadays the specific tensile strength are still not comparable to Aluminum alloy or Titanium alloy[7-9]. Different studies already demonstrated that the addition of Nickel in the FeAl-Mn-C class has beneficial effects increasing the specific yield strength and making it comparable to Aluminum and Titanium alloys[9]. Further, also Chromium resulted as a promoter of tensile strength and elongation to fracture in previous studies[7,10]. In this work, the interesting potential of the combined addi38

tion of Nickel and Chromium in low weight percentage to the Fe-Al-Mn-C class is analyzed. In particular, this study focus on wide assessment of the resistance to chemical degradation phenomena of the Fe-15%Mn-9.5%Al-6.5%Ni-1%Cr0.43%C alloy, a new steel grade. The aim of this work is to compare this new steel grade with its natural rivals, commercial stainless steels, which have in the resistance to chemical degradation phenomena one of their predominant features. Among chemical degradation phenomena corrosion plays the predominant role. Corrosion is a predominantly surface damage phenomenon of metals, which leads to a loss in weight, a detachment of the material and the consequent malfunction of components. The study of the corrosion resistance is a necessary step in order to have a direct comparison with the stainless steels grades, which are materials of selection for many applications mainly in order to prevent or limit La Metallurgia Italiana - n. 1 2018

Attualità industriale these phenomena. In order to derive useful values of the corrosion resistance of this low density steel grade, it has been chosen the realization of iso-corrosion curves, which represents an effective tool for the evaluation of the general corrosion resistance of a metal or an alloy, taking into account the overall corrosive phenomena [11]. The response of the low density Fe-15%Mn-9.5%Al-6.5%Ni1%Cr-0.43%C alloy against a different degradation phenomenon has been investigated in this work. The chemical corrosion, or to hot oxidation, is a degradation phenomenon of the materials that occurs in gaseous environments at high temperature. It is also defined as dry corrosion and occurs in the absence of aqueous solutions. The first leads to the formation of metal ions, and it occurs at the metal-oxide interface. On the other hand, the reduction half-reaction produces oxygen ions and it occurs at the oxidegas interface. Common practice in typical applications of steels resistant to this phenomenon (such as furnaces and industrial reactors, facilities for combustion processes ...) is to use: common steels up to 550 °C, chromium-molybdenum steels up to 650 °C, chrome-molybdenum-silicon alloyed steels up to about 750 °C, stainless steels up to 900 °C and Nickel-based superalloys for higher working temperatures. The great expectations, regarding low density steels, concerns their possible exploitation as stainless steels substituted. So these steel grades should result extremely resistant to corrosion and hot oxidation. In detail, low density steels should rival with the austenitic stainless steels, which show excellent resistance to high temperatures [12]. EXPERIMENTAL PROCEDURE In this research, a Fe-15%Mn-9.5%Al-6.5%Ni-1%Cr0.43%C alloy was melted in 10 kg electric arc furnace, using the same raw materials of the steel production cycle. The chemical composition of the alloy has been checked through O.E.S. measurements. Selected samples of the as-cast alloy have been undergone to an annealing thermal treatment in an Argon-protected atmosphere for 1 hour at 1300 °C. The cooling path was executed through a water quench. The assessment of the microstructures was achieved via specimen preparation, chemical etching with Beraha’s tint etch or Nital 5% optical microscopy and SEM/EDS and SEM/EBSD analysis. The corrosion resistance in chloride-rich environments of this Fe-15%Mn-9.5%Al-6.5%Ni-1%Cr-0.43%C alloy have been evaluated, the details of the procedures are

La Metallurgia Italiana - n. 1 2018

exposed following. The corrosion resistance in chloride-rich environments has been established through the identification of a corrosion rate. The corrosion rate has been calculated in millimeters per year. The same conversion table used by Oguike et al.[13] was used, which defines the conversion factor in order to derive the iso-corrosion curves from quantities directly measurable. In particular, the needed data are easily acquirable as the change in weight, the time and the surface exposed to the corrosive environment: it has been therefore used the conversion proposed by Oguike et al.[13], which provides the corrosion rate in mm per year, normalized on the density of the material. The procedure used to collect the data is indicated in the standard ISO 17475. Since the influence of the microstructure plays a key role in the corrosion resistance, especially in chlorine-rich environments, first it has been studied the influence of the annealing thermal treatments at fixed concentration of HCl (2.3%) and temperature (ambient temperature). Further, the experimental plan (Fig. 1) has been realized using four HCl concentration and four temperatures, in order to derive an iso-corrosion curve which can be directly compared to those of stainless steels. For each point of the experimental plan, it was then calculated the corrosion rate. This experimental plan has been followed for specimens realized in as-cast material and in 1300 °C annealed material, in order to show the considerable possibilities of improvement, obtainable with an appropriate heat treatment, in the corrosion rate and to extrapolate the two iso-corrosion curves. Further, to give a more complete description of the corrosion resistance behavior of this alloy potentiodynamic polarization tests have been performed. The analysis have been executed on the as-cast alloy and the 1300 °C annealed material, which exhibits the best performance in terms of corrosion resistance. The potentiodynamic curves have been obtained by means of a “three-electrode” potentiostat working with a saturated calomel electrode (S.C.E.) in a solution of water and 35 g/l of NaCl. The reference standard for the execution of the tests is the ISO 17475, using specimens previously polished and left in contact with the atmosphere for 1 day, to allow the complete development of any passive film on their surface. RESULTS The tests has been performed on both the as-cast material and the 1300 °C annealed alloy, in order to characterize both the initial material and its best achieved performance. Thus, the iso-corrosion curves have been derived, to have a clear

Industry news understanding of the corrosion phenomenon in chlorine-rich environment as function of both temperature and HCl concentration. The results, obtained for the as-cast material, (Fig. 1) show

the trends of the corrosion rate as a function of the concentration of HCl at constant temperature, and then as a function of temperature at a constant HCl concentration.

Fig. 1 - Corrosion rates of the as-cast material as a function of temperature (left) and HCl concentration (right) In detail, there are several aspects, which can be underlined. All trends regarding the variations of the HCl concentration are with good approximation parabolic with a downward concavity. Thus, it is possible to observe that the effect of HCl concentration gradually becomes less important, tending to a saturation value. On the other hand, the trends of the rate of corrosion are rather linear as function of temperature. Consequently, it can be highlighted a stronger influence of the temperature on the corrosion rates especially for the larger HCl concentrations.

It is therefore possible to conclude that above a concentration of HCl of about 2% the corrosive phenomenon is mainly driven by the temperature rather than by the concentration of HCl. In order to be able to observe in a single graph the combined effects of temperature and concentration of HCl, the regression of the entire experimental plan has been calculated, and represented in the three axes diagram (Fig. 2), where several iso-corrosion lines have been highlighted.

Fig. 2 - Corrosion rates of the as-cast material as a function of temperature and HCl concentration

La Metallurgia Italiana - n. 1 2018

Attualità industriale Then, the corrosion resistance tests in chlorine-rich environments of the 1300 °C annealed alloy provide the following results. Also for this thermal treated material, the charts show respectively the

corrosion rate as a function of the concentration of HCl at constant temperature, and then as a function of temperature at constant HCl concentrations (Fig. 3).

Fig. 3 - Corrosion rates of the 1300° C annealed alloy as a function of temperature (left) and HCl concentration (right) Observing these results, the strong decrease in the corrosion resistance, whenever the temperature exceeds the 60 °C, appears evident. In detail, all the measured values are quite low, featuring this thermally treated alloy by high corrosion resistance properties. Further, all the trends, occurring at the variation of the HCl concentrations and temperatures, appear linear. However, the tests performed at 80 °C show a much lower corrosion resistance with respect to the other collected data. This implies the overcoming of a temperature threshold

value, the critical pitting temperature, featuring the activepassive corrosion behavior of this 1300 °C annealed alloy. This feature appears more pronounced in the thermally treated alloy than in the as-cast material. As for the previous case, in order to be able to observe the combined effects of temperature and HCl concentration, the regression of the entire experimental plan has to be calculated, and represented in the three axes diagram (Fig. 4), where several iso-corrosion lines have been highlighted.

Fig. 4 - Corrosion rates of the 1300° C annealed alloy as a function of temperature and HCl concentration

La Metallurgia Italiana - n. 1 2018

Industry news Moreover, potentiodynamic polarization tests have been performed and the results are reported following (Fig. 5). These tests compared AISI 304 austenitic stainless steel, the as-cast low density alloy and the 1300 °C annealed mate-

rial. The initial potential value was considered as the difference between the potential value of the electrode and the reference electrode potential, equal to 0.244 V.

Fig. 5 - Anodic polarization curves (in 3.5 wt.% of NaCl) of: AISI304 (as reference), Fe-Al in the as-cast condition and annealed (1300 °C for 1 hour) DISCUSSION The results obtained for the corrosion resistance in chlorinerich environments appears in line with the experiments executed on the Fe-Al-Mn-C steel grade[14]. It was observed that the ferritic phase is always subject in a preferential corrosion in the case of a corrosive environment with the presence of chlorides. This occurs regardless the Chromium addition (as b.c.c.-stabilizer element). The reason of this behavior is related to the poor resistance to chlorine-rich environments of Aluminum. On the other hand, the trend in the corrosion resistance displayed by the annealed alloy appears in contrast to the previous consideration. Indeed, the specimen, that shows a higher corrosion resistance, is the sample thermally treated at 1300 °C, which microstructure is composed by the highest content of b.c.c. β’ phase. This behavior is explained by the higher chemical and microstructural homogeneities of the 1300 °C annealed alloy with respect to the other specimens. Although the corrosion resistance at high temperatures (60-80 °C) is quite low, at low temperatures (40 °C) and in particular at room temperature the corrosion resistance is 42

excellent and competitive against those of stainless steels, showing values similar to AISI 304. In accordance to this the slope of the iso-corrosion curve remains low with increasing concentration of HCl, demonstrating the great influence of the temperature. In general, the effect of the solubilization treatment at 1300 °C can be summed up as a considerable increase of the corrosion resistance compared to the as-cast material, in particular at low temperatures (20-40 °C). At high temperatures (60-80 °C) the benefit, although present, is less evident, reflecting the great influence of temperature in the corrosion process. At low temperatures (20-40 °C), in particular, the resistance is excellent, being able to compete with the best stainless steels already available on the market. As the temperature increases, however, the corrosion resistance worsens systematically, then coming equal to the values featuring the austenitic stainless steel AISI 316. The progressive and rapid deterioration of the corrosion resistance is detectable around 60 °C: it is therefore probable that this temperature value coincides with the breaking of a passive film and, then, with the critical pitting temperature. Consequently, it is possible La Metallurgia Italiana - n. 1 2018

Attualità industriale to divide the iso-corrosion curve of the 1300 °C annealed alloy into two areas. In the first, indicatively below 60 °C, the corrosion rate is primarily driven by the temperature displaying a very low slope. On the other hand, above 60 °C the slope increase drastically and the HCl concentration results to be the rate governing factor. This steel, therefore, ensures exceptional performance in the corrosion resistance, resulting as a potential rival for commercial stainless steels and as a promising material of choice in several industries. The better corrosion response of the annealed condition is probably due to the microstructure, in particular the absence of a big amount of secondary phase enhances the corrosion resistance respect to the as-cast one. The stainless steel has a better corrosion resistance but the different in the corrosion current is restrained. However, further investigation will be needed for a better comprehension of the corrosion resistance, the presence of a protective layer and the pitting

mechanism[15]. CONCLUSIONS In this work, the corrosion resistance and the hot oxidation resistance of the Fe-15%Mn-9.5%Al-6.5%Ni-1%Cr-0.43%C alloy, produced in laboratory tests as low-density steel grade, has been characterized. The corrosion resistance in chlorine-rich environments was evaluated both through iso-corrosion curves, derived by gravimetric measurements, and through potentiodynamic polarization analysis. In detail, it has been verified that the austenitic phase is more resistant than the b.c.c. β' matrix, due to the lower presence of aluminum, which has a poor resistance to chlorides. The iso-corrosion curves show good corrosion resistance for the as-cast material, and excellent resistance of the 1300 °C annealed material, especially at low temperatures but also good at high temperatures.

