La Metallurgia Italiana - Febbraio 2018 by aimnet3

Metallurgia Italiana

International Journal of the Italian Association for Metallurgy

n. 2 Febbraio 2018 Organo ufficiale dellâ&#x20AC;&#x2122;Associazione Italiana di Metallurgia. Rivista fondata nel 1909

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

introduction

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.

registration is open!

registration fees

(early bird registration fees by april 30, 2018) MEMBER

NON MEMBER

€ 560

€ 680

€ 700

€ 820 € 380 € 720

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.

Speaker (Presenter) Session Chairperson Committee member Participant (non-Presenter) Student Exhibitor / Sponsor

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

Congress registration fees include • Admittance to technical sessions and to the exhibition • Congress bag with electronic proceedings • Social event on June 14 • Coffee breaks • Lunches For non-members the fee includes AIM Membership for the last three quarters of 2018 and for the year 2019.

topics

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

FREE

€ 600

congress venue

The Conference will be held in Venice Mestre mainland, at NH LAGUNA PALACE (Viale Ancona 2, Mestre, Venice - Italy).

language

The official language of the Congress will be English.

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

www.aimnet.it/ics2018.htm

La Metallurgia Italiana

Metallurgia Italiana

International Journal of the Italian Association for Metallurgy

n. 2 Febbraio 2018 Organo ufficiale dell’Associazione Italiana di Metallurgia. Rivista fondata nel 1909

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

n. 2 Febbraio 2018

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

Anno 110 - ISSN 0026-0843

Leghe leggere / Light metals Influence of Mg and Ti on both eutectic solidification and modifying efficiency in Sr-modified Al-7Si cast alloys L. Lattanzi, A. Fortini, M. Giovagnoli, M. Merlin

Theoretical and Experimental Evaluation of the Effectiveness of Aluminum Melt Treatment by Physical Methods V.B. Deev, E.S. Prusov, A.I. Kutsenko 16 Analisi del comportamento a corrosione di campioni di alluminio AA6012 sottoposti ad ECAP e trattamento criogenico A. Viceré, M. Cabibbo, C. Paoletti, G. Roventi, T. Bellezze

LA COMMUNITY DELL’ACCIAIO

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)

Electrochemical corrosion behaviour of binary magnesium heavy rare earth alloys F. Rosalbino, S. De Negri, G. Scavino, A. Saccone 34 Attualità industriale / Industry news Manifestazioni AIM

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A comparative cradle-to gate impact assessment: primary and secondary aluminum automotive components case S. Cecchel, M. Collotta, G. Cornacchia, A. Panvini, G. Tomasoni 46 PLANBLUE – una visione della crescita sostenibile e della sua realizzazione a cura di: Ufficio Stampa Ronal Group 56 Scenari / Experts’ Corner Intervista a Roberta Niboli - Raffmetal

Atti e notizie / Aim news Calendario degli eventi internazionali / International events calendar 60 Rubrica dai Centri

Normativa

l’editoriale La Metallurgia Italiana L’attenzione che il mondo industriale, ed in particolare il settore dei trasporti, pone nei riguardi delle leghe leggere è continuamente crescente. Il quantitativo di alluminio mediamente utilizzato per applicazioni automotive è raddoppiato negli ultimi vent’anni, e svariate, autorevoli fonti stimano un significativo incremento anche nel futuro. Questo successo dipende da una concomitanza di fattori: requisiti sempre più stringenti di eco-sostenibilità (ottenibili con riduzione di pesi e di emissioni dei veicoli), sviluppo di soluzioni di processo sempre più innovative (sia in fonderia che

Prof. Franco Bonollo Università di Padova

nell’ambito delle tecnologie di deformazione plastica), messa a punto di leghe di elevate prestazioni (grazie a nuovi alliganti e a trattamenti metallurgici sempre più raffinati), contenimento dei costi (anche grazie a migliorate metodologie di riciclo). In questo scenario, l’industria italiana gioca un ruolo di rilievo in ambito europeo, grazie a prodotti di notevole qualità metallurgica, destinati a svariati ambiti applicativi, e ad una supply chain (leghe, impianti, materiali di consumo) altamente innovativa. Tutto ciò porta l’Italia ad avere, insieme alla Germania, un ruolo di leadership europea, in termini di quantitativi prodotti e di aziende presenti sul mercato. Tale ruolo è sempre più spesso associato a svariate forme di collaborazione tra imprese e mondo universitario, che vanno dalle semplici tesi di laurea ai dottorati, da contratti di ricerca a complessi progetti multi-partner, spesso cofinanziati a livello regionale, nazionale, europeo. Questo numero della Metallurgia Italiana vuole testimoniare il dinamismo che anima il settore delle leghe leggere, con una serie di contributi scientifici e tecnici incentrati sui trattamenti dell’alluminio liquido, sui processi innovativi di estrusione, sulla fonderia, su nuove leghe a base magnesio, sulle tematiche di sostenibilità ambientale, sugli scenari attuali e futuri del settore.

La Metallurgia Italiana - n. 2 2018

Leghe leggere Influence of Mg and Ti on both eutectic solidification and modifying efficiency in Sr-modified Al-7Si cast alloys L. Lattanzi, A. Fortini, M. Giovagnoli, M. Merlin Magnesium and titanium are the main alloying elements always present in a commercial A356 alloy, and strontium is commonly added to achieve a good degree of modification of the eutectic structure. Whilst most of the studies have been focused on the role of strontium on commercial A356 alloy, little attention has been paid to understand the possible interaction of magnesium and titanium with the modifying efficiency of strontium. In the light of these aspects, the present study is aimed at investigating the effects of magnesium and titanium on the microstructural evolution of Sr-modified Al-7Si, Al-7Si-0.4Mg and Al-7Si-0.4Mg-0.12Ti alloys. To assess the role of Sr on the eutectic solidification path, the analysis of the cooling curves combined with a quantitative microstructural evaluation of the eutectic silicon particles has been carried out. Thermal analysis results highlight how the alloys that contain magnesium show a significant decrease, higher than 5 Â°C, of the thermal parameters of the eutectic solidification, with respect to the Al-7Si alloy. On the other hand, titanium seems to have only a slight effect on the same solidification characteristics. Metallographic investigations and the related statistical analysis of some geometrical parameters of the silicon particles indicate that the presence of magnesium and titanium induces a variation of both dimension and morphology of the particles. As a result, these experimental findings suggest that the effect of alloying elements, such as titanium and magnesium, on the thermal parameters obtained from cooling curves has to be taken into consideration when the thermal analysis is used to assess the strontium modification efficiency.

KEYWORDS: AL-SI ALLOYS, SR MODIFICATION, THERMAL ANALYSIS, SI PARTICLES, ELEMENTS INTERACTION

INTRODUCTION Al-Si alloys pertain to the predominant aluminium alloys used for a huge variety of automotive and aircraft cast components. This is mainly due to their low weight, good castability, low cost and favourable mechanical properties which depend on the microstructure resulting from the solidification process. In this respect, a considerable amount of studies has concerned the improvement of their characteristic by melt inoculation [1,2], by alloying [3] and by tuning the parameters of the heat treatment process [4]. In order to control the solidification behaviour of an alloy, Thermal Analysis (TA) has been proved to be an effective technique for the metal casting industry. This non-destructive and rapid on-line monitoring method enables, in fact, to evaluate the melt quality and to monitor the processing parameters prior to casting. TA can provide information about the degree of grain refinement and modification [5,6], the characteristic temperatures related to the solidification regions of both primary and eutectic phases [7,8] and the intermetallic phases formation [9]. As regards the relationship between the refinement of the eutectic silicon phase and the related changes in the cooling curves, this aspect has been extensively investigated in the literature. The depression of the eutectic growth temperature has La Metallurgia Italiana - n. 2 2018

been used to assess the modification level of the melt, thus suggesting a correlation between thermal and microstructural parameters of the eutectic phase [10]. By contrast, there is a relative paucity of scientific studies focused on the interaction between modifiers and alloying elements and its effect on TA cooling curves [11,12]. The present study sets out to experimentally investigate the influence of magnesium and titanium on the microstructural evolution of strontium modified Al-7Si, Al-7Si-0.4Mg and Al-