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

G. FROMMEYER and U. BRÜX, Steel Res. Int. 77-9, (2006), p. 627. L. ZHANG, R. SONG, C. ZHAO, F. YANG, Y. XU and S. PENG, Mater. Sci. Eng. A 643, (2015), p. 183. Z.Q. WU, H. DING, X.H. AN, D. HAN and X.Z. LIAO, Mater. Sci. Eng. A 639, (2015), p. 187. S. LEE, J. JEONG and Y.-K. LEE, J. Alloys Compd. 648, (2015), p. 149. F. YANG, R. SONG, Y. LI, T. SUN and K. WANG, Mater. Des.76, (2015), p. 32. S.-H. KIM, H. KIM and N.J. KIM, Nature 518, (2015), p. 77. I. OHNUMA et al, Thermodynamic Database High-Strenght Low-Density Steels. Tohoku University, Department of Materials Science (2010). H. KIM, D.-W. SUH and N.J.KIM, Sci. Technol. Adv. Mater. 14-1, (2013), 014205. J.W. LEE and T.F. LIU, Mater. Chem. Phys. 69-1-3, (2001), p. 192. K. CHOI, C.-H. SEO, H. LEE, S.K. KIM, J.H. KWAK, K.G. CHIN et al. Scr. Mater. 63-10, (2010), p. 1028. U.H. KIVISÄKK, Corrosion 61-6, (2005), p. 602. C.J. WANG, J.W. LEE and T.H. TWU, Surf Coatings Technol. 163-164, (2003), p. 37. R.S. OGUIKE, Adv. Mater. Phys. Chem. 4-8, (2014), p. 153. M. RUSCAK and T.-P. PERNG, J. Mar. Sci. Technol. 1-1, (1993), p. 1 V.F.C. LINS, M.A. FREITAS and E.M. PAULA E SILVA, Appl. Surf. Sci. 250-1-4, (2005), p. 124.

La Metallurgia Italiana - n. 1 2018

Industry news Caratterizzazione microstrutturale ed elettrica di Cu1.2Mn1.8O4 applicato come rivestimento protettivo in pile a combustibile ad ossidi solidi (SOFC) a cura di: R.Spotorno, P.Piccardo, S. Barison, S. Fasolin In questo lavoro, un rivestimento protettivo di Cu1.2Mn1.8O4 è stato depositato su un acciaio inossidabile ferritico AISI 441 mediante magnetron sputtering ottenendo uno strato denso e adeso al substrato. I campioni sono stati invecchiati in forno per 250 ore a 800°C in flusso di aria applicando una corrente continua di 0.5 A∙cm-2 con lo scopo di valutarne il comportamento alle condizioni di esercizio nel comparto catodico di una pila a combustibile ad ossidi solidi (SOFC). Il composto ha dimostrato un’ottima stabilità chimica, compatibilità meccanica con l’acciaio ed efficacia nel limitare la diffusione del cromo proveniente dal substrato. Inoltre, durante l’invecchiamento e il raffreddamento è stata misurata la resistenza elettrica specifica (ASR) del campione in modo da valutarne l’evoluzione nel tempo e l’effetto della temperatura. La conducibilità elettrica del campione è risultata termicamente attivata, e direttamente proporzionale alla temperatura. I valori di ASR si sono stabilizzati durante l’invecchiamento a 0.23 Ω∙cm2.

PAROLE CHIAVE: RIVESTIMENTI PROTETTIVI - OSSIDAZIONE A CALDO - SPINELLI - CU-MN - SOFC

R.Spotorno, P.Piccardo Università degli Studi di Genova S. Barison, S. Fasolin Istituto ICMATE - Consiglio Nazionale delle Ricerche

INTRODUZIONE Le celle a combustibile sono dispositivi in grado di trasformare elettrochimicamente l’energia chimica contenuta nell’idrogeno o idrocarburi in energia elettrica senza passaggi intermedi. Per questa caratteristica esse presentano un’elevata efficienza energetica che, associata alla modularità, flessibilità nell’utilizzo del combustibile e assenza di rumori, vibrazioni ed emissioni nocive, le rendono dispositivi

interessanti per sopperire ai fabbisogni energetici sempre in aumento con un ridotto impatto ambientale. Le SOFCs sono una classe di celle a combustibile che utilizzano un elettrolita solido, il quale permette la conduzione di ioni di ossigeno dal catodo all’anodo. La loro temperatura di esercizio è compresa tra 600°C e 1000°C, garantisce un buon rendimento e rende superfluo l’uso di catalizzatori [1]. La densità di potenza erogata da una singola cella (circa 1

La Metallurgia Italiana - n. 1 2018

Attualità industriale W/cm2) ne limita le applicazioni dirette rendendo necessaria la costruzione di pile, sistemi formati da più celle connesse che consentono di raggiungere tensioni e correnti elettriche adeguate per qualsiasi applicazione. L’assemblaggio di più celle prevede l’introduzione di componenti aggiuntivi tra cui gli interconnettori, che permettono di collegare elettricamente più celle mantenendone separati i flussi dei gas reagenti. Per il corretto funzionamento della pila, l’interconnettore deve possedere diverse caratteristiche tra cui: (i) elevata conducibilità elettrica; (ii) resistenza all’ossidazione a caldo; (iii)impermeabilità ai gas; (iv)stabilità meccanica; (v)compatibilità chimica e meccanica con gli altri componenti della pila; (vi) elevata conducibilità termica; (vii) bassi costi di produzione e lavorazione [2,3]. Gli acciai inossidabili ferritici sono i materiali più utilizzati come interconnettori per SOFC date le loro caratteristiche meccaniche, coefficiente di espansione termica simile ai materiali costituenti le celle e la formazione di ossidi protettivi a base cromo che li rendono resistenti alle alte temperature mantenendone elevata la conducibilità elettrica [3]. Tuttavia, alle condizioni operative delle SOFC, per la presenza di ossigeno e vapore d’acqua al comparto catodico, il cromo tende a reagire formando specie volatili che impoveriscono lo strato di ossido superficiale e si depositano sulla superficie catodica compromettendone l’attività catalitica [4]. La contromisura a questo fenomeno consiste nell’applicare dei rivestimenti protettivi sull’acciaio che, oltre a mantenere adeguate caratteristiche chimiche e fisiche dell’interfaccia interconnettore/elettrodo, limitano la diffusione di ossigeno e acqua verso il substrato e la migrazione di cromo verso l’elettrodo prevenendone l’avvelenamento. Tra i diversi tipi di rivestimenti recentemente studiati, gli spinelli conduttivi sono spesso applicati per le loro adeguate proprietà elettriche, capacità di absorbire varie specie di cromo, stabilità e compatibilità con i substrati ed elettrodi [5,6]. In questo lavoro uno spinello a base Cu-Mn è stato applicato mediante la tecnica di magnetron sputtering su un acciaio inossidabile ferritico AISI 441. L’evoluzione delle proprietà elettriche e microstrutturali è stata caratterizzata alle condizioni operative per 250 ore.

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MATERIALI E METODI Uno spinello di rame e manganese con un rapporto stechiometrico rispettivamente di 1.2-1.8 è stato applicato su substrati di AISI 441 (composizione in tabella 1) [7] mediante magnetron sputtering. La scelta del rapporto stechiometrico è stata dettata dalla stabilità del composto [8], e dalla convenienza nel trovarsi in eccesso di manganese, nel caso in cui composizioni come il Cu1.5Mn1.5O4 dovessero risultare privilegiate, dal momento che l’Mn 3O4 presenta migliori proprietà elettriche e meccaniche rispetto al CuO. Il metodo di deposizione ha previsto l’impiego di due target a base rame e manganese come sorgenti dei rispettivi elementi per la deposizione sui substrati di acciaio. Per la procedura è stato scelto di applicare prima lo strato di manganese con uno spessore di 1.5µm e in seguito 0.97µm di rame per ottenere la composizione Cu 1.2Mn1.8O4. A seguito della deposizione dei due strati metallici i campioni sono stati trattati per 2 ore a 800°C in aria statica per garantire l’interdiffusione degli elementi nei due strati e la loro ossidazione per la formazione della fase desiderata. Un campione è stato collegato mediante un circuito a 4 fili a un galvanostato/potenziostato (Solartron SI 1287) per l’effettuazione di misure elettriche durante il test. Il contatto elettrico sul campione è stato realizzato mediante reti di platino. Il campione è stato testato per 250 ore a 800°C in flusso di aria (50 cc∙min-1) applicando una corrente continua di 0.5 A∙cm -2 con lo scopo di valutarne il comportamento alle condizioni di esercizio nel comparto catodico di una SOFC. Le misure elettriche sono state inoltre effettuate durante il raffreddamento del campione al termine dell’esperimento per valutare la resistenza elettrica in funzione della temperatura. La qualità dei rivestimenti dal punto di vista microstrutturale, adesione al substrato e effetto barriera alla diffusione del cromo è stata valutata a seguito di osservazioni al microscopio elettronico a scansione (SEM EVO 40, Zeiss equipaggiato con analizzatore EDX PentaFET) sulle sezioni trasversali dei campioni preparate dopo essere inglobati a freddo in resina epossidica, lucidati con carte abrasive e panni insieme a sospensioni di alcol e pasta diamantata con granulometrie decrescenti fino a 0.25 µm.

Industry news Tab. 1 - Composizione chimica dell'AISI 441 / AISI 441 chemical composition.

Designazione

Fe wt%

Cr wt%

C wt%

Mn wt%

Si wt%

Nb wt%

Ti wt%

441ss-std

Bal.

0.009

0.35

0.34

0.5

0.22

RISULTATI E DISCUSSIONE La sezione trasversale del rivestimento applicato mediante magnetron sputtering ha evidenziato la presenza di due strati densi e ben distinti corrispondenti ai materiali depositati in sequenza (Figura 1a). L’analisi composizionale ha riscontrato la presenza di ferro e cromo in percentuali minori del 2 e 0.5% rispettivamente (Spettri 1 e 2, Tabella 2). A seguito del trattamento termico di ossidazione degli strati depositati, la sezione trasversale del campione ha mostrato uno strato di ossido omogeneo, adeso al substrato metallico e di elevata densità (Figura 1b). Alcune porosità chiuse e di dimensioni ridotte sono state riscontrate nella parte più esterna del rivestimento, lontane dall’interfaccia con il metallo. L’analisi composizionale (Spettro 3, Tabella 2) ha confermato la presenza di una singola fase a base Cu-Mn con minori contenuti di ferro, e cromo. Del primo è stato riscontrato un aumento rispetto al campione non ossidato evidenziandone la tendenza a diffondere dall’acciaio per effetto della temperatura. Dopo l’invecchiamento a 800°C per 250 ore il rivestimento è risultato ancora ben adeso al substrato ma caratterizzato da una maggior porosità nella sezione, vicina all’interfaccia metallo/ossido (Figura 1c). La composizione del rivestimento è rimasta stabile mantenendo costante il rapporto tra rame e manganese lungo la sezione come evidenziato dalle analisi EDX (Figura 1d). La presenza di ferro è risultata omogenea lungo la sezione del rivestimento mentre il cromo ha presentato un picco all’interfaccia tra i due materiali e un andamento decrescente nel rivestimento allontanandosi dal substrato. Il picco relativo al cromo evidenzia la formazione di uno strato di ossido sulla superficie dell’acciaio per effetto dell’ossidazione a caldo [3]. Tale ossido, già visibile nel campione ossidato (Figura 1b), corrisponde allo strato nero all’interfaccia tra i due materiali e risulta stabile e compatto sulla superficie dell’acciaio. L’andamento del profilo del cromo nel rivestimento ne indica l’efficacia nel limitarne la diffusione [9]. Il coefficiente di diffusione del 46

cromo nel Cu 1.2Mn 1.8O 4 è stato stimato risolvendo la seconda legge di Fick mediante interpolazione della funzione errore complementare del profilo EDX [10], risultando di 1.53∙10 -14 cm 2·s -1, un ordine di grandezza superiore rispetto a dati in letteratura [9,11]. L’ASR (Figura 2a) ha presentato un aumento molto rapido nelle prime ore di misura raggiungendo un picco di 1.37 Ω∙cm 2 dopo due ore a 800°C per poi diminuire e stabilizzarsi a 0.26 Ω∙cm 2 dopo 17 ore. L’andamento dell’ASR nelle successive ore di misura ha presentato oscillazioni attorno a tale valore fino alla fine dell’esperimento a 250 ore con un valore dì 0.23 Ω∙cm 2. L’ ASR misurata è risultata eccessiva per l’applicazione in SOFC, che prevede un massimo di 0.1 Ω∙cm 2 [3], tuttavia l’ottimizzazione degli spessori depositati può portare alla diminuzione della resistenza elettrica al di sotto del valore limite. Il grafico di Arrhenius (Figura 2b) evidenzia l’andamento decrescente dell’ASR all’aumentare della temperatura indicando una mobilità delle cariche elettriche termicamente attivata con energie di attivazione superiori rispetto a quelle del Cu 1.2Mn 1.8O 4 [12,13], probabilmente a causa dell’elevato contributo dello strato di ossido ricco in cromo formatosi sul substrato [14]. Un ulteriore effetto della presenza di ossidi di Cr all’interfaccia tra il substrato e il rivestimento può essere la variazione della pendenza dell’ASR nel grafico di Arrhenius al di sopra dei 600°C, associata a un aumento nell’energia di attivazione da 0.52 a 0.67eV [11,14].