L. Lattanzi, A. Fortini, M. Giovagnoli, M. Merlin University of Ferrara, Department of Engineering (DE) via Saragat 1, I-44122 Ferrara, Italy

Light metals 7Si-0.4Mg-0.12Ti alloys. The analysis of the cooling curves and their derivatives, examined with a tailored MatlabÂŽ code, provides evidence of the role of magnesium and titanium additions. The depression of the eutectic growth temperature has been considered in the light of the microstructural investigations, conducted on the eutectic silicon particles. A comprehensive quantitative image analysis, supported by an extensive statistical approach which concerns the comparison of the distributions of equivalent diameter and roundness of the particles, clarified the mutual effect of the alloying elements. LITERATURE SURVEY The control of the microstructure is a key aspect to enhance mechanical properties and casting quality of aluminium alloys. Within this frame, it is well-established that the size and the morphology of the silicon particles can impair the mechanical properties of Al-Si alloys, especially ductility, due to the needle-like structure arising from nucleation and growth of the eutectic phase. As a result, the chemical modification has become a common foundry practice to promote the structural transformation of the silicon particles from a coarse plate-like structure into a fine fibrous one. The modification is commonly performed by the addition of modifying agents in the molten metal to determine changes in the growth kinetics of the eutectic phase. In recent years, strontium has become the most widely used modifying agent thanks to its good modification rate, long fading time, high recovery efficiency and ease of handling [8,13,14]. A great deal of efforts has been made to optimise the modification treatment and many studies have been addressed to investigate the microstructural changes related to strontium additions, modification level assessment and mechanical properties of strontium modified hypoeutectic Al-Si alloys [15,16]. To promote the precipitation of Mg2Si during the heat treatment and, in turn, improve the mechanical properties of Al-Si foundry alloys, small amounts of magnesium are usually introduced. There is no dearth of experimental studies dedicated to understanding the influence of magnesium on the microstructural evolution and tensile properties of hypoeutectic Al-Si alloys. Some studies have evaluated the effects of magnesium additions on the mechanical behaviour of both unmodified and strontium modified alloys [3,17]. More recently, some studies also investigated the role played by magnesium in influencing the formation of intermetallic compounds [18,19] and the fracture behaviour [20]. Besides the benefits on mechanical properties, it has been suggested that magnesium additions up to 1 wt. % slightly increase the modification level of silicon particles [21]. Furthermore, whether sodium or strontium are not present, magnesium enables a change in morphology, from coarse lamellar to acicular, without achieving a fibrous structure though, and thus shows a weak modifying effect [11,22]. Another way to improve mechanical properties of Al-Si alloys is reducing the primary aluminium grain size. Therefore, the addition of titanium to the melt is a common foundry prac6

tice because of its potential grain refining effect [23]. Despite this, some recent works show how titanium actually refines the grains of primary aluminium phase but has no significant effect on Secondary Dendrite Arm Spacing (SDAS), thus it determines only a slight enhancement of the mechanical properties [7]. The presence of titanium also shows some influences on thermal analysis parameters. Xu et al. reported how the addition of titanium, from 0.2 wt. % to 0.8 wt. %, to an A357 alloy causes the rise of primary phase characteristic temperatures and a depression of recalescence [7]. Other works highlight how titanium can also influence the eutectic region of the cooling curve, decreasing the characteristic temperature parameters [24]. For what concerns quantitative microstructural investigations, in the last years the effect of both alloying elements and heat treatment processes have been deepened by image analysis coupled with statistical approaches. In these regards, particular attention has been paid to quantitative image analysis and silicon particle distribution. Alexopoulos et al. found a correlation between silicon particle size and average elongation, concluding that the addition of alloying elements is reflected by variations of size distribution of silicon particles [25]. TiryakioÄ&#x;lu investigated solution treatments at 540 Â°C for different treatment durations and evaluated their effect on size and aspect ratio distributions of eutectic silicon. In particular, the reported studies find that the 3-parameter lognormal distribution provides the best fit for both equivalent diameter and aspect ratio [26]. Otherwise, a certain number of studies evaluated the effect of different grain refiners [27], alloying elements [22], combined modifying elements and solidification rate [8,28] simply considering variation in the mean values of characteristic parameters of silicon particles, coupled with their standard deviation. On the other hand, some authors considered the median values of the above-mentioned parameters, because of the large scattering of mean values [10,29]. In the present study, comparisons between Si particles parameters distributions were evaluated considering, as reported by Wilcox [30,31], the differences between the deciles of the distributions, i.e. the values that divide data into ten equal parts so that each part represents 1/10 of the population. The analysis of the cooling curves has shown to be an effective approach to control and optimise the solidification process and TA is widely used in foundry practice to evaluate the degree of modification of hypoeutectic silicon alloys. In particular, the difference between the eutectic growth temperature of the unmodified and of the modified alloy is widely used to evaluate the modification level [9,10]. Furthermore, other temperature and time-related parameters, e.g. recalescence and duration of the eutectic plateau, have been suggested for the control of silicon modification [8,32,33]. Among the experimental variables and issues that can affect the results (e.g. possibility of comparison with the cooling curve of the unmodified melt, cooling rate variability, melt and crucible temperature stability), the interaction of alloying elements on the solidification path has not been extensively examined so far. Heusler and Schneider [11] performed a systematic investigation by means of cooling curve La Metallurgia Italiana - n. 2 2018

Leghe leggere analysis on the influence of magnesium on the modification efficiency of sodium and strontium in an Al-11%Si alloy. Tahiri et al. [12] explored the influence of the combined addition of grain refiners and strontium on cooling curves and microstructure of the A356 alloy, reporting that the partial reaction between TiB2 and strontium would lead to a partial decrease of the modifying efficiency of strontium. Despite the fact that the influence of alloying elements on the microstructural and mechanical features of Al-Si alloys it is well-established and TA has been gaining increasing acceptance as an effective melt quality control, few studies have been focused on the effects of chemical composition on cooling curves parameters. In particular, very little attention has been paid to the assessment of the modifying efficiency of strontium via TA with respect to the interaction of the alloying elements [11,34]. In the light of these aspects, this study examines the changes in the eutectic phase solidification of strontium modified alloys arising from magnesium and titanium additions. In an attempt to provide a quantitative evidence of the eutectic phase changes, a combined approach based on cooling curves analysis and quantitative metallographic investigation on eutectic silicon particles has been adopted.

EXPERIMENTAL PROCEDURE Melt preparation Three different reference alloys were prepared: Al-7Si, Al-7Si0.4Mg and Al-7Si-0.4Mg-0.12Ti. For the Al-7Si alloy, primary aluminium ingots were melted in an electric resistance furnace and pure silicon was then added to the bath. For the Al-7Si0.4Mg alloy, pure magnesium was also added to reach the targeted nominal content of ~ 0.4 wt. %. For the Al-7Si-0.4Mg0.12Ti, pure magnesium and AlTi10 master alloy rods were added to reach the targeted magnesium and titanium nominal content of ~ 0.4 wt. % and ~ 0.12 wt. %, respectively. All the melts were degassed for 480 s through a rotary degasser supplied with nitrogen inert gas. The chemical composition of the reference alloys, evaluated with Optical Emission Spectrometer (OES) analysis, is reported in Tab. 1. The melts were then transferred to a heated ladle and kept at a temperature of 735 ± 15 °C. After the melt transfer, AlSr15 master alloy rods were added to obtain the targeted strontium content of 100 ppm. The actual strontium contents, measured with the OES, are equal to 103 ppm, 96 ppm and 100 ppm for the alloys Al-7Si, Al-7Si-0.4Mg and Al-7Si-0.4Mg-0.12Ti, respectively.

Tab. 1 – Chemical composition [wt. %] of the reference alloys

Alloy

Al-7Si

6.86

0.01

0.009

0.10

0.0011

0.0014

0.0010

0.0031

Bal.

Al-7Si-0.4Mg

7.16

0.37

0.012

0.11

0.0015

0.0019

0.0020

0.0026

Bal.

Al-7Si-0.4Mg0.12Ti

7.17

0.38

0.123

0.12

0.0016

0.0020

0.0017

0.0029

Bal.

Thermal analysis Thermal parameters of the modified alloys were evaluated by pouring the melts into a pre-heated steel cup (40 mm height, 47 mm upper diameter and 30 mm lower diameter). Cooling curves were recorded for all the samples by means of a mineral insulated K-type thermocouple (1.5 mm diameter) located in the centre of the cup, 15 mm far from the bottom. The same thermocouple was used for all tests: it was placed inside a stainless-steel sheath and then removed from solidified samples. This procedure enables a good comparison of the results obtained for each alloy. Temperature and time data were recorded at a frequency of 20 Hz by a data acquisition system (Pico Technology TC-08 Thermocouple Data Logger) linked to a personal computer, and the acquisition stopped when a temperature of 400 °C was reached during cooling. Cooling curves and their derivatives were processed by means of a tailored Matlab® code. Experimental data processing comprised smoothing, curve fitting and plotting of the first derivative curve for the determination of the characteristic solidifiLa Metallurgia Italiana - n. 2 2018

cation temperatures. The cooling rate (CR) was evaluated in the liquid region prior to the nucleation of α-Al primary phase, in the temperature range 630 - 645 °C. The main solidification parameters of the Al-Si eutectic phase were determined, i.e. minimum temperature Tmin, and growth temperature TG. As displayed in Fig. 1, Tmin is the minimum temperature prior to recalescence, TG is the maximum temperature after Tmin and the recalescence undercooling is defined as ΔTE = TG - Tmin.

Light metals

Fig. 1 – Eutectic solidification region of a cooling curve and its first derivative and related thermal parameters

ΔTG was calculated as the difference between the growth temperature of the unmodified reference alloy (TG,0) and the growth temperature of the alloys modified with 100 ppm of strontium. TG,0 of the unmodified alloy was estimated by Eq. 1, proposed by Apelian et al. [35], that considers the effect

of alloying elements on the eutectic growth temperature. The TG,0 values of the reference alloys, evaluated by Eq. 1 using the chemical compositions reported in Tab. 1, are showed in Tab. 2.

(1)

Tab. 2 – Eutectic growth temperature of the reference alloys calculated by Eq. 1

Alloy

Al-7Si

Al-7Si-0.4Mg

Al-7Si-0.4Mg-0.12Ti

TG,0 [°C]

576.6

573.8

It has been extensively reported in the literature that the depression in eutectic growth temperature is closely linked to the Sr content in the alloy and several studies have investigated the influence of strontium content on ΔTG [8-10, 36]. Image analysis and statistical evaluation Samples from TA were sectioned transversely to the axis of the thermocouple and prepared using standard metallographic procedures. Quantitative Image Analysis (IA) was performed by means of a Leica DMi8 A optical microscope, equipped with Leica Application Suite 4.9 image analysis software. The investigated area was a square of 4 mm2 chosen close to the centre of the sample surface (i.e. close to 8

the tip of the thermocouple) and included enough silicon particles to be representative of the entire sample. Indeed, in accordance with the BS ISO 13322-1:2014, a preliminary study was conducted to determine the minimum number of particles to investigate. The analysis of 400 micrographs enabled to trustworthy identify the distribution of geometrical parameters for the entire population. The minimum number of particles needed was calculated according to Eq. 2:

(2)

La Metallurgia Italiana - n. 2 2018

Leghe leggere where α and β are constants defined from the geometrical parameter to be measured, s is the standard deviation of the population geometrical parameter, δ is the % relative error (fixed to 5 %), u is an intermediate parameter of 1.96 for a given probability of 95 %. The resulting n* of 3000 particles ensures a negligible variability of geometrical features of Si particles. To ensure the suitability of the employed optical magnification and the sufficient number of pixel for the smallest particle to be measured, a minimum threshold of 15 pixels was fixed within the IA software in order to obtain the required accuracy of measurement. In the present study, 5 composite images were considered for each specimen. Each composite image was made up of 25 micrographs, taken at a magnification of 500 ×, and comprised an investigation area of about 1920000 mm2. The described method ensured the analysis of a suitable number of particles, i.e. comprised within the range of 6000 - 10000 particles. Equivalent Diameter (ED = (4A/π 0.5) and Roundness (R = Aπ/4p 2) of eutectic silicon particles were statistically analysed by means of a customised Matlab® code. In particular, distribution fitting was performed on ED and R of eutectic silicon particles and the 3-parameters lognormal distribution resulted to be the best fit to the experimental data, since it was the distribution with the lowest Anderson-Darling statistic (AD) value. Usually, the effect of different chemical compositions on the eutectic microstructure of Al-Si alloys is evaluated by comparing the mean values of some geometrical and morphological parameters of Si particles and considering the standard deviation of data, for example by a Student’s t-test. However, it is worth noting that a simple comparison of means would be poorly efficient with distributions that exhibit a similar central tendency and, in general, it would not be able to catch modification of spread and shape of distribution due to ‘treatment’ effect. For these reasons, comparisons between distributions of silicon particle parameters were evaluated considering the differences of deciles of the distributions themselves, according to the work of Wilcox [30,31]. Deciles were calculated using the Harrell-Davis (HD) quantile estimator [37], which is the weighted average of all the order statistics (Eq. 3):

(3)

where X (i) is the i-th order statistic of the sample and W n,i is the weighting function of the i-th term and it derives from a beta cumulative distribution function. The HD estimator enables, in combination with a bootstrap estimation of the standard error of deciles, to derive the confidence intervals of the difference between deciles of two groups. The bootstrap estimation is a resampling method that relies on random sampling with replacement from an approximating La Metallurgia Italiana - n. 2 2018

distribution. The confidence level for hypothesis testing to evaluate if decile differences are statistically significant was fixed to 95 %. Multiple comparisons were necessary to perform a set of statistical inferences simultaneously, one for each decile. Nevertheless, when performing multiple comparisons, the probability of incorrect rejection increases with increasing number of comparisons. Thus, the method used to evaluate the rate of Type I errors in null hypothesis testing, when conducting multiple comparisons, was the False Discovery Rate (FDR) and it was controlled by the BenjaminiHochberg procedure [38]. This procedure controls the FDR so that the confidence level could be maintained around 5 % across the nine confidence intervals. RESULTS AND DISCUSSION TA curves and parameters Figure 2 compares the cooling curves, restricted to the Al-Si eutectic formation, of the investigated alloys with a strontium content of 100 ppm. The horizontal lines at 576.6 °C and 573.8 °C represent the TG,0 evaluated by Eq. 1, as reported in Tab. 2. It can be observed in Fig. 2 that in the Al-7Si alloy strontium determines a decrease of the eutectic TG of about 2.8 °C. What stands out in Fig. 2 is that the addition of magnesium leads to a significant decrease of both minimum and growth temperatures during the Al-Si eutectic solidification, of about 8 °C. The presence of titanium seems to have a slightly similar effect on the cooling curve since it determines a further slight decrease of Tmin. What can be concluded from these cooling curves is that magnesium and titanium additions appear to not produce an adverse effect on the strontium modification performance since they do not determine a rise of both minimum and growth temperatures of the eutectic solidification. Table 3 reports the principal thermal parameters of the eutectic solidification in the strontium modified investigated alloys. From the data, it can be seen that the depression of TG related to the presence of magnesium and titanium, also depicted in Fig. 2, is reflected in increased values of ΔTG for the Al-7Si-0.4Mg and Al-7Si-0.4Mg-0.12Ti alloys. The Al-7Si alloy modified with 100 ppm of strontium displays a ΔTG of 2.8 °C, whilst the alloys containing magnesium and titanium show a significantly higher ΔT G, of about 7.8 °C. Changes in the recalescence undercooling, defined as ΔTE = TG - T min, also occur, since the alloying elements seem to have an influence also on T min. As illustrated in Tab. 3, for the Al7Si alloy ΔT E has a value of about 3.2 °C but it decreases to 2-2.3 °C in presence of magnesium and of both magnesium and titanium. ΔTG has been often used as an index to be correlated with the modification level [9,10,39]. In fact, it has been reported that modifying agents like sodium and strontium increase twinning density of eutectic silicon and thus allow silicon itself to branch more easily and to form fibrous particles rather than flakes or lamellae [40]. 9