La Metallurgia Italiana - n. 1 2018

AttualitĂ industriale

Fig. 1 - Micrografia SEM-BSE su sezioni trasversali dei campioni a) dopo deposizione; b) ossidato; c) invecchiato 250 ore; d) analisi EDX effettuata lungo la freccia indicata in c) / SEM-BSE images of cross sections of samples a) as deposited; b) oxidized; c) aged for 250 hours; d) EDX analysis performed along the arrow indicated in c).

Tab. 2 - Analisi EDX nei punti indicati in Figura 1; i risultati sono espressi in percentuale atomica / EDX analysis performed in the positions indicated in Figure 1; results are in atomic percentage.

Fe Cu Mn

Spettro

Spettro 1

4.83

0.48

1.17

1.23

92.28

Spettro 2

7.65

0.67

71.76

1.94

17.99

Spettro 3

53.91

0.28

25.14

5.65

15.02

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Industry news

Fig. 2 - a) Evoluzione dell’ASR del rivestimento durante l’invecchiamento di 250 ore; b) Dipendenza dalla temperatura dell’ASR del rivestimento / a) Evolution of the coating ASR during aging for 250 hours at 800°C; b) Temperature dependence of the coating ASR.

CONCLUSIONI L’applicazione di Cu1.2Mn1.8O4 mediante magnetron sputtering su acciaio inossidabile ferritico AISI 441 ha portato all’ottenimento di un rivestimento monofasico, denso e adeso al substrato anche a seguito di trattamenti termici di invecchiamento in aria a 800°C per 250 ore dimostrando un’ottima stabilità chimica e compatibilità meccanica con l’acciaio. Tale composto è inoltre risultato efficace nel limitare la diffusione di cromo dalla superficie dell’acciaio. Per quanto riguarda le proprietà elettriche, l’acciaio rivestito con il Cu1.2Mn1.8O4 ha

mostrato un meccanismo di conducibilità termicamente attivato con dei valori di ASR stabilizzati a 0.23 Ω∙cm2 durante il test di 250 ore a 800°C. L’andamento della curva di ASR ha confermato la stabilità del composto applicato come rivestimento. Tuttavia per l’applicazione in SOFC sono necessari valori inferiori a 0.1 Ω∙cm2, probabilmente ottenibili mediante un’ottimizzazione degli spessori applicati.

RINGRAZIAMENTI Il lavoro presentato è stato svolto nell’ambito di un assegno di ricerca finanziato dall’Associazione Italiana di Metallurgia.

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Attualità industriale BIBLIOGRAFIA [1] J. LARMINIE, A. DICKS, Fuel Cell Systems Explained, 2nd Edition, Wiley (2003). [2] K.C. WINCEWICZ, J.S. COOPER, Journal of Power Sources, nr. 140 (2005), p. 280. [3] J. WU, X. LIU, Journal of Material Science & Technology, nr. 26 (2010), p. 293. [4] S.P. JIANG, Y.D. ZHEN, Solid State Ionics, nr. 179 (2008), p. 1459. [5] N. SHAIGAN, W. QU, D. G. IVEY, W. CHEN, Journal of Power Sources, nr. 195 (2010), p. 1529. [6] R. SPOTORNO, P. PICCARDO, F. PERROZZI, S. VALENTE, M. VIVIANI, A. ANSAR, Fuel Cells nr. 15 (2015) p. 728. [7] P. D. JABLONSKI, C. J. COWEN, J. S. SEARS, Journal of Power Sources, nr. 195 (2010), p. 813. [8] A.-M. AZAD, A. HEDAYATI, M. RYDÉN, H. LEION, T. MATTISSON, Energy Technology, nr. 1 (2013), p. 59. [9] M. GALBO, K. J. YOON, U. B. PAL, S. GOPALAN, S. N. BASIL, Energy Technology 2015: Carbon Dioxide Management and Other Technologies, The Minerals, Metals & Materials Society, (2015). [10] K. RIGHTER, A. J. CAMPBELL, M. HUMAYUN, Geochimica et Cosmochimica Acta, nr. 69 (2005), p. 3145. [11] W. Huang, S. Gopalan, U. B. Pal, S. N. Basu, Journal of Electrochemical Society, nr. 155 (2008) p.B11617. [12] N. Hosseini, F. Karimzadeh, M.H. Abbasi, G.M. Choi, Ceramics International nr. 40 (2014) p. 12219. [13] B. E. Martin, A. Petric, Journal of Physics and Chemistry of Solids, nr. 68 (2007) p. 2262. [14] K. Huang, P. Hou, and J. B. Goodenough, Solid State Ionics, nr. 129 (2000), p.237.

Microstructural and electrical characterization of Cu1.2Mn1.8O4 coating for application in Solid Oxide Fuel Cells (SOFC) stacks KEYWORDS: COATINGS - HOT CORROSION - SPINELS - CU-MN - SOFC

In this work, Cu1.2Mn1.8O4 was applied by means of magnetron sputtering on AISI 441 ferritic stainless steel as protective coating. Samples were aged for 250 hours at 800°C in flowing air applying a DC electric load of 0.5 A∙cm-2 to reproduce the operating conditions of the cathode compartment of a Solid Oxide Fuel Cell (SOFC) stack. The coating resulted mechanically compatible with the steel and chemically stable exhibiting also good efficiency in limiting the chromium diffusion from the substrate. Additionally, the Area Specific Resistance (ASR) of the sample was measured during aging and cooling in order to evaluate its evolution over time and the temperature effects. The electrical conductivity resulted thermally activated, proportionally to the temperature. The ASR was stabilizing during aging to a value of 0.23 Ω∙cm2 after 250 hours.

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Industry news Comparative investigation of deep drawing formability in austenitic (AISI 321) and in ferritic (DIN 1.4509) stainless steel sheets a cura di: C. de Paula Camargo Pisano, H. J. B. Alves, T. Reis de Oliveira, C. G. Schön One of the most challenging activities of the today’s industry is to balance productivity and quality of the production. Across the years there was an evolution on the technologies of production, of raw materials fabrication and on the qualification of the working force. In some circuits, however, there still remain gaps on the knowledge of various stages of the production chain, which need to be solved in order to the improve product quality. One of these gaps is the knowledge of the basic properties of raw material that is being used on the stamping production. The present work discusses and compares formability, as evaluated by the forming limit curve, of two different stainless steel 2mm sheets: an austenitic class (AISI 321) and a ferritic class (DIN 1.4509). The main mechanical properties and a differentiation between the grades on the deep drawing ability are presented by the results of some well known tests, such as the Tensile, Erichsen, Swift and Nakazima. The differences in formability between the two materials are discussed in terms of physical metallurgical reasoning, microstructure, and texture evolution, and allows to understand the characteristics and limitations of formability, concerning deep drawing of these steel grades.

KEYWORDS: STAINLESS STEEL – DEEP DRAWING – FORMABILITY – FORMING LIMIT DIAGRAMS – MECHANICAL BEHAVIOUR

Caio de Paula Camargo Pisano Aperam South America and Universidade de Sao Paulo, Brazil Hélio José Batista Alves, Tarcísio Reis de Oliveira Aperam South America Cláudio Geraldo Schön Universidade de Sao Paulo, Brazil INTRODUCTION According to a study that was made in 2002 by the Federal Highway Administration (FHWA), from United States, entitled as "Corrosion Costs and Preventive Strategies in United States", the annual costs due to corrosion problems represented around US$ 276 billion, which represented around 3,1% of the GDP of that country [1]. Considering these results, there is an urgent need to use alternative materials, such as stainless 50

steel, that can resist to the corrosive actions from the environment. Whenever it is necessary to discuss about the application of stainless steel the first property that is taken into consideration is the corrosion resistance, however in an industrial process other properties are desired to achieve a better result, such as better mechanical properties or an ease of forming or other operations. Considering this innovative scope, it is possible to mention some examples of the usage La Metallurgia Italiana - n. 1 2018

AttualitĂ industriale of stainless steel in industry in which corrosion resistance is not the single desired property, : structural components for buildings, vehicles, boats and many other cases, where structural integrity is required , in refractory coatings in furnaces, in which high resistance at high temperature is necessary for the good performance of the material, in surgery equipments [2], for which it is necessary that the material avoid the growth of noxious bacterias [3], in the exhaust system of vehicles in which the material needs to combine the corrosion resistance with the capacity to keep its mechanical properties in high temperatures [4] and many other applications. The application of stainless steel in exhaust systems, which has already been previously quoted , was a great evolution to increase the efficiency of automobiles and also reduced the gas emissions on the environment, which also increased the sustainability of the vehicles that are being produced [5]. All these applications were only possible due to the development and usage of the many kinds of production processes, where it is possible to highlight the deep drawing process, since the evolution of the market led to an increase on the complexity

of the parts and therefore the deep drawing process had to adapt its parameters and practices to maintain its utility. However the deep drawing process also depends on the material performance, therefore the steel mills had to develop their processes in order to achieve better mechanical, metallurgical and microstructural properties, according to a Management and Strategic Content Center from the Brazilian Government study, it is estimated that since the world crisis of 2008, 75% of the steel grades that were being used in the industry, have been newly developed [6]. The mechanical tests can be defined as a simplified procedure, wherein which the performance in service of the material will be evaluated, in a static or dynamic condition, therefore the portability of these properties is crucial on this procedure. The tensile test is one of these simplified tests, however its results, such as Yield Strength, Tensile Strength, Elongation and Stress x Strain Diagram others, in some cases it s not possible to evaluate the real mechanical behavior of the materials, leading to unsatisfactory results on the dimensioning of parts or tools.

Fig. 1 - Comparative Stress vs Strain Diagram considering an austenitic and a ferritic stainless steel.

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Industry news Considering this, a new concept to evaluate the formability of materials have been developed, this concept is named as Forming Limit Diagrams (FLD’s) and was introduced by Keeler [7], who worked on the stretching area of the diagram and Goodwin [8], who subsequently developed the deep drawing area of the diagram. The FLD’s can provide a broader idea of the performance of the material in a deep drawing process [9], and also can provide to engineering projects a more precise result for the cor-

rect dimensioning of raw material, parts and tooling. Besides the fact that the FLD’s can provide a correct safety margin of usage to the materials, they can also indicate the critical areas for fracture of a specific part. With this kind of information it is possible to direct the efforts of an analysis and work on the process correctly, avoiding future failures or embrittlement of components due to the excessive reduction of thickness [10].

Fig. 2 - Theoretical Forming Limit Diagram, indicating the safe and failure zones, the stretching area, deep drawing area and the forming limit.

Considering this background, the steel mills must keep a close eye on the performance of the materials on these many processes, always using the correct information, and by doing that they will be able to understand and evaluate the market necessity and, as mentioned above, direct the production efforts to have better grades. This work will approach these analysis for two different types of stainless steel, an austenitic grade (AISI 321) and a ferritic grade (DIN 1.4509), using the FLD methodology, tensile test results and simulative stamping tests, such as Erichsen and Swift. By correlating these results it will be possible to discuss the best process practices for each material and also discuss about the particularities of these grades on the deep drawing process.