Light metals At the same time, strontium addition seems to deactivate some favourable sites for the nucleation of eutectic phase, typically AlP nuclei, which are located at the tips of primary aluminium dendrites or in the interdendritic space [40-42]. Therefore, eutectic cells require a higher degree of undercooling to nucleate, and this leads to the depression of both nucleation and growth eutectic temperatures [40,42]. These considerations clarify the common use of Î&#x201D;T G as an indicator of strontium modification efficiency. Figures 3a, 3b and 3c depict the microstructures of the three reference alloys modified with 100 ppm of strontium. Comparing the three microstructures, it can be observed that strontium addition to Al-7Si alloy (Fig. 3a) determines a fully modified eutectic morphology, with fine silicon particles. For what concerns the Al-7Si-0.4Mg (Fig. 3b), the simultaneous presence of magnesium and strontium leads to a partially modified eutectic structure, characterised by regions with coarse silicon particles and other regions with silicon particles that show a fine morphology. The titanium presence in the Al-7Si-0.4Mg-0.12Ti (Fig. 3c) does not determine a significant variation of the microstructure already observed in the alloy with sole magnesium. It has been reported that an increase of Î&#x201D;TG is a representative index of the goodness of the strontium modification treatment. For this reason, one could be induced to conclude, from the cooling curves displayed in Fig. 2 and data showed in Tab. 3, that magnesium additions determine a better modification than the one obtained with strontium. Nevertheless, this could be a misled observation, because

it does not take into account the effect of every single alloying element on the cooling curves. In fact, previous studies [11,22,24] reported that the effect of alloying elements on the eutectic temperatures has to be considered when the thermal analysis is used to assess and control the efficiency of a modification treatment. Heusler and Schneider [11] affirmed that the effect of magnesium content is difficult to determine and it cannot be easily correlated with the observed microstructure in an Al-7Si alloy. A similar conclusion is reported by Joenoes and Gruzleski [22], who found that magnesium reduces the degree of microstructure homogeneity despite it determines a decrease of eutectic minimum and growth temperatures. Magnesium alone neither clearly refines nor coarsens eutectic silicon particles in an Al-7Si alloy. Furthermore, magnesium has a negative effect on strontium modification since it changes the eutectic morphology from a fully modified to a partially modified one. For what concerns titanium, Golbahar et al. [24] reported that eutectic growth temperature in a strontium modified A356 alloy is decreased due to titanium additions and thus it does not affect the strontium modification performance. According to the above-reported observations, the effect of alloying elements on both cooling curves and microstructural features cannot be overlooked. In an attempt to further investigate and provide a quantitative evaluation of the role played by alloying elements on microstructure, a detailed analysis of silicon particle geometrical parameters has been carried out.

Fig. 2 â&#x20AC;&#x201C; Cooling curves, restricted to the eutectic solidification region, of the investigated alloys modified with 100 ppm of strontium. The horizontal lines represent the TG,0 values related to each alloy

La Metallurgia Italiana - n. 2 2018

Leghe leggere Tab. 3 – Eutectic thermal parameters [°C] of the alloys modified with 100 ppm of strontium

Alloy

Tmin

ΔTE = TG - Tmin ΔTG = TG,0 - TG

Al-7Si

570.5

573.8

3.3

2.8

Al-7Si-0.4Mg

563.9

566.0

2.1

7.8

Al-7Si-0.4Mg-0.12Ti

563.7

566.0

2.3

7.7

(a)

(b)

(c) Fig. 3 – Comparison of microstructures of the three different alloys with 100 ppm of strontium: a) Al-7Si; b) Al-7Si-0.4Mg; c) Al-7Si-0.4Mg-0.12Ti

La Metallurgia Italiana - n. 2 2018

Light metals Silicon particles Figure 4a represents the 3-parameters log-normal distributions related to the dimensional parameter Equivalent Diameter of eutectic silicon particles, for the investigated alloys, with the addition of 100 ppm of strontium. It can be observed that all distributions are located in a range of very small values of equivalent diameter. In addition, the peak of the curves, that corresponds to the maximum probability density of the selected parameter, is always located in the range between 2 ÷ 3 µm. These observations seem to be in line with the fact that addition of 100 ppm of strontium induces a good degree of modification for all the analysed alloys and thus reduces silicon particle dimension. Nevertheless, the distribution related to the binary Al-7Si alloy clearly shows a higher peak and, consequently, a lower right tail when compared to the other distributions.

(a)

The presence of magnesium leads to a decrease of the maximum value of the distribution and an increase of the right tail. Further, the addition of titanium seems to enhance this effect. These qualitative observations suggest that the presence of both magnesium and titanium induces some modification on the dimension of eutectic silicon particles. Figure 4b illustrates the quantitative comparison between the distributions depicted in Fig. 4a. The trend with solid circles shows the differences between the deciles of the distribution related to Al-7Si-0.4Mg alloys with respect to the deciles of the distribution related to Al-7Si alloy, assumed as the reference. Likewise, the trend with solid squares represents the same differences related to Al-7Si-0.4Mg-0.12Ti alloy. These values are plotted as a function of Al-7Si alloy deciles.

(b) Fig. 4 – a) Distributions of equivalent diameter for the alloys with 100 ppm of Sr; b) Comparison of distributions, Al-7Si taken as reference alloy

It is noteworthy to observe that differences are negative for both the represented curves. This means that all deciles for alloys with sole magnesium and with both magnesium and titanium additions are shifted toward slightly higher values of equivalent diameters, with respect to the binary alloy. Furthermore, the reported trends reveal that each decile difference is statistically significant, as it can be observed in Fig. 4b, since neither of the confidence intervals overlaps the zero line. It is worth to note that the differences of first deciles, although statistically significant, are very close to zero for both the represented trends. Instead, the absolute value of the differences increases to 2 ÷ 3 µm for the last deciles. This increasing trend is directly linked to the modification of the distribution spread and, in particular, it explains the decreased distribution peak and the increased size of the right tail observed with the addition on magnesium and titanium to the binary alloy. The quasizero values of the first decile differences in Fig. 4b reflect the fact that the probability of finding particles with a small size (2 - 2.5 µm, corresponding to first deciles values) in the microstructure is quite the same for the three investigated alloys. This observation is in line with the knowledge reported in the 12

literature that an amount of 100 ppm of strontium is enough to achieve a good degree of modification of silicon particles [8]. On the other hand, the increased value of last deciles suggests that there is a higher probability of finding also particles with a larger size in the alloy containing sole magnesium and even more in the alloy containing both magnesium and titanium, with respect to the binary Al-7Si alloy. As a result, a loss of microstructure homogeneity, depicted in the microstructures of Fig. 3a, 3b and 3c and reported in the literature [22], can be related to a decreased efficiency of the strontium modification treatment. Since strontium modification induces a variation not only on the size of the silicon particles but also on their morphology, a single geometrical parameter is not sufficient to understand its overall effect. Thus, the dimensionless parameter Roundness has been also evaluated. The 3-parameters log-normal distributions for the roundness of the analysed alloys are reported in Fig. 5a. Similarly to Fig. 4a, it shows that the distribution related to the binary Al-7Si alloy appears to be again the one characterised by the maximum value of probability density and the minimum spread, whilst the other distributions exhibit a La Metallurgia Italiana - n. 2 2018

Leghe leggere larger dispersion of data and a taller right tail. Figure 5b reports the comparison between decile differences for the alloy

(a)

with magnesium and titanium addition with respect to the base Al-7Si alloy.