EXPERIMENTAL PROCEDURE In order to obtain the FLD’s, mechanical properties and simulative stamping tests results, such as Erichsen and Swift, samples of both grades, AISI 321 and DIN 1.4509, were collected in a cold rolled, annealed and pickled condition. It is very important to reinforce that the thickness used in this present work is a typical thickness used by the automobile industry, which is 2 mm. The material thickness has a direct influence on the results of the FLD [11], therefore different thicknesses for the same grades that are being presented on this work, may present different results and performances.

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Attualità industriale

Fig. 3 - Thickness influence on the results of the FLD [Adapted from 11]

Chemical composition results were obtained with two methods, metallic elements measurement was made with X-Ray fluorescence spectrometry and non-metallic elements measurement was made with infrared absorption after fusion. These results are important since they can support on the understanding of some mechanical and metallurgical behaviors, such as hardness and martensite transformation after cold deformation. The mechanical properties, such as yield strength, tensile strength, uniform elongation, total elongation, hardness coefficient, anisotropy and stress vs strain diagram were obtained by a tensile test, using AVE technology (Advanced Video Extensometer) as a data collector. The strain is measured with the support of a high resolution camera which tracks two contrast marks on the specimen, on the longitudinal and transversal directions. The specimens were collected in three different directions of the sheet: 0°, 45° and 90°, in function of the rolling direction. Fig. 4 - Tensile test with AVE technology as data collector

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Industry news The Erichsen test was performed in an Erichsen Press (model 142-40) considering ASTM E643-15 [12] guidelines, where it is determined that 90 mm is the minimum dimension of the specimen, so the material doesnâ&#x20AC;&#x2122;t flow into the tooling affecting the results, also it is important to keep the blank holder with a high pressure and lubricate the punch. Considering these directives, 2,00 mm x 100 mm x 100 mm specimens were collected for each grade and the test was performed with two lubricant conditions, with only lubricant and with lubricant and PVC coating.

Fig. 5 - Tooling with lubricant and PVC coating before Erichsen test. The Swift test is a cup stamping test and simulates the deep drawing condition of the materials determining the Limit Drawing Ratio (LDR), which is a relation between the maximum

blank diameter without rupture on the test and the punch diameter.

(1)

The test was performed in the same press of the Erichsen Test (model 142-40) and the tooling used for it already cuts the circular blanks before positioning and stamping them. The punch had a diameter of 33 mm and the set of blanks had a minimum diameter of 63 mm with an increment to other blanks diameters of 1 mm, until reaching 70 mm. The most common test to obtain FLDâ&#x20AC;&#x2122;s is the Nakazima test [13] which was the one used in this work, the test was performed considering all guidelines of ASTM E2218-15 [14], with the addition of some conditions to increase the test results. The standard determines that for a hemispherical punch with a diameter of 100 mm, which was the one used, the specimens must have their widths varying from 12 mm to 200 mm,

with an increase varying from 12 to 25 mm, and the length varying between 180 mm and 225 mm, therefore the specimens used for the Nakazima test had a fixed length of 200 mm and the widths varying from 40 mm to 200 mm, with an additional specimen of 200 mm x 200 mm with the addition of PVC coating. The ASTM standard also determines a few models and dimensions for the surface mashes, for the square mash, which was the one selected for this test, the size of the sides of the square need to have 2,5 mm or less, the size for the square mash of the specimens on this test had 2,0 mm and was marked with an electrolytic marking procedure.

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AttualitĂ industriale

Fig. 6 - Representation of the tooling used in the Nakazima test.

Fig. 7 - Representation of a specimen after Nakazima test, with the reference used by the software ASAME.

To keep the representativeness of the results it is necessary to keep the pressure high on the blank holder, in order to avoid the flow of the material, for this work a force of 100 kN was used, also the specimens must have a good edge condition, with a maximum burr of 10% of the thickness. The test must end as soon as the load drops, which is the spot where the necking begins and the measurement must be performed in the mash near the necking. To support with the measurement of the mesh on the specimens Nakazima test a software, named ASAME, was used and as mentioned above, the spots selected for these measurements were the ones near the necking area. To proceed with the measurement is necessary to place a cubic reference, which is supplied by the software and can be seen in figure 7, near the selected spots and take two photos, with an angular variation between them. After that these photos must be uploaded to the software that will convert the lines of the mash in digital elements and therefore will be possible to measure the distance between the intercepts and obtain Îľ1 and Îľ2 and consequently obtain the FLD.

RESULTS AND DISCUSSION Chemical Composition The chemical composition evaluation is an important initial parameter due to the influence of some elements on the properties of the materials. One of these properties is the hardness and it is known that as higher as the quantity of some chemical elements, higher the hardness will be.

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Industry news

Fig. 8 - Influence of the quantity of some chemical elements on the hardness of the material [15]. Other important parameter is the stability of the austenitic structure of AISI 321, the stability of this grade will determine the quantity of martensite formation after cold deformation. A concept to quantify the martensite formation is the calcu-

lation of a temperature called Md30/50 (°C), which is the temperature that 50% of martensite will be generated in the structure after 30% deformation [16], and the equation to calculate this temperature was proposed by Angel [17]:

Md30/50 (°C) = 413 – [462(%C+%N) + 9,2(%Si) + 8,1(%Mn) + 13,7 (%Cr) + 9,5(%Ni) + 18,5 (%Mo)]

(2)

Fig. 9 - Martensite formation in three different grades, after 35% of deformation and in the same temperature. a) UNS S30153, b) UNS S20153 and c) UNS S30403

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Attualità industriale The results of chemical composition, that can be seen in table 1, showed that both grades were according to ASTM A240/ A240M specifications. It was possible to notice that the car-

bon, nitrogen and phosphor levels are very low, but this is justifiable since both grades were produced through argon oxygen decarburization and vacuum oxygen decarburization.

Tab. 1 - Chemical composition results of DIN 1.4509 and AISI 321. Chemical Composition Results (%) Element

DIN 104509

AISI 321

0,01

0,0123

0,0118

0,4489

0,5039

0,2152

0,8243

0,0326

0,0293

0,0004

0,0009

17,5766

17,0093

0,2879

9,012

0,3921

0,0121

0,1353

0,1492

Using equation 3 to calculate Md30/50 (°C) for the AISI 321 it was possible to see that the temperature is around 72°C which indicates that this material has a low stability and after deformation will tend to form lots of martensite, this is an important information since the martensite formation in great quantities will affect the deep drawing performance and may lead to premature failures, mainly in processes with lots of stages.

Mechanical Properties The first mechanical property evaluated was not obtained with the tensile test, but through a hardness test. The hardness was obtained by the Rockwell methodology a test that was introduced in 1922 and measures the resistance of the material to the penetration of the tooling, the results are read on the equipment which leads to fast results, with small chances of human mistakes [15]. It was possible to see that the ferritic grade (DIN 1.4509) have a higher hardness when compared to the austenitic grade (AISI 321).

Tab. 2 - Hardness evaluation for both grades Average Hardness (HRb)

DIN 1.4509

AISI 321

Difference

Specimen 0°

77,3

71,4

5,9

Specimen 45°

77,4

69,7

7,7

Specimen 90°

77,4

71,4

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Industry news By only analyzing the chemical composition, these results would not be justifiable, however two different grades are being evaluated in this work, with different production processes and metallurgical structures, which justifies the difference on the results. These results can be seen on table 2. The typical mechanical properties were obtained with the tensile test, already described. It is possible to see in figure 10 that the austenitic stainless steel (AISI 321) have a higher capacity of absorbing energy due to a higher tenacity

(area below the diagram), when compared to the austenitic stainless steel (DIN 1.4509), this can be justified due to the atomic sliding plans with higher density of the FCC structure, which is the structure of the austenitic stainless steel, when compared to the BCC structure, which is the metallurgical structure of the ferritic stainless steel. Besides that, the austenitic stainless steel has a lower stacking-fault energy (SFE) when compared to the ferritic stainless steel, which is another indicative for the easiness of deformation of this grade.

Fig. 10 - Stress x Strain diagrams for both grades, AISI 321 and DIN 1.4509. The mechanical properties of the tensile test, as already mentioned, are a good initial perspective of the material behavior, however it shouldnâ&#x20AC;&#x2122;t be the only evaluated properties, since they are obtained in a unidirectional test, thatâ&#x20AC;&#x2122;s why these properties are obtained in three different directions, so complementary evaluations can be made, such as the verification of the Lankford Coefficient (R). This coefficient measures the capacity of the material to deform preferentially in the plane directions than on the thickness [18], in other words, as higher as this coefficient is most likely the material

will maintain its thickness during the deep drawing operation [19], which is a great capacity for the deep drawing process, mainly for the DIN 1.4509, that doesnâ&#x20AC;&#x2122;t have high elongation values, when compared to AISI 321. The reference for collecting the deformations to calculate the Lankford Coefficient (R) is 15% of strain on the tensile test [20] and its result is completely dependent of the rolling direction. The R value can be calculated according to equation 4:

(3)

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Attualità industriale After calculating R for each direction, it is possible to obtain an average value of these measurements called normal ani-

sotropy ( ), and it can be calculated according to equation 5:

(4) The anisotropy reference is an important indicator to predict the material performance on a process, since if this coefficient is lower than 1 the material will most likely tend to lose its thickness during stamping operations. And by evaluating the-

se results by the rolling direction it is possible to determine what is the best position to stamp a part and obtain the best result as possible, in a stamping operation.

Tab. 3 - Lankford Coefficients for DIN 1.4509 and AISI 321 Lankford Coeficient - R

DIN 1.4509

AISI 321

Specimen 0°

1,33

0,78

Specimen 45°

0,98

1,03

Specimen 90°

1,79

0,84

Normal Anisotropy

1,27

0,92

It is possible to verify that the ferritic stainless steel (DIN 1.4509) has a much higher influence of the rolling direction on the Lankford Coefficient, where 90° from it is the best direction to deep draw and 45° from the rolling direction is the worst direction to deep draw, with an anisotropy lower than 1. In a general analysis the results indicate that the ferritic stainless steel (DIN 1.4509) will most likely keep the thickness during a stamping operation when compared to the AISI 321, however this analysis shouldn’t be done without consider other parameters since the AISI 321 has specific metallurgical characteristics. As already mentioned, to enrich the comparison between these grades, two typical stamping tests were performed.

The Erichsen test simulates the stretching capability of the material and the advance of the punch, in millimeters, is measured until the machine load drops [21]. This distance is known as Erichsen Index and as higher as it is, it indicates that the material has a better stretching capacity. As already expected, the austenitic stainless steel AISI 321 had a higher Erichsen Index, when compared to the ferritic stainless steel DIN 1.4509, which is another indicative that the AISI 321 parameters processes, such as blank holder pressure, design of the tooling and advance speed of the punch, must be done considering this thickness reduction.

Fig. 11 - Erichsen Index evaluation between AISI 321 and DIN 1.4509.

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Industry news As a test without PVC coating was also made, it was possible to see that the extra lubrication increase the IE with an average value of 2,5%, however the effect on the load is more significant, since the addition of PVC coating reduced the machine load about 10% to obtain a better result, which means that for an industrial process the correct use of lubrication

can increase performance and reduce costs. The other performed test was the Swift Test, that measures the capacity of deep drawing of the materials on a cup stamping operation, as higher as the LDR on this test, better deep drawing capacity the material will have.

Tab. 4 - Swift Test results for AISI 321 and DIN 1.4509 DIN 1.4509

AISI 321

Punch Diameter

33 mm

Max Diameter w/o failure

69 mm

65 mm

Limit Drawing Ratio

2,09

1,97

As expected, the ferritic stainless steel DIN 1.4509 presented a higher LDR when compared to AISI 321, however there isnâ&#x20AC;&#x2122;t a big distance between the results indicating that both materials should have a similar performance in a deep drawing operation.

Finally the FLDâ&#x20AC;&#x2122;s were obtained to conclude the analysis, since its results are directed to a process behavior and provide a better understanding of the material performance than the other tests.