(b) Fig. 5 – a) Distributions of roundness for the alloys with 100 ppm of Sr; b) Comparison of distributions, Al-7Si taken as reference alloy

As previously observed for equivalent diameter, the higher differences are related to the last deciles, whilst for the first deciles differences are close to the zero line. These results suggest that, despite no confidence interval intersects the zero line, there are only slight differences between first deciles. In other words, a large number of particles with low Roundness, i.e. with a shape similar to the circular one, is present in all three alloys. Nevertheless, the addition of magnesium causes the distribution to become flatter and the right tail of the distribution to become taller. Therefore, considering its effect on microstructures showed in Fig. 3a, 3b and 3c, magnesium increases the number of silicon particles which exhibit high roundness values and thus reduces the modification efficiency of strontium, causing the retention of a lamellar silicon structure rather than the formation of a fibrous one. A slightly further deterioration of the eutectic silicon modification level is achieved with the addition of titanium to the alloy with magnesium. The poisoning effect of magnesium on the strontium modification treatment is in good agreement with literature. Some authors [22] reported that magnesium additions to an Al-7Si alloy cause the eutectic silicon morphology to become progressively less modified, as the magnesium content increases. The result is an increased value of geometrical and morphological parameters, i.e. perimeter and aspect ratio respectively, in comparison with the alloy without magnesium. Furthermore, the final microstructure is characterised by a mixture of fibrous, lamellar and acicular silicon particles. This loss of homogeneity of eutectic silicon morphology is linked to higher values of standard deviation for silicon particles descriptors parameters. The same authors suggest that a possible explanation for this poisoning effect is the reaction of magnesium with strontium that leads to the formation of a complex intermetallic compound, with formula Mg2Sr(Si3Al4). The formation of this intermetallic comLa Metallurgia Italiana - n. 2 2018

pound, which occurs at a temperature higher than the eutectic one, reduces the total amount of strontium available for the eutectic silicon modification. In summary, these statistical outcomes are consistent with the observations arisen from the microstructural investigation, showing that magnesium and titanium affect the efficiency of the strontium modification treatment. CONCLUSIVE OBSERVATIONS The present study was carried out with the aim of deepening the interaction between magnesium and titanium, the main alloying elements of Al-7Si alloy, and strontium, one of the most commonly employed modifying agents. According to the experimental findings, the following conclusions can be drawn: • Cooling curves related to the investigated alloys, i.e. Al-7Si, Al-7Si-0.4Mg and Al-7Si-0.4Mg-0.12Ti, modified with 100 ppm of strontium revealed that magnesium determines a substantial decrease of both Tmin and TG of the eutectic solidification. Furthermore, titanium addition also seems to have a slightly similar effect on the eutectic temperatures. • Microstructural observations showed that strontium determines a fine and fibrous morphology of eutectic silicon particles and thus a fully modified microstructure. The presence of magnesium leads to a partially modified eutectic structure and, further, titanium does not determine a significant variation of the microstructure already observed in the alloy with sole magnesium. • Statistical analysis of geometrical parameters of the eutectic silicon particles, i.e. Equivalent Diameter and Roundness, indicated that in both cases data were fitted with a 3-parameters log-normal distribution. The binary Al-7Si alloy shows the higher peak in comparison with the other alloys. The comparison of distributions, performed by quantiles comparison, revealed 13

Light metals that the presence of magnesium leads to a decrease of the maximum values of the distribution and the further addition of titanium seems to enhance this effect. These results show that the presence of both magnesium and titanium induces a variation of the dimension of eutectic silicon particles. In summary, these results could suggest that the Î&#x201D;TG parameter should be used cautiously for the assessment of the modification level when comparing different alloys, since each alloying element could have a certain influence on its numerical value and thus could conduce to a misled interpretation of the efficiency of the modification treatment. Nevertheless, thermal analysis is a valuable on-line instrument to control liquid alloy treatments, forasmuch as microstructural evaluation is operator dependent and requires time-consuming preparation techniques. For this reason, it is important to improve the reliability of thermal analysis technique and it would be helpful to develop

a reference database of thermal parameters, established from a large number of experimental investigations on melts with different alloying elements. In this way, the effect of each alloying element on cooling curves could be then easily tracked and a better correlation between thermal parameters and modification efficiency could be obtained. Acknowledgements: The authors wish to acknowledge Fonderie Mario Mazzucconi Spa for the provision of research facilities at Tekal Spa foundry of San Giovanni Teatino (CH). Special thanks are due to Stefano Pirletti and Stefano Spreafico MorĂ¨ of Fonderie Mario Mazzucconi Spa of Ponte San Pietro (BG) for all the collaborating efforts made during the experimental work.

REFERENCES [1] [2] [3] [4] [5] [6]

[7] [8] [9] [10] [11] [12] [13] [14] [15] [16]

Greer A.L., Cooper P.S., Meredith M.W., Schneider W., Schumacher P., Spittle J.A., Tronche A., Grain Refinement of Aluminium Alloys by Inoculation. 2003, Advanced Engineering Materials, 5, 1-2, 81-91. Quested T.E., Understanding mechanisms of grain refinement of aluminium alloys by inoculation. 2004, Materials Science and Technology, 20, 11, 1357-1369. Darlapudi A., McDonald S.D., StJohn D.H., The influence of Cu, Mg and Ni on the solidification and microstructure of Al-Si alloys. 2016, IOP Conference Series: Materials Science and Engineering, 117, 1, 1-7. Tocci M., Pola A., Raza L., Armellin L., Afeltra U., Optimization of heat treatment parameters for a non-conventional Al-Si-Mg alloy with Cr addition by DoE method. 2016, La Metallurgia Italiana, 6, 141-144. Farahany S., Ourdjini A., Idris M.H., The usage of computer-aided cooling curve thermal analysis to optimise eutectic refiner and modifier in Al-Si alloys. 2012, Journal of Thermal Analysis and Calorimetry, 109, 1, 105-111. Farahany S., Idris M.H., Ourdjini A., Faris F., Ghandvar H., Evaluation of the effect of grain refiners on the solidification characteristics of an Sr-modified ADC12 die-casting alloy by cooling curve thermal analysis. 2015, Journal of Thermal Analysis and Calorimetry, 119, 3, 1593-1601. Xu C., Wang C., Yang H., Liu Z., Yamagata H., Ma C., Solidification Behavior and Mechanical Properties of Al-Si-Mg Alloy with Ti Addition. 2016, Materials science forum, 850, 594-602. Zamani M., Seifeddine S., Determination of optimum Sr level for eutectic Si modification in Al-Si cast alloys using thermal analysis and tensile properties. International Journal of Metalcasting, 2016, 10, 4, 457-465. Shabestari S., Ghodrat S., Assessment of modification and formation of intermetallic compounds in aluminum alloy using thermal analysis. 2007, Materials Science and Engineering A, 467, 1, 150-158. Djurdjevic M., Jiang H., Sokolowski J., On-line prediction of aluminum-silicon eutectic modification level using thermal analysis. 2001, Materials Characterization, 46, 31-38. Heusler L., Schneider W., Influence of alloying elements on the thermal analysis results of Al-Si cast alloys. 2002, Journal of Light Metals, 2, 17-26. Tahiri H., Mohamed S.S., Samuel F.H., Doty H.W., Valtierra S., Effect of Sr-Grain Refining-Si Interactions on the Microstructural Characteristics of Al-Si Hypoeutectic Alloys. 2017, International Journal of Metalcasting, 1-19. Manente A., Timelli G., Influence of strontium modification and solidification history on the microstructure of as cast Al-Si alloys. 2008, La Metallurgia Italiana, 10(10), 1-14. Timelli G., Caliari D., Rakhmonov J., Influence of process parameters and Sr addition on the microstructure and casting defects of LPDC A356 alloy for engine blocks. 2016, Journal of Materials Science and Technology, 32, 515-523. Shabestari S.G., Shahri F., Influence of modification, solidification conditions and heat treatment on the microstructure and mechanical properties of A356 aluminum alloy. 2004, Journal of Materials Science, 39, 2023-2032. Rakhmonov J., Timelli G., Bonollo F., Influence of melt superheat, Sr modifier and Al-5Ti-1B grain refiner on microstructural evolution of secondary Al-Si-Cu alloys. 2016, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science, 47, 5510-5521. La Metallurgia Italiana - n. 2 2018