Fig. 10 - Stress x Strain diagrams for both grades, AISI 321 and DIN 1.4509. As expected, the ferritic stainless steel presented a better performance on the deep drawing area, which is the left side of the diagram and the austenitic stainless steel presented a better result on the stretching area, which is the right side 60

of the diagram. However both materials have a very similar performance, since the distance between the forming limit for them is not big.

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Attualità industriale CONCLUSIONS By the presented results it is possible to conclude that in an uniaxial condition, simulated by the tensile test, the austenitic stainless steel AISI 321 have a greater ease of deformation when compared to the ferritic stainless steel DIN 1.4509, mainly due to its metallurgical structure. However both materials had a very similar performance on the stamping tests despite their metallurgical and price differences. On the Erichsen test it was possible to conclude that a correct usage of lubrication can increase the performance and reduce costs in a stamping operation.

The FLD’s reinforced the similarity of the grades that was already being shown in the other tests, the main reason for this similarity is that the ferritic stainless steel was produced under process conditions that privileged the acquisition of better deep drawing results, this process conditions are mainly the slab production with electromagnetic stirrer, in order to break equiaxial grains, and a correct reduction plan from hot rolled coils to cold rolled coils to obtain as much as possible of the better texture, which for the ferritic stainless steels is the γ fiber, with planes {111} parallel to the normal direction [22].

REFERENCES [1] Koch GH, Brongers MPH, Thompson NG, et al. Corrosion Costs and Preventive Strategies in the United States (Report No. FHWA-RD-01-156). Dublin, Ohiohttps://www.nace.org/uploadedFiles/Publications/ccsupp.pdf (2002, accessed 29 September 2016). [2] Carbó HM. Aço Inoxidável – Aplicações e Especificações. 2001; 1–38. [3] Rebello M de C. O Aço Inoxidável Como Garantia da Segurança Alimentar Dos Produtos Minimamente Processadoshttp:// www.esalq.usp.br/departamentos/lpv/eventos/palestras/ProcessamentoMinimo2006.pdf (2006, accessed 29 September 2016). [4] Sato E, Tanoue T. Present and Future Trends of Materials for Automotive Exhaust Systemhttp://www.nssmc.com/en/tech/ report/nsc/pdf/6403.pdf (1995, accessed 29 September 2016). [5] Itoh I, Fukaya M, Hisatomi R, et al. Development of Ferritic Stainless Steel Foil as Metal Support for Automotive Catalytic Converterhttp://www.nssmc.com/en/tech/report/nsc/pdf/6412.pdf (1995, accessed 29 September 2016). [6] Bolt I de S. Tendências e Inovações em Aços. Estudo Prospectivo do Setor Siderúrgico. Brasíliahttp://www.abmbrasil.com. br/epss/arquivos/documentos/2011_4_19_8_46_4_33043.pdf (2008). [7] Keeler SP. Plastic instability and fracture in sheets stretched over rigid punches. 1961. [8] Goodwin GM. Application of Strain Analysis to Sheet Metal Forming Problems in the Press Shop. Epub ahead of print 1 February 1968. DOI: 10.4271/680093. [9] Chinouilh G, Toscan F, Santacreu PO, et al. Forming Limit Diagram Prediction of Stainless Steels Sheets. Epub ahead of print 16 April 2007. DOI: 10.4271/2007-01-0338. [10] Bastos AL. Análise do processo de estampagem de chapas de aço através da curva limite de conformação. Universidade Federal de Santa Catarina, Centro Tecnológico, Programa de Pós-Graduação em Ciência e Engenharia de Materiaishttp:// repositorio.ufsc.br/xmlui/handle/123456789/93161 (2009, accessed 29 September 2016). [11] Schaeffer L. Fundamentos do projeto de ferramentas para o processo de estampagem. Ferramental, 2006, pp. 31–36. [12] ASTM E643-15, Standard Test Method for Ball Punch Deformation of Metallic Sheet Material, ASTM International, West Conshohocken, PA, 2015, www.astm.org [13] SANAY B. Prediction of plastic instability and forming limits in sheet metal forming. MIDDLE EAST TECHNICAL UNIVERSITYhttps://etd.lib.metu.edu.tr/upload/12612486/index.pdf (2010). [14] ASTM E2218-15, Standard Test Method for Determining Forming Limit Curves, ASTM International, West Conshohocken, PA, 2015, www.astm.org [15] Bain EC, Paxton HW. Alloying elements in steel. American Society for Metals, 1961. [16] Padilha AF, Plaut RL, Rios PR. Annealing of Cold-worked Austenitic Stainless Steels. ISIJ Int 2003; 43: 135–143. [17] Angel T. Formation of martensite in austenitic stainless steel. J Iron Steel Inst 1954; 177: 165–174. [18] Charca Ramos G, Stout M, Bolmaro RE, et al. Study of a drawing-quality sheet steel. I: Stress/strain behaviors and Lankford coefficients by experiments and micromechanical simulations. Int J Solids Struct 2010; 47: 2285–2293. [19] Hosford WF, Caddell RM. Metal Forming: Mechanics and Metallurgy. 3rd ed. New York: Cambridge University Press, 2007. [20] Hosford WF. Mechanical Behavior of Materials. Cambridge: Cambridge University Press. Epub ahead of print 2009. DOI: 10.1017/CBO9780511810923. [21] Schaeffer L. Conformação de Chapas Metálicas. 1st ed. Porto Alegre, RS, 2004. [22] Costa ALN. Formação da Textura de Recristalização dos Aços Inoxidáveis Ferríticos AISI 430A e 430E. Rio de Janeiro: Instituto Militar de Engenharia, 2006.

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Industry news Strain induced martensite evolution in a rolling contact of SS AISI 304 a cura di: M. Werschler, P. Gümpel, K. Werner The evolution of strain induced martensite in austenitic stainless steel AISI 304 was investigated in a rolling contact on a two-discstribometer. The effects of surface roughness, slip and normal force as well as the number of load cycles were examined. In comparison to the investigations of martensitic phase transformation during cold rolling, the applied stresses are considerably lower. The formation of strain induced martensite was detected in-situ by means of a FERITSCOPE® MP30 and ex-situ by optical microscopy after etching with Kane etchant. Both number of load cycles and magnitude of normal force appeared to be the main influencing factors regarding strain induced martensitic evolution in low stress rolling contacts.

KEYWORDS: STRAIN INDUCED MARTENSITE – ROLLING CONTACT – SS AISI 304 – TRIBOMETER – MARTENSITIC EVOLUTION – AUSTENITIC STEEL

M. Werschler, P. Gümpel, K. Werner HTWG Konstanz, Germany

INTRODUCTION The behaviour of metastable austenitic steels under loading conditions such as tensile testing or low and high cycle fatigue (LCF/ HCF) has been widely studied. More complex conditions like rolling contacts have been the subject of less focus even though, these investigations are carried out on high levels of strain such as cold rolling [1]. In order that the evaluated hardening of austenitic steels due to the strain induced martensitic phase transformation can be helpful in modern production technologies [2], the phase transformation must be caused by much lower strain rates in comparison to the deformation by cold rolling. Therefore, the aim of this research project is to evaluate the evolution of martensite in the metastable austenitic stainless steel SS AISI 304 under relatively low stresses in a rolling contact and derive the influence of normal load (normal force), slip and surface roughness on the formation of martensite. Austenitic stainless steels feature a face-centred cubic (fcc) crystal structure. However, the austenitic phase is not sta-

ble, so that an applied mechanical load can cause a phase transformation to non-magnetic hexagonal close-packed (hcp) ε-martensite or ferro-magnetic body-centred cubic (bcc) α’-martensite. The austenitic phase’s stability can be described by means of the Md30-temperature, which stands for the temperature at which at least 50 % of the microstructure is transformed into martensite at a plastic deformation of 30 %. Whether the γ-austenite transforms to ε- or α’-martensite depends on the conditions at which the phase transformation takes place. The transformation can arise by one of the following transformation paths [3, 4]: a.) γ-austenite → ε-martensite b.) γ-austenite → α’-martensite c.) γ-austenite → ε-martensite → α’-martensite Low temperatures and small strain favour the formation of ε-martensite [5], higher strain the formation of α’-martensite, respectively. In the early stages of deformation, the strain

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Attualità industriale induced α’-martensite nucleates at the intersections of shear bands, twin boundaries, grain boundaries and grain boundary triple points [3, 4, 6]. Shear band is the global term for planar defects precipitated by plastic deformation [7]. With an increasing amount of deformation, the number of α’nuclei rises. The coalescence of α’-nuclei results in the formation of single martensite-laths [8]. The morphology of the martensitic phase depends on the grade of deformation. An increase in deformation first leads to a coalescence of single martensite-laths, after which the formation of blocky-shaped martensitic structures takes place in the particular grains and eventually leads to the evolution of a nanocrystalline microstructure with no clearly visible martensite-laths [9].

stress rolling contact a special two-discs-tribometer is used. This tribometer is designed to simulate the kinetics of a tooth flank contact and offers the possibility to apply different variations of complex rolling contacts. The main part of the tribometer is the containment, which houses the specimens (Fig. 1; a). The containment is grouped in three parts. An active, respectively movable part with specimen A, a passive, respectively static part with specimen B and the sensor cap (Fig. 1; b) with modular slots for a FERITSCOPE®, an infrared temperature sensor and a digital microscope. Each specimen is connected to a servomotor (Fig. 1; c and d) for rotational movements. The testing load is applied infinitely variable and in highly dynamic manner by a piezoelectric linear actuator (Fig. 1; e) [10].

EXPERIMENTAL DESIGN Tribometer In order to investigate the strain induced martensitic phase transformation of metastable austenitic steels in a low

Fig. 1 - Specialized two-discs-tribometer for the simulation testing of tooth flank contact. a) Containment with specimen; b) sensor cap with FERITSCOPE®, infrared temperature sensor and microscope; c) and d) rotational servo actuator of specimen A and B; e) linear piezoelectric servo actuator. A two-discs-tribometer is based on the tribological contact between the lateral surface of two axially symmetric and rotatable specimens (Fig. 2, a) [11]. Corresponding to the Hertzian theory of contact, different types of stress form in the tribological contact, depending on the normal force FN and the shape of the lateral surface of the specimen. A coordinated rotation of specimen nA and nB enables continuously variable relative La Metallurgia Italiana - n. 1 2018

slip s in the contact area [10]. The used tribometer additionally provides the function to vary the load parameters (FN, s) depending on the rotational angles φA/B during a single revolution (single turn profile) (Fig. 2, b). Furthermore, the test rig automation allows a fully automated variation of a single turn profile over a test (Fig. 2, c).

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Fig. 2 - a) basic principle of a two-discs-tribometer with differently shaped specimen; b) Example for varying load parameters during a single revolution FN = f(φ), s = f(φ); c) Variation of load parameter FN = f(φ) and s = f(φ) during a test after 1, 100 and 500 revolutions, respectively load cycles. Sensor system and measurands As a standard setting the tribometer is equipped with sensors to measure the normal force FN and the friction torque MR between the specimens. Furthermore, the rotational angles φA/B, rotational speeds nA/B, ambient temperature TU and position sp of the containments active part are recorded. Based on position sp the approach of the specimens ∆sp, respectively the deformation Σ can be calculated. Thus, Σ is relative to the absolute value sp at loaded state during the first load cycle. Each measurand can be displayed continuously for one revolution

(load cycle) or for a defined rotational angle φA/B over a complete test run [10]. Fig. 3 shows the geometrical arrangement of the additional sensors for the investigation of metastable austenitic steels in low stress rolling contacts. The sensors are an infrared temperature sensor to monitor the surface temperature of specimen B (Fig. 3, a), a FERITSCOPE® to capture the martensitic transformation (Fig. 3,b) and a digital microscope (200x magnification) for documentation of the specimen surface during a test run [12].

Fig. 3 - Sensor application on the two-discs-tribometer for the investigation of metastable austenitic steels in low stress rolling contacts. a) infrared temperature sensor; b) sensor head of FERITSCOPE® MP30 c) digital microscope [13].

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Attualità industriale Because of its special method of application, the FERITSCOPE® requires a detailed discussion. It was used with a defined space of 0.05 mm between the sensor head and the surface of the specimen (Fig. 3, b). The original purpose of the FERITSCOPE® (Helmut Fischer GmbH, FERITSCOPE® MP30) is the determination of the amount of ferrite (Fe%) in austenitic stainless steels and duplex-steels. Thereby a low-frequency magnetic field, generated by an excitation coil, interacts with the specimen’s volume and detects the difference in relative permeability of ferromagnetic and paramagnetic phases (ferrite and austenite or α’-martensite and austenite). The percentage of ferromagnetic phases within the included specimen volume depends on the induced voltage in a measurement coil [13]. Talonen et. al. [14] revealed that the amount of α’-martensite in austenitic stainless steels can be derived from the measured value in Fe% by multiplication with a factor of 1.7. Nevertheless, the measurement of the martensitic phase was deducted by using the raw reading, referred to as count rate, since the calculation algorithms, which convert the count rate into Fe%, are based on the assumption of a homogenous microstructure. Based on results of previous investigations [12] it is known that the resulting martensite shows clearly inhomogeneous distribution. In order to compensate differences in the specimens’ relative permeability the martensitic evolution is described by means

of the change of count rate (ΔX), which is calculated as the difference of the specimen’s count rate at current reading points and its count rate before the start of experiment. Furthermore Werner [15] pointed out that the measurement of the martensitic phase with the FERITSCOPE® on the two-discs-tribometer depends significantly on the specimen’s surface roughness, the distance between specimen and FERITSCOPE® as well as the distribution of the martensitic phase within the specimen’s microstructure. Specimen To investigate the martensitic transformation two types of specimen are utilized. Type one is a hard metal counterbody (KXF®) further labelled as HM-SPC (hard metal – specimen) (Fig. 4, a). Type two is the actual item of investigation and made of SS AISI 304 metastable austenitic stainless steel further labelled as SIM-SPC (strain induced martensite – specimen) (Fig. 4, b). The geometrical dimensions of the two specimens are given in Figure 4. Table Tab. 1 shows the exact composition of the used material of the SIM-SPC measured with spectral analysis. The calculated Md30 – temperature [16] without consideration of grain size is Md30 = -9°C.

Tab. 1 - Chemical composition of tested SS AISI 304

CHEMICAL COMPOSITION SS AISI 304 %C

% Si

% Mn

% Cr

% Ni

% Cu

% Mo

% Nb

0.018

0.089

0.555

1.04

18.21

7.88

0.332

0.405

0.032

All SIM-SPC are fabricated on a turning machine and passed through a solution annealing afterwards. The annealing process is executed in a continuous furnace with nitrogen atmosphere at a temperature of 1,100°C. The quenching is carried out with nitrogen, too. During a test run a SIM-SPC is charged with a constant load for a defined number of load cycles (LC). To determine the effect of different numbers of load cycles the specimens are seg-

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mented in eight areas, further designated as overrun-sections. The overrun-sections, as shown in the cross-section of SIMSPC (Fig. 4, c), pass through different amounts of load cycles from one up to 1,000. Moreover, one section (Fig. 4, c; 0 LC) is left out by the load and used as reference of texture after finishing the loading tests.

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Fig. 4 - a) Geometry of hard metal – specimen (HM-SPC); b) Geometry of strain induced martensite – specimen (SIM-SPC); c) Cross-section of SIM-SPC with marked overrun-section of different load cycles. An examination of the surface evolution, documented with the digital microscope, in regards of asymmetrical degradation marks ensures a uniform distribution of the loading in the contact of SIM-SPC and HM-SPC. Preparation and etching Subsequent to cutting and embedding, the microsections were mechanically ground and polished, using silicon carbide abrasive discs with grain sizes from 120-1200 for grinding and diamond slurry from 9-1 micrometre on synthetic-fibre cloths for polishing. The polished microsections were submerged in Kane etchant, consisting of 96 ml H2O, 60 ml of 37% HCl and 6 g CuCl2, for 40 seconds at a temperature of 25°C. In comparison to Kalling I etchant (50 ml H2O, 33 ml 96% ethanol, 33 ml 37% HCl and 1,5 g CuCl2), Lichtenegger-Bloech etchant (100 ml H2O, 20 g NH4HF2, 0.5 g K2S2O5) and modified Vilella etchant (150 ml 96% Ethanol, 10 ml 37% HCl and 5 ml 65% HNO3), Kane etchant led to a better visible martensitic phase as well as grain boundaries. In the digital micrographs α’-martensite becomes visible as black shaped laths, and shear

bands appear as more slightly etched thin lines in the austenitic microstructure. Microscopy was operated utilizing an optical microscope (DMREM by Leica). Mounting the microsections on a rotary positioning stage, such that the positioning notch (Fig. 4, c) and the positioning stage's 0° mark are congruent, enabled the linking of the digital micrographs with the equivalent overrun-section. Design of experiment Due to results of former investigations [12] the testing field is based on the three influencing factors, normal force FN, slip s and surface roughness Ra. The associated factor levels are given in Table 2. According to Karas [17] the theoretical maximum shear stress τmax determined by the shape of specimen and stages of normal force FN are τmax = 394 MPa for FN= 454 N, τmax = 363 MPa for FN = 373 N and τmax = 326 MPa for FN = 300 N. All factors are tested in a full factorial design of experiment.

Tab. 2 - Testing levels of influencing factors FACTOR LEVELS

factor

level 1

level 2

level 3

normal force FN

300 N

373 N

454 N

slip s

roughness Ra

0.3 µm

0.6 µm

0.9 µm

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Attualità industriale Further settings are rotational speed of specimen B with a constant level of nB = 0.6 rps and an ambient temperature of about TU = 23 °C. The tribological contact between the specimens is not lubricated. Each revolution of the specimen is followed by a break of two seconds to avoid a heating of the specimen and to ensure a proper measurement of the martensitic content with the FERITSCOPE® MP30. RESULTS Exemplary martensitic evolution Figure 5 shows the change of count rate ΔX, deformation of specimen Σ and friction torque MR represented from zero up to 1,000 load cycles, exemplarily for a test run at FN = 454N, s = 0 % and surface roughness Ra = 0.3 µm. During a test run the change of count rate ΔX increases in a slightly declining manner, which indicates the advancing generation of martensitic structures. As well there is a significant increase of deformation Σ during the first 40 to 60 load cycles. After that running-in phase no further deformation can be detected. The friction torque remains constant besides a slightly decreasing

behavior during the test run. During the running-in phase there is a correlation between the deformation Σ and the change of count rate ΔX. As Σ is increasing rapidly ΔX stays constant. The measured increase of ΔX and with it the generation of martensitic structures are confirmed by the metallurgical evaluation of SIM-SPC. As seen in Figure 6 for a test with FN = 454N, s = 0 % and surface roughness Ra = 0.3 µm, the overrun-sections with higher numbers of passed load cycles show higher amounts of martensitic structures. Sections with less than 50 load cycles show no significant martensitic structures like the process of ΔX already showed. Martensitic evolution starts as strayed laths in detached grains, as Figure 6 (50 cycles) displays. Driven by further load repetitions more laths occur in a higher amount of grains. Also, the area of concerning grains spreads deeper into the specimen. The figures furthermore indicate a small band near the specimen surface which shows no martensitic structures.

Fig. 5 - Development of the magneto – inductive measured change of count rate ΔX, the deformation of specimen Σ in µm and the friction torque MR in Nm.

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Fig. 6 - Evolution of strain induced martensitic structures seen in an etched cross section after 10, 50, 100, 500 and 1,000 load cycles at FN = 454 N, s = 0 % and roughness Ra =0.3 Âľm.

In order to evaluate the information of the polished and etched cross-section the measurands tmin, tmax and LD are defined. tmin describes the minimum and tmax the maximum depth of martensitic structures relative to the specimen surface (Fig. 7). The

lath density LD is a semi-quantitative measurand for the count of lath per area and is determined by the references LD 0 to LD 4 shown in Figure 7.

Fig. 7 - References for measurands tmin and tmax as well as lath density LD 0 to LD 4 of strain induced martensitic structures in an etched cross section.

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Attualità industriale Influence of testing parameters The statistical analysis of the experimental design shows significant influences on the martensitic evolution for normal force FN and number of load cycles LC. The influence of slip s and surface roughness Ra are indifferent or not detectable within the parameters of the performed tests. Due to this fact, the analysis is reduced to the levels of normal force FN = (454/ 373/ 300) N. The results of all runs on each stage of FN are ave-

raged. Figure 8 shows these force-averaged results for tmin, tmax, LD and ΔX displayed over increasing number of load cycles LC. Each measurand is calculated for discrete supporting points, respectively load cycles (0, 1, 5, 10, 50, 100, 500 and 1,000). The error bar represents a confidence level of 68.3%.

Fig. 8 - Evolution of lath density LD, change of count rate ΔX, tmin and tmax averaged over the factor levels of the normal force FN and display over load cycles. Like the exemplary results shown in Figure 5 the analysis of the complete experimental design displays an increasing amount of martensitic structures with increasing number of load cycles. In detail a higher normal force FN leads to a higher level of lath density LD and change of count rate ΔX. The lath density proceeds in a declining manner. The maximum depth of martensitic structures tmax behaves similarly and shows a significant dependency on normal force FN, too. Additionally, it can be observed that the difference between tmax for FN = 454 N and FN = 373 N is significantly smaller than between FN = 373 N and FN = 300 N. tmin shows an analogous but inverted behaviour concerning the influence of increasing normal force as tmax. Thus, the minimum depth of martensitic structures decreases with increasing normal force and number of load cycles. An overall analysis of the experimental design shows that the first ascertainable occurrence of martensitic structures is after the previously described deformation in the running-in phase. This means during the main evolution of martensite no further

La Metallurgia Italiana - n. 1 2018

measurable deformation occurs. Additional investigations of the polished and etched texture of specimens’ cross-sections show a significant difference between the hardness of the identified martensitic structures and the original austenitic texture. Grains with a high amount of martensite-lath show a hardness of H = (296 ± 5) HV0.1 and the original texture a hardness of H = (180 ± 1) HV0.1 with a confidence level of 68.3% each. Nucleation of martensite Figure 9 displays the observed nucleation sites of the martensitic phase in the austenitic microstructure, which are in accordance with the nucleation sites studied in former research projects [4]. Figure 9 (a) depicts the nucleation of SIM at grain boundaries. It is evident that the martensite-laths, starting at the grain boundary, spread solely into one grain, whereas in the other grain martensitic structures cannot be observed. The nucleation in twins is apparent from Figure 9 (b); here the

Industry news martensite-laths stand approximately perpendicular on the twin boundary. In Figure 9 (c) the nucleation of α’-martensite at shear band intersections is illustrated, and the α’-crystals are visible as punctate structures at the crossover of two shear bands. The observed angles included by a pair of shear bands

varies in the range of 60-70°. The α’-nuclei form predominantly in the direction of one shear band. Furthermore, the beginning evolution of a martensite-lath by coalescence of several nuclei can be observed in the highlighted microsection region in Figure 9 (c).