Leghe leggere [17] Caceres C.H., Davidson C.J., Griffiths J.R., Wang Q.G., The effect of Mg on the microstructure and mechanical behavior of Al-SiMg casting alloys. 1999. Metallurgical and Materials Transactions A, 30, 10, 2611-2618. [18] Yildirim M., Özyürek D., The effects of Mg amount on the microstructure and mechanical properties of Al-Si-Mg alloys. 2013, Materials and Design, 51, 767-774. [19] Sanna F., Fabrizi A., Ferraro S., Timelli G., Ferro P., Bonollo F., Multiscale characterization of AlSi9Cu3(Fe) die casting alloys after Cu, Mg, Zn and Sr addition. 2013, La Metallurgia Italiana, 105(4), 13-24. [20] Li Q., Li B., Li J., Xia T., Lan Y., Guo T., Effects of the addition of Mg on the microstructure and mechanical properties of hypoeutectic Al-7%Si alloy. 2017, International Journal of Metalcasting, 11, 4, 823-830. [21] Aguilera-Luna I., Castro-Román M.J., Escobedo-Bocardo J.C., García-Pastor F.A., Herrera-Trejo M., Effect of cooling rate and Mg content on the Al-Si eutectic for Al-Si-Cu-Mg alloys. 2014, Materials Characterization, 95, 211-218. [22] Joenoes A.T., Gruzleski J.E., Magnesium effects on the microstructure of unmodified and modified Al-Si alloys. 1991, Cast Metals, 4, 2, 62-71. [23] Rakhmonov J., Timelli G., Bonollo F., The influence of AlTi5B1 grain refinement and cooling rate on the formation behavior f Fe-rich compounds in secondary AlSi8Cu3 alloys. 2016, La Metallurgia Italiana, 108(6), 109-112. [24] Golbahar B., Samuel E., Samuel A.M., Doty H.W., Samuel F.H., On thermal analysis, macrostructure and microstructure of grain refined Al-Si-Mg cast alloys: role of Sr addition. 2014, International Journal of Cast Metals Research, 27, 5, 257-266. [25] Alexopoulos N.D., Tiryakioğlu M., Vasilakos A.N., Kourkoulis S.K., The effect of Cu, Ag, Sm and Sr additions on the statistical distributions of Si particles and tensile properties in A357-T6 alloy castings. 2014, Materials Science and Engineering A, 604,40-45. [26] Tiryakioğlu M., Si particle size and aspect ratio distributions in an Al–7%Si–0.6%Mg alloy during solution treatment. 2008, Materials Science and Engineering A, 473,1-6. [27] Samuel A.M., Doty H.W., Valtierra S., Samuel F.H., Effect of grain refining and Sr-modification interactions on the impact toughness of Al-Si-Mg cast alloys. 2014, Materials and Design, 56, 264-273. [28] Ibrahim M.F., Elgallad E.M., Valtierra S., Doty H.W., Samuel F.H., Metallurgical parameters controlling the eutectic silicon characteristics in Be-treated Al-Si-Mg alloys. 2016, Materials, 9, 78. [29] Abdelrahman M.S., Abdu M.T., Khalifa W., Assessment of eutectic modification level in Al-Si alloys via thermal analysis. 2017, Light Metals, 885-895. [30] Wilcox R.R., Comparing two independent groups via multiple quantiles. 1995, The Statistician, 44, 1, 91-99. [31] Wilcox R.R., Erceg-Hurn D., Clark F., Carlson M., Comparing two independent groups via the lower and the upper quantiles. 2014, Journal of Statistical Computation and Simulation, 84, 7, 1543-1551. [32] Niklas A., Abaunza U., Fernández-Calvo A.I., Lacaze J., Suárez R., Thermal analysis as a microstructure prediction tool for A356 aluminium parts solidified under various cooling conditions. 2011, In: 69th World Foundry Congress (WFC), 16-20 Oct 2010, Hangzhou, China. [33] Malekan M., Dayani D., Mir A., Thermal analysis study on the simultaneous grain refinement and modification of 380.3 aluminum alloy. 2014, Journal of Thermal Analysis and Calorimetry, 115, 1, 393-399. [34] Moustafa M.A., Lepage C., Samuel F.H., Doty H.W., Metallographic observations on phase precipitation in strontium-modified Al-11.7% Si alloys: Role of alloying elements. 2003, International Journal of Cast Metals Research, 15, 6, 609-626. [35] Apelian D., Sigworth G.K. , Whaler K.R., Assessment of grain refinement and modification of Al-Si foundry alloys by thermal analysis. 1984, AFS Transactions, 297-307. [36] McDonald S.D., Nogita K., Dahle A.K., Eutectic nucleation in Al-Si alloys. 2004, Acta Materialia, 52, 14, 4273-4280. [37] Harrell F.E., A new distribution-free quantile estimator. 1982, Biometrika, 69, 3, 635-640. [38] Benjamini Y., Hochberg Y., Controlling the false discovery rate: a practical and powerful approach to multiple testing. 1995, Journal of the Royal Statistical Society B, 57, 1, 289-300. [39] Chen X., Geng H., LI Y., Assessment of modification level of hypoeutectic Al -Si alloys by pattern recognition of cooling curves. 2005, China Foundry, 2, 4, 246-253. [40] Nogita K., McDonald S.D., Dahle A.K., Eutectic solidification and its role in casting porosity formation. 2004, Journal of Minerals, Metals and Materials Society, 52-58. [41] Laslz G., Dual microstructure in hypoeutectic Al-Si alloys: dendrites and eutectic grains. Its effects on shrinkage behaviour. 1995, Proceedings of the 4th International Conference on Molten Aluminum Processing, 459-480. [42] McDonald S.D., Nogita K., Dahle A.K., Eutectic grain size and strontium concentration in hypoeutectic aluminium-silicon alloys. 2006, Journal of Alloys and Compounds, 422, 184-191.

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Light metals Theoretical and Experimental Evaluation of the Effectiveness of Aluminum Melt Treatment by Physical Methods V.B. Deev, E.S. Prusov, A.I. Kutsenko Physical methods of aluminum melts treatment have a number of advantages in comparison with the traditional methods of grain refinement: they do not require the use of expensive additives, they do not influence the change in the alloy chemical composition, and they are more environmentally favorable. Comparative evaluation of various physical methods of aluminum melt processing according to the calculated and experimental values of the solid fraction close to the solidus temperature and the total solidification time on the example of the A356 alloy has been carried out in this paper. It has been shown that when melt is processed by physical methods, the size of the dendritic cell and the structure of the eutectic are significantly reduced. Fluidity of test alloys has increased after the treatment by 18 ... 25% by an average. The most significant increase in the strength characteristics of the A356 alloy has been noted after the combined thermo-temporal treatment of melts and the application of electric current during crystallization (+ 20.5%), as well as thermo-temporal treatment followed by the inert gas purging (+ 20%), herewith more than two-fold increase in ductility has been indicated in each of the above-noted treatment modes.

KEYWORDS: CAST ALUMINUM ALLOYS, GRAIN REFINEMENT, PHYSICAL MELT TREATMENT, FRACTION SOLID, TOTAL SOLIDIFICATION TIME, FLUIDITY, MICROSTRUCTURE, MECHANICAL PROPERTIES INTRODUCTION Grain refining treatment is one of the effective ways to control the structure and properties of aluminum alloys. Traditional widely used grain refining additives (titanium, boron, strontium, sodium, phosphorus, etc.) can significantly improve the quality of castings from both hypoeutectic and hypereutectic alloys [1-3]. However, the search for new grain refiners and the development of more efficient methods for grain refinement of aluminum alloys continue to be of interest to metallurgists and casters [4-6]. More and more attention recently has been paid to the development and investigation of methods that make it possible to obtain a refined alloy structure without introducing special additives into the melt. One of the promising ways of ensuring the quality of castings and reducing the material and energy costs for their production is the working out and mastering of advanced methods of melting and casting using physical effects on liquid and crystallizing melt. Such physical methods as vibrational [7], ultrasonic [8-10], electric [11-13], electromagnetic [14-19], electron beam [20, 21] etc. have been successfully tested for grain refining treatment of aluminum melts so far. At the same time different variants of high-temperature furnace treatment (thermo-temporal treatment, superheat treatment, etc.) are used to homogenize the melt and to refine the grain structure of the castings [22-25]). However, for the industrial production of shaped castings, these technologies are still used to a limited extent due to the lack of knowledge of 16

the processes and the lack of the information on the optimal processing regimes for various alloys and technological processes of melting and casting. Despite the fact that some of the above mentioned methods require the use of complex and expensive equipment, as well as the strict control of the temperature and time modes of melting, they contribute to the increase in the number of crystallization centers and grain structure refinement in comparison with the

V.B. Deev Wuhan Textile University, Wuhan, China; National University of Science & Technology (MISIS), Moscow, Russia

E.S. Prusov Vladimir State University named after Alexander and Nikolay Stoletovs, Vladimir, Russia

A.I. Kutsenko Siberian State Industrial University, Novokuznetsk, Russia

La Metallurgia Italiana - n. 2 2018

Leghe leggere methods of chemical grain refinement where the effect of their application strongly depends on the uniform distribution of the additive in the melt is limited in time in some cases, and the grain refining elements themselves can influence the change in the chemical composition of the alloy [26]. In addition, the methods of physical impact on the melt allow the use of up to 100% scrap and waste products in the furnace charge in some cases [27, 28]. The best results from the point of view of obtaining cast aluminum alloys and castings of a given quality, especially with an increased amount of scrap and waste in the charge, can certainly be achieved with the complex application of various physical methods of the grain refinement, for example, by complex melt processing using thermo-temporal treatment and vibration, or thermo-temporal treatment and electric current, etc. However, in industrial casting technologies, insufficient attention is paid to the complex methods of treating melts by physical action. Before the beginning of the melt treatment by physical methods,

it is desirable to analyze various variants of its processing and to choose the most suitable method for given castings production conditions. Within this framework, the actual task is the elaboration and experimental proof of a reliable computational and analytical technique that makes it possible to assess the effectiveness of various physical methods of melt processing. The purpose of this paper is a comparative evaluation of various methods of aluminum melt processing by superimposing physical actions on the results of theoretical calculations and experimental data. MATERIALS AND METHODS A hypoeutectic aluminum alloy A356 (equivalent to the designation EN AB-42100 according to EN 1676-2010 specifications) has been used as a research subject. Chemical composition of alloy is given in Table. 1.

Tab. 1 – Chemical composition of the A356 alloy (EN AB-42100), wt.%

bal.

6.5-7.5

0.300.45

≤0.15

≤0.03

≤0.10

≤0.07

≤0.18

The charge has consisted of 10 ... 15% of pig alloys and 85 ... 90% of recycled materials (scrap and waste of A356 alloy). Experimental meltings have been carried out in the alumina crucible in the electric resistance furnace. The alloy has been subjected

Others each

total

≤0.03

≤0.10

to various physical actions during melting and molding for the comparative evaluation of their effectiveness. The conditions of experimental meltings and melt treatment parameters are given in Table 2.

Tab. 2 – Technological modes for melt treatment by physical actions

N°

Method and modes of melt treatment

Initial melt (untreated)

Thermo-temporal treatment : T = 1000 0C, τ = 8…12 min

Vibration in vertical direction: amplitude 1…1,2 mm; frequency 50 Hz

Thermo-temporal treatment (T = 1000 0C, τ = 8…12 min) + vibration (amplitude 1…1.2 mm, frequency 50 Hz)

Direct electric current with density j = 0,92 . 105 A/m2 at solidification

Thermo-temporal treatment (T = 1000 0C, τ = 8…12 min) + direct electric current with density j = 0.92 . 105 A/m2 at solidification

Argon: purging time 5…6 min at pressure 0.3 MPa

Thermo-temporal treatment (T = 1000 0C, τ = 8…12 min) + argon: purging time 5…6 min at pressure 0,3 MPa

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Light metals Thermo-temporal treatment modes have been the same for all melt processing variations: superheating to 1000 °C, holding for 8 ... 10 min in the inert gas atmosphere (argon), cooling to the treatment temperature by the second physical action in the combined processes (750 °C) or to the pouring temperature (740 °C). In all cases, the melt have been degassed with manganese chloride (0.2%) before pouring. During crystallization processing of the molding with a direct electric current of j = 0.92 · 10^5 A/ m^2 current density and I = 30 A current strength using a standard current source delivered to the sand form has been carried out. Argon blowing has been carried out by means of a lance injected into the melt; the blowing time has been 5 ... 6 min at a pressure of 0.3 MPa. To create vibration (1 ... 1.2 mm amplitude, 50 Hz frequency), a special vibrating table whose vibratory impulse transferred to a casting form through a metal platform in a vertical axial direction has been used.

To determine the theoretical amount of the solid phase falling out near the solidus temperature during the melt crystallization, and the total time of its crystallization, a special calculation technique has been developed. This technique is based on the determination of the alloys crystallization parameters related to fluidity. Fluidity of casting alloys is the most important property determining the production of castings of a desired quality, and depends on the amount of solid phase falling out near the solidus temperature. At a critical amount of the solid phase (from 25 to 30% for various alloys), the melt ceases to flow. Therefore, knowledge of the solid phase amount, as well as the total crystallization time of the alloy τn is essential in the castings production. The fraction of the solid phase mOM, falling out near the solidus temperature, has been calculated by the derived formula:

(1)

where ΛM, Λ0 is the fluidity of the modified (treated) and unmodified (untreated) alloy, respectively; ΔT is the overheating interval above the liquidus temperature of the untreated alloy; m0 is the solid fraction at crystallization of the untreated alloy (m0 = 0.30 has been assumed for the calculations); ΔTLS is the interval of the untreated alloy crystallization; Lf is the latent heat of the melt crystallization; c is the melt heat capacity. The formula (1) for calculating the mass fraction of the solid phase mOM for a modified alloy at a known value of m0 for the initial

alloy was derived by mathematical transformations of the conventional heat balance equations for the heat transfer processes between melt and mold [29], classical dependences of the theory of fluidity of metallic melts [30], and also taking into account the lever rule for the Al-Si phase diagram under the assumption that the melt crystallizes under equilibrium conditions [31]. The calculated value of the casting total crystallization time has been determined by the Chvorinov’s rule [32]:

(2)

where λm is the thermal conductivity of the mold; ρm is the density of the mold material; cm is the mold material heat capacity; ρs is the density of the solid metal; Lf is the latent heat of crystallization; Tf is the melting point of the metal; T0 is the initial temperature of the mold; A is the form and metal separation surface area; V is the volume of the casting. The experimental determination of the solid fraction mOM and the alloy total crystallization time τn has been carried out using direct thermal and differential thermal analysis methods at the solidification of cylindrical samples of 300 mm of length and 26

mm of diameter in the sand form. The experimental technique is suitable when performing computer thermal analysis of melts crystallization using Computer Aided Cooling Curve Analysis (CA-CCA) methods [33]. To determine of the solid phase volume fraction in the solidifying sample by the CA-CCA method, the following mathematical model has been used. At each timepoint the solid phase volume fraction in the crystallization interval of the alloy can be calculated according to [34] by the following equation:

(3)

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Leghe leggere where V0 and VS are the volume of the melt sample and the volume of the solid phase, respectively; tL is the timepoint when the melt reaches the liquidus temperature; T is the temperature of the melt sample; TL and TS are the liquidus and solidus temperatures of the alloy, respectively; T(t) is the sample tempera-

ture as a function of time; T0 is the ambient temperature; α is the coefficient of heat transfer; tS is the end of the alloy solidification; k(T) is a coefficient that can be determined empirically for the sample state, when φ(tS)=1 is a fully solidified sample or more precisely calculated according to equation