Fig. 9 - Exemplary nucleation of martensitic structures: a) grain boundaries; b) twins; c) shear band crossings

DISCUSSION Several investigations like Talonen [14] or Hecker [18] show that the amount of strain induced martensite increases with increasing degree of plastic strain. In contrast to these results, which were mostly generated by tensile testing, there is no observable increasing plastic strain within the tests detailed in this paper. But the discovery of shear bands in the etched cross-sections and the identification of martensitic nucleation in shear band crossings points to the possibility of an advancing and discontinuous strain process which leads to a similar result as known from continuous tensile tests. During that process the dislocation density increases and generates enough shear bands for the nucleation and growth of martensitic structures. A similar accumulation of strain is known from HCF tests, which leads to an increase of martensitic content with increasing number of load cycles [19], too. An interesting phenomenon is the running-in behavior of the martensite evolution. The strict order of enclosed deformation Σ and occurrence of first martensitic structures also indicates that a basic level of dislocation density is needed for the strain induced martensite transformation which fits to the model of Olson and Cohen [20]. Due to the uncertainties of the used measuring methods (dependence of the FERITSCOPE® on change of surface roughness and uncertain detection of shear bands by etching with Kane) further investigations are required for a verified statement. The change of position, the expansion and the lath density of

the martensitic structures correlate with the change of position, and expansion of the stress field generated in the specimens contact. Due to the shape of specimen and value of normal force FN the stress field can be calculated according to Karas [17]. Conformable to the calculation, higher normal forces lead to higher maximal stresses which reach in deeper areas. Beneath the influence of the macroscopic factor of the design of experiment the martensitic structures show a strong dependency on grain orientation. This is represented through the few and scattered grains showing martensitic laths at low number of load cycles as through single austenitic grains in an elsewise martensitic area remain after 1,000 load cycles. Moreover, the orientation of laths in single grains shows that some shear systems accommodate deformation easier than others or fit better to the applied stress field. The detected angle between crossing shear bands or crossing laths is between 60 and 70 degrees and fits to the angle between the densest sphere packing of γ-austenite. An open-ended question is the small remaining band of austenite near the surface above the martensitic structures described by tmin. An explanation with the calculated field of stress is not adequate. The stress levels between the surface and tmin is mostly equivalent or higher than the stress level of deeper areas which pass through the strain induced martensite transformation. The best approach is a combination of the influence of grain structure, austenite stability and stress field. At the actual

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Attualità industriale state of investigation, a proper conclusion is not possible.

densities and maximum depth of martensitic structures. Lath density and maximum depth of martensitic structures increase with increasing number of load cycles. Further investigations concerning different kinds of stress variation and austenite stability are in process and will provide a more detailed view on the topic. Moreover, analyses concerning the grain structure like local distribution of grain size and micro hardness profiles will be carried out as well as numerical analyses of the stress field for more detailed correlations. •

CONCLUSION Metastable austenitic steels form significant martensitic structures in low stress rolling contacts. The occuring structures show different stages of lath density LD and maximum depth of martensitic structures tmax. These descriptive parameters are mainly influenced by the factors normal force FN and number of load cycles. This investigation leads to the following conclusions: • High normal force or high level of stress leads to high lath

ACKNOWLEDGEMENT This study is part of the research project "AREWESI" supported by the German Federal Ministry of Education and Research (BMBF).

REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]

M. BIGDELI KARIMI, H. ARABI, A. KHOSRAVANI and J. SAMEI, Mat. proc. 203(1), (2008), p. 349 D. MEYER, Forschungsbericht Stiftung Institut für Werkstofftechnik Bremen 57, (2012) G. B. OLSON and M. COHEN, journal of less-common metals 28, (1972), p. 107 A. DAS, S. TARAFDER, int. jou. of plast., 25(11), (2009), p. 2222 M. R. DA ROCHA, C. A. S. DE OLIVIERA, mat. sc. and eng., 517(1-2), (2009), p. 281 P. L. MANGONON, G. THOMAS, met. trans. 1(6), (1970), p. 1577 J. TALONEN, H. HÄNINEN, acta. met. 55(18), (2007), p. 6108 L. E. MURR, K. P. STAUDHAMMER, S. S. HECKER, met. trans. 13(4), (1982), p. 627. A. HEDAYATI, A. NAJAFIZADEH, A. KERMANPUR, mat. proc. 210(8), (2010), p. 1017 M. WERSCHLER and P. GÜMPEL, Proc. 56. Tribologie Fachtagung 2, (2015), p. 79 N. N., Tribologische Prüfstände – Zweischeibenprüfstand; www.gft-ev.de, (2016) M. WERSCHLER, K. LUTHER and P. GÜMPEL, Proc. 12. Arnold-Tross-Kolloquium, (2016), p. 10 N. N., FERITSCOPE® manual, (2016) J. TALONEN, A. ASPEGREN and H. HÄNNINEN, mat. science and technology 20, (2004), p. 1506 K. WERNER, HTWG Konstanz project thesis, (2016) K. NOHARA, Y. ONO and N. OHASHI, Tetsu-to-Hagane, (1977), p. 5 F. KARAS, Forschung auf dem Gebiet des Ingenieurwesens A 12.6, (1941), p. 266 S. S. HECKER, M. G. STOUT, K. P. STAUDHAMMER, J. L. SMITH, met. trans. 13(4), (1982), p. 619 H. J. CHRIST, U. KRUPP, C. MUELLER-BOLLENHAGEN, I. ROTH, Proc. ICF12, (2009), p. 2013. G. B. OLSON and M. COHEN, metallurgical transactions 6A, (1975), p. 791

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Experts’ corner Le norme EN 10088 per gli acciai inossidabili a cura di Mario Cusolito

Dr. Ing. Mario Cusolito • Presidente della SC 25 Unsider "Acciai da trattamento termico, acciai legati, acciai automatici e acciai inossidabili" • Membro di ECISS TC105 "Steels for heat treatment, alloy steels, free-cutting steels and stainless steels" • Membro di ISO TC17 SC4 "Heat Treatable and Alloy Steels"

e norme EN 10088 per gli acciai inossidabili Le EN 10088-1, -2, -3 sono le norme base europee per gli acciai inossidabili: dal 1995 hanno sostituito le norme nazionali, quali la DIN 17440 – fino ad allora la norma di riferimento europea – e, tra le altre, la UNI 6901. I tedeschi utilizzavano già il loro “Werkstoffnummer”, diventato oggi il “numero acciaio”, che permette di identificare facilmente il tipo di materiale con un semplice numero a 4 cifre, mentre le designazioni alfanumeriche sono particolarmente complesse per gli acciai inossidabili, tanto che in Italia e in Francia si usavano le designazioni americane AISI, con composizioni non perfettamente corrispondenti. La prima edizione delle norme EN 10088 nel 1995 Già nella prima edizione del 1995, tra72

dotta in italiano nel 1997, la struttura delle norme era simile a quella attuale, con un elenco degli acciai inossidabili (parte 1), una parte 2 dedicata ai prodotti piani e una parte 3 riferita ai prodotti lunghi. La norma, nata prevalentemente da compromessi tra produttori tedeschi, svedesi e inglesi, ebbe un immediato successo e rapidamente sostituì le norme nazionali, secondo gli accordi e i regolamenti europei. La revisione delle norme EN 10088 nel 2005 Dopo i cinque anni di validità, le tre parti della norma EN 10088:1995 furono messe in revisione nel 2000 e le nuove norme furono pubblicate nel 2005: il lungo periodo di gestazione, pur trattandosi di una semplice revisione, fa capire l’interesse e gli sforzi fatti per pubblicare un documento valido e

condiviso. La nuova edizione fu recepita da UNI in lingua inglese nel 2005 e successivamente tradotta in italiano nel 2008: la struttura rimase la stessa, ma diversi miglioramenti vennero apportati aggiungendo nella parte 1 parecchi acciai tra cui quelli resistenti al creep; la parte 2 rimase sostanzialmente invariata mentre le parte 3 fu sensibilmente ampliata per comprendere i requisiti dei prodotti finiti a freddo (caratteristiche meccaniche nelle finiture 2H, 2B, 2G, 2P per i prodotti lunghi). In aggiunta, nella parte 3 è stato introdotto per la prima volta in concetto, mutuato dalla allora recente norma EN 10277, che non si possono avere acciai con zero difetti superficiali: questa piccola rivoluzione ha cambiato l’approccio degli utilizzatori di prodotti lunghi in acciaio inossidabile, costringendoli a pensare in termini più realistici alle problematiche di sanità superficiale. La Metallurgia Italiana - n. 1 2018

Scenari La attuale norma EN 10088:2014 La successiva revisione delle EN 10088:2005, messe in revisione nel 2010, fu pubblicata nel 2014, ed è attualmente in vigore. Durante la revisione sono stati proposti moltissimi nuovi acciai: per evitare un inutile proliferare di doppioni e di acciai non utilizzati, è stato necessario definite alcune regole, adottate poi in generale da tutti i comitati di normazione. Lo sviluppo tecnologico degli acciai inossidabili in relazione alle possibili applicazioni specifiche ha portato in effetti a mettere a punto molti nuovi tipi, con composizione chimica mirata, che gli inevitabili interessi industriali avevano già portato alla protezione tramite brevetti. Una norma deve comprendere anche i nuovi tipi di acciaio, che vanno nella direzione dello sviluppo delle applicazioni, anche considerando che i tempi di revisione sono lunghi e un acciaio non introdotto nella norma ne resta escluso per 7-8 e anche 10 anni; talvolta però la richiesta di introdurre un acciaio nuovo potrebbe essere prematura in quanto quel tipo è stato messo a punto di recente, magari brevettato, ma prodotto in quantità minime e per applicazioni specifiche che non possono essere considerate “standard”, cioè normate. Alla fine il criterio di non lasciare fuori dalla norma gli acciai “nuovi” ha prevalso, con un sensibile incremento del numero degli acciai rispetto all’edizione 2005 (ben 22 acciai in più – da 160 a 182) e soprattutto di quelli brevettati (da 5 a 12). Nel contempo alcuni acciai desueti (11, al momento) sono stati spostati in una speciale lista (“Uncommon Grades”): se non torneranno di moda nel prossimo futuro, saranno totalmente eliminati con la prossima revisione. L’eliminazione d’altra parte non è semplice perché molti degli acciai più vecchi sono entrati in

La Metallurgia Italiana - n. 1 2018

leggi e regolamenti europei (per esempio Regolamento PED per recipienti in pressione) o dei singoli stati (per esempio, lista positiva di acciai destinati al contatto con gli alimenti) e non possono essere eliminati dalla norma base. La norma EN 10088-1:2014 contiene molti dati utili ma poco consultati dagli utenti come, ad esempio, la classificazione degli acciai inossidabili in base alle proprietà di utilizzo, alla microstruttura e agli elementi di lega significativi (allegato C), le formule empiriche per la classificazione della microstruttura e la valutazione della resistenza al pitting (allegato D), alcune proprietà fisiche come densità, modulo di elasticità, coefficiente di espansione termica, conducibilità termica, capacità termica specifica, resistività elettrica (allegato E).

“armonizzate” alla direttiva CPD (Construction Product Directive) e pertanto contengono un allegato ZA che elenca i requisiti per il rilascio della dichiarazione di conformità. Le due norme sono state emesse nel 2009 e confermate nel 2013, in attesa di precisazioni circa lo sviluppo dei regolamenti europei. Saranno messe in revisione durante il 2018.

Il rapporto con la norma ISO 15510 e la possibile unificazione L’allegato A della EN 10088-1:2014 riporta una comparazione tra le designazioni degli acciai delle più importanti norme: internazionali (ISO), statunitensi (ASTM/UNS), europee (EN), giapponesi (JIS) e cinesi (ISC) e fa riferimento ISO 15510:2014 “Elenco degli acciai inossidabili” per la spiegazione della designazione ISO. Contrariamente a quanto si sta già facendo per altri acciai, nonostante tutti gli sforzi, appare altamente improbabile che in tempi brevi si possa arrivare ad una unificazione degli acciai inossidabili EN con quelli ISO, sia per la mancanza di cooperazione con gli enti di unificazione statunitensi, sia per le citate oggettive difficoltà dovute all’inserimento degli acciai europei nei regolamenti ufficiali.

Conclusioni Grazie alla norma EN 10088 nelle sue tre parti, oggi l’Europa dispone di una norma forte, condivisa dalla maggior parte delle nazioni più industrializzate e assolutamente all’avanguardia nel mondo. La norma, nel suo sviluppo ormai ventennale, ha raggiunto una diffusione capillare ed è diventata un punto di riferimento europeo spendibile anche in sede extracomunitaria grazie alla sua chiarezza ed efficacia.

Il futuro delle EN 10088 A seguito della richiesta di alcuni stati membri di introdurre nuovi acciai e di effettuare alcune piccole modifiche, le norme EN 10088-1, -2 e -3 sono state poste in revisione poco dopo la loro pubblicazione; al momento però nessuna azione pratica è stata effettuata per iniziare il processo di revisione.

Le norme EN 10088-4 e -5 Queste due norme riguardano i prodotti piani e lunghi in acciaio inossidabile destinati alle costruzioni: sono norme

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La Metallurgia Italiana - n. 1 2018

Atti e notizie AIM – UNSIDER

Norme pubblicate e progetti in inchiesta (aggiornamento 29 dicembre 2017)

Norme UNSIDER pubblicate da UNI nel mese di dicembre 2017 UNI EN 12681-2:2017 Fonderia - Controllo radiografico - Parte 2: Tecniche con rilevatori digitali UNI EN 10207:2017 Acciai per recipienti a pressione semplici - Condizioni tecniche di fornitura per lamiere, nastri e barre UNI EN 10263-3:2017 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Parte 3: Condizioni tecniche di fornitura degli acciai da cementazione UNI EN 10263-2:2017 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Parte 2: Condizioni tecniche di fornitura degli acciai non destinati al trattamento termico dopo lavorazione a freddo UNI EN 10263-1:2017 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Parte 1: Condizioni tecniche generali di fornitura

EN ISO 19901-2:2017 Petroleum and natural gas industries - Specific requirements for offshore structures - Part 2: Seismic design procedures and criteria (ISO 19901-2:2017) CEN ISO/TS 17969:2017 Petroleum, petrochemical and natural gas industries Guidelines on competency management for well operations personnel (ISO/TS 17969:2017) ISO 4545-2:2017 Metallic materials -- Knoop hardness test -- Part 2: Verification and calibration of testing machines ISO 4701:2017 Iron ores and direct reduced iron -- Determination of size distribution by sieving ISO 4545-4:2017 Metallic materials -- Knoop hardness test -- Part 4: Table of hardness values ISO 4545-1:2017 Metallic materials -- Knoop hardness test -- Part 1: Test method

Norme UNSIDER ritirate da UNI nel mese di dicembre 2017

ISO 35106:2017 Petroleum and natural gas industries -- Arctic operations -- Metocean, ice, and seabed data

EC 1-2014 UNI EN 10263-3:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura degli acciai da cementazione

ISO 945-1:2017 Microstructure of cast irons -- Part 1: Graphite classification by visual analysis

UNI EN 10207:2005 Acciai per recipienti a pressione semplici - Condizioni tecniche di fornitura per lamiere, nastri e barre UNI EN 10263-3:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura degli acciai da cementazione UNI EN 10263-2:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura degli acciai non destinati al trattamento termico dopo lavorazione a freddo UNI EN 10263-1:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura generali Norme UNSIDER pubblicate da CEN e ISO nel mese di dicembre 2017

La Metallurgia Italiana - n. 1 2018

ISO 4545-3:2017 Metallic materials -- Knoop hardness test -- Part 3: Calibration of reference blocks ISO 35103:2017 Petroleum and natural gas industries -- Arctic operations -- Environmental monitoring Progetti UNSIDER messi allo studio dal CEN (Stage 10.99) – dicembre 2017 EC102099 Iron and steel - European standards for the determination of chemical composition Progetti UNSIDER in inchiesta prEN e ISO/DIS – gennaio 2018 prEN – progetti di norma europei prEN ISO 19904-1 Petroleum and natural gas industries - Floating offshore structures - Part 1: Ship-shaped, semi-submersible, 75

Aim news spar and shallow-draught cylindrical structures (ISO/DIS 19904-1:2017)

systems for floating offshore structures and mobile offshore units

prEN ISO 19901-7 Petroleum and natural gas industries - Specific requirements for offshore structures - Part 7: Stationkeeping systems for floating offshore structures and mobile offshore units (ISO/DIS 19901-7:2017)

ISO/DIS 19904-1 Petroleum and natural gas industries -- Floating offshore structures -- Part 1: Ship-shaped, semi-submersible, spar and shallow-draught cylindrical structures

prEN ISO 19345-2 Petroleum and natural gas industry - Pipeline transportation systems -Pipeline integrity management specification - Part 2: Full-life cycle integrity management for offshore pipeline (ISO/DIS 19345-2:2017) prEN 15655-1 Ductile iron pipes, fittings and accessories - Requirements and test methods for organic linings of ductile iron pipes and fittings - Part 1: Polyurethane lining of pipes and fittings prEN 598 Coated and lined ductile iron pipes, fittings and their joints for sewerage and drainage applications - Products characteristics and test and assessment methods EN 13480-1:2017/prA1:2017 Metallic industrial piping - Part 1: General prEN ISO 4885 Ferrous materials - Heat treatments - Vocabulary (ISO/ FDIS 4885:2017) prEN ISO 6892-2 Metallic materials - Tensile testing - Part 2: Method of test at elevated temperature (ISO/FDIS 6892-2:2017) prEN ISO 4945 Steel - Determination of nitrogen - Spectrophotometric method (ISO/DIS 4945:2017) prEN 10225-4 Weldable structural steels for fixed offshore structures - Technical delivery conditions - Part 4: Cold formed welded hollow sections prEN 10253-2 Butt-welding pipe fittings - Part 2: Non alloy and ferritic alloy steels with specific inspection requirements

ISO/DIS 4978 Steel sheet and strip for welded gas cylinders ISO/DIS 19345-2 Petroleum and natural gas industry -- Pipeline transportation systems -- Pipeline integrity management specification -- Part 2: Full-life cycle integrity management for offshore pipeline ISO/DIS 19345-1 Petroleum and natural gas industry -- Pipeline transportation systems -- Pipeline integrity management specification -- Part 1: Full-life cycle integrity management for onshore pipeline ISO/DIS 20915 Life cycle inventory calculation methodology for steel products ISO/DIS 4945 Steel -- Determination of nitrogen -- Spectrophotometric method ISO/DIS 1143 Metallic materials -- Rotating bar bending fatigue testing ISO/DIS 20064 Metallic materials -- Steel -- Method of test for the determination of brittle crack arrest toughness, Kca Progetti UNSIDER al voto FprEN e ISO/FDIS – gennaio 2018 FprEN – progetti di norma europei FprEN ISO 4829-1 Steel and cast iron - Determination of total silicon contents - Reduced molybdosilicate spectrophotometric method - Part 1: Silicon contents between 0,05 % and 1,0 % (ISO/FDIS 4829-1:2016)

prEN 10253-4 Butt-welding pipe fittings - Part 4: Wrought austenitic and austenitic-ferritic (duplex) stainless steels with specific inspection requirements

FprEN 10277 Bright steel products - Technical delivery conditions

ISO/DIS – progetti di norma internazionali

ISO/FDIS 4829-1 Steel and cast iron -- Determination of total silicon contents -- Reduced molybdosilicate spectrophotometric method -- Part 1: Silicon contents between 0,05 % and 1,0 %

ISO/DIS 21809-11 Petroleum and natural gas industries -- External coatings for buried or submerged pipelines used in pipeline transportation systems -- Part 11: Coating repairs on rehabilitation ISO/DIS 19901-7 Petroleum and natural gas industries -- Specific requirements for offshore structures -- Part 7: Stationkeeping 76

ISO/FDIS – progetti di norma internazionali

ISO/FDIS 6892-2 Metallic materials -- Tensile testing -- Part 2: Method of test at elevated temperature

La Metallurgia Italiana - n. 1 2018

trentasettesimo convegno nazionale

AIM

Bologna 12-13-14 settembre 2018 PRIMO ANNUNCIO E RICHIESTA DI MEMORIE Gli interessati a presentare memorie scientifiche (sia per le sessioni orali che per la sessione poster) dovranno inviare entro il 27 marzo 2018, il titolo della memoria, i nomi degli autori e la loro affiliazione ed un sommario di circa 300 parole. Ci sono due modi per sottoporre le proposte di memorie: - compilando il form online presente sul sito dell’evento: www.aimnet.it/37aim.htm - inviando tutte le informazioni (titolo, autori, recapiti del relatore e sommario) a mezzo e-mail: info@aimnet.it

SPAZIO AZIENDE E SPONSORIZZAZIONE È previsto uno spazio per l’esposizione di apparecchiature, per la presentazione dei servizi e per la distribuzione di materiale promozionale. Le aziende interessate potranno richiedere informazioni più dettagliate sullo spazio aziende e sulle diverse possibilità di sponsorizzazione dell’evento alla Segreteria AIM (info@aimnet.it – tel. 02 76021132).

www.aimnet.it/37aim.htm

7th international congress on science and technology of steelmaking the challenge of industry 4.0 13-15 June 2018 Venice - Italy

introduction

call for papers - abstract submission

AIM will host as an integral part of ICS 2018 its traditional Conference on Heat Treatments - XXVI Convegno Nazionale Trattamenti Termici - which will start on 13 June, with a whole day in Italian language for Italian technicians and researchers and will go on the following days with technical sessions exclusively in English language. Therefore the topics of ICS 2018 will be integrated with the topics of Heat Treatments and Surface Engineering.

deadlines

The 7th International Congress on Science and Technology of Steelmaking (ICS 2018) will be organized by AIM, the Italian Association for Metallurgy, in Italy in June 2018. ICS 2018 will provide a forum for researchers and manufacturers involved in the scientific and technical developments of steelmaking. This meeting is aimed at creating an opportunity for a technical exchange at an international level among the numerous experts involved in the steelmaking.

We kindly invite you to participate in ICS 2018 and are looking forward to meeting you in Venice!

topics

Prospective authors wishing to present papers are invited to submit a tentative title and an abstract of about 300 words (in English), specifying a maximum of two topics for each proposal, to the Organising Secretariat (aim@aimnet.it). The abstract should provide sufficient information for a fair assessment and include the title of the paper, the author’s names and contact details (address, telephone and fax numbers and e-mail address). The name of the presenting author should be underlined. A poster session might be organized as well. There are two ways to submit papers: • fill in the form on the Congress website at: http://www.aimnet.it/ics2018.htm • send the requested information by e-mail to: aim@aimnet.it.

Deadline for submission of abstracts Information on Acceptance Opening of the online registration Deadline for Full Paper Submission

January 31, 2018 February 28, 2018 February 28, 2018 April 16, 2018

language

The official language of the Congress will be English.

The following general topics will be involved in the program of ICS 2018: Science & Technology in Steelmaking • Fundamentals of Steelmaking • Thermodynamics • Thermophysical properties, Thermochemistry & Kinetics • Solidification • Slags and fluxes • Process modelling and Process control • Sensors, Measurement & Process characterisation • Electric arc furnace • Basic Oxygen Furnace • Primary and secondary steelmaking • Continuous casting • Refractory • Quality • Sustainability & Environment, Recycling and use of by-products • Industry 4.0 • Automation Heat Treatment & Surface Engineering • Thermo-chemical treatment (carburizing, nitriding, nitrocarburising,…) • Surface hardening (induction, laser,…) • Coating technology and coatings (PVD, CVD, plasma, thermal spray,…) • Design and construction of industrial heat treatment equipment • Equipment for measurement and process control • Quenching technology, equipment and quenchants • Residual stress and distortion • Tribology and tribological testing methods • Wear and wear protection • Modeling and simulation of heat treatment and surface engineering related aspects • Reliability and process control • Cost analysis and reduction in manufacturing • Energy saving • Bulk heat treatment • Cryogenic treatment • Mechanical properties

proceedings

The full texts of all papers will be published in the electronic form proceedings and issued to delegates on arrival at the Congress. A selection of the best papers will be also published in “La Metallurgia Italiana - International Journal of the Italian Association for Metallurgy” - the scientific journal of AIM, or in “Ironmaking and Steelmaking: Processes, Products and Applications”, published by Taylor & Francis on behalf of the Institute of Materials, Minerals and Mining (IOM3). Both journals are covered in the Science Citation Index Expanded by Thomson Reuters and in Scopus by Elsevier B.V.

exhibition & sponsorship opportunities

As an integral element of the event, ICS 2018 will feature an Exhibition that will enable excellent exposure for company products, technologies, innovative solutions or services. At this opportunity the Organizers will set an area strategically located. This area will be a focal point of the Congress, so that enough time will be available to guarantee a perfectly targeted potential customer’s environment. Companies will be able to reinforce their participation and enhance their corporate identification by taking advantage of the benefits offered to them as Sponsor of the Congress. Companies interested in exhibiting and/or sponsoring the event may contact:

e-mail: commerciale@siderweb.com tel: +39 0302540006 - fax +39 0302540041

organising secretariat ASSOCIAZIONE ITALIANA DI METALLURGIA ASSOCIAZIONE ITALIANA DI METALLURGIA via Filippo Turati, 8 • 20121 Milan • Italy phone: +39 0276021132 • fax +39 0276020551 e-mail: aim@aimnet.it website: www.aimnet.it