(4) where F is the surface area of the sample; ρ is the density of the alloy; Lf is the latent solidification heat; α is the heat tran-

sfer coefficient which is depends on the sample temperature and can be calculated by following equation:

(5)

To determine the temperatures of the beginning and the end of solidification of alloys, the curves of the differential thermal analysis of the crystallizing sample have been used (Fig. 1). The solidus temperature has been calculated on the basis of the analysis of the first derivative of obtained thermogram T’(t), liquidus – based on the analysis of the second derivative of thermogram T’’(t). The wavelet transform of the signal based on the normalized second derivative of Gaussian has been used to determine the liquidus temperature at high noise interference of temperature measurements. The alloys fluidity has been determined from a horizontal round rod sample of 5 mm in diameter and 450 mm in length, poured into the steel mold in a gravity casting conditions (commercially available test method for measuring fluidity; one of its examples described by F. Binczyk et al. [35]). The mold temperature before pouring was kept constant at 200 °C (measured by a calibrated K-type thermo-couple installed at the mold center). Pouring temperature for all variants has

been 725 ... 730 °C. The mold was coated with a zinc oxide to the thickness about 0.3 mm. To ensure reproducible hydraulic conditions of the mold filling in all experiments, the head height was the same and equal to 100 mm. To realize this, a stopper rod was installed in the pouring cup, which was raised after complete filling the cup with a melt. The number of test samples for calculating the mean value of fluidity for each variant of melt processing technology was 10. The samples mi-crostructure has been examined in as-cast state by a Carl Zeiss Axiovert 200M inverted microscope. Mechanical properties of the alloys have been determined on standard samples with a working diameter of 10 mm on the universal testing machine Electronic Universal Testing Machine WDW100E. When studying the properties of the experimental alloys in the course of the direct measurements, the arithmetic average and confidence intervals of the obtained values of the measured quantity have been calculated for each series of experiments.

Fig. 1 – Typical thermogram of A356 alloy

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Light metals RESULTS AND DISCUSSION The calculated and experimental values of the total solidification time and the solid fraction as exemplified in A356 alloy treated by different physical methods are shown in Fig. 2 and 3. The analysis of the obtained data indicates that physical actions affect the melt crystallization parameters, increasing the total solidification time and the solid fraction falling out near the solidus temperature. It has been assumed that external physical effects on melts contribute to a decrease of the se-

parating diffusion rate near the liquidus temperature and a shift of the crystallization process to the lower temperatures area so the fraction of the solid phase formed near the solidus temperature increases with respect to the base alloy for the treated alloys. After the treatment by physical action the melt is able to retain fluidity at a relatively high fraction of the formed solid phase. In melt treatment the most effective are variants 2 and 8, in which mOM corresponds to 0.41 and 0.42 according to the calculations, 0.38 and 0.39 according to the experimental data.

Fig. 2 – Calculated and experimental values of the total solidification time τn of the alloy sample, s (numbers of melt treatment variations correspond to those indicated in Table 2).

Fig. 3 – Calculated and experimental values of the solid phase fraction mOM (numbers of melt treatment variations correspond to those indicated in Table 2).

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Leghe leggere Fig. 5 shows the optical micrographs of the A356 alloy structure in the initial state (variant 1) and after the melt treatment (variants 2 and 8). It can be seen that during treatment there is a substantial reduction of both the size of the dendritic cell

and of the eutectic structure. These observations testify to the complex effect of the applied treatment methods and the effectiveness of their application for refining the grain structure of the alloys of the Al-Si system.

c Fig. 4 – Microstructure of the experimental alloys: a – initial alloy (untreated); b – melt processing technology according to variant 2 (Table 2); c – melt processing technology according to variant 8 (Table 2).

Results of the investigations of mechanical properties (ultimate tensile strength UTS and elongation δ) and the alloy fluidity (Λ)

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at using of different physical methods of melt treatment are presented in Table. 3.

Light metals Tab. 3 - Mechanical and technological properties of A356 alloy (as-cast condition) depending on the melt processing method

Variant 1 2

Method of melt treatment Initial alloy (untreated) Thermo-temporal treatment

UTS, MPa

δ, %

Λ, mm

190 ± 3

2.8 ± 0.3

96 ± 2.80

200 ± 5

3.5 ± 0.1

109 ± 2.72

Vibration

211 ± 7

4.2 ± 0.1

109 ± 3.26

Thermo-temporal treatment + vibration

219 ± 5

4.7 ± 0.3

115 ± 3.68

Electric current

222 ± 3

5.2 ± 0.1

Thermo-temporal treatment + electric current

229 ± 4

5.9 ± 0.2

Inert gas purging

223 ± 5

5.8 ± 0.1

120 ± 2.88

Thermo-temporal treatment + inert gas purging

228 ± 2

6.1 ± 0.3

126 ± 3.17

The obtained experimental data confirmed the positive effect of complex treatment by physical methods. Fluidity of the experimental alloys, determined by reference to the test sample, increased by an average by 18 ... 25%, herewith the best results have been achieved for variant No. 8 (+ 31.3%). The most significant increase in the strength characteristics of the A356 alloy has been recorded after the combined thermo-temporal treatment and application of electric current during crystallization (+ 20.5%) as well as thermo-temporal treatment and inert gas purging (+ 20%), with that two-fold ductility increase in each of the indicated treatment variants has been mentioned. Thus, with the employment of the complex technologies for melts treating, the use of the preliminary thermo-temporal treatment enhances the grain refining effect of the second physical action greatly. This turned out to be characteristic for all melt processing options. The best results have shown the options of the combined use of thermo-temporal treatment with an inert gas (No. 8), or with electric current (No. 6). The stated increase of the alloy mechanical properties are conditioned by the refining of its structural constituents. Treatment by electric current during crystallization is promising in the production of castings from secondary aluminum alloys containing iron impurity. The positive effect of such treatment can be expressed in a change in the dispersability and morphology of iron-containing phases. However, the mechanism of the electric current impact on the process of alloys crystallization has not yet been sufficiently investigated both in the experimental and in the theoretical aspects. The possibility of using thermo-temporal treatment in the production of castings from 22

low-grade charge materials with elevated iron content should also be considered [36]. A promising direction for further research is the investigation of the physical actions influence the formation of cast aluminum matrix composites structure and properties [37-39]. It has been assumed that the application of external physical action during melting and casting will allow to control the composites structure purposefully, providing a desired degree of interphase interaction between the matrix material and the reinforcing phase, as well as the uniform distribution of reinforcing components in the melt volume. CONCLUSIONS It has been shown that processes of casting aluminum alloys crystallization can be determined to a great extent by the technology of processing by physical influences during melting and casting. These technologies, provided they are used rationally, can significantly improve the quality of the obtained cast products. The technique for estimating the effectiveness of various physical actions on melts based on the determination of the solid phase fraction falling out near the solidus temperature during the alloys crystallization has been proposed and tested. The predictive calculations by this technique are in good agreement with the experimental data of thermal and differential thermal analysis. The technique is recommended for selecting the most rational and effective physical methods for melts processing on the predicted values of the crystallization parameters of aluminum alloys. La Metallurgia Italiana - n. 2 2018

Leghe leggere Despite the fact that the use of physical methods for melts processing in industrial conditions leads to some increase in the production cost of castings due to additional power energy costs, this increase is compensated due to the absence the need to use expensive refining additives to provide the desired structure and properties of castings with increased operational requirements. The obtained data on the changes in the structure and properties of aluminum alloys indicate that the physical methods of melt treatment (thermo-temporal treatment, electric current during crystallization, etc.) can be successfully applied in foundries in the production of castings for critical duty appli-

cations, for which strict limitations on the content of impurity elements have been established. Physical methods make it possible to produce high quality castings with a modified structure without using refining additives influencing the chemical composition of the alloy. ACKNOWLEDGMENTS The work has been carried out within the framework of the state work "Organization of Scientific Research" of the Ministry of Education and Science of Russia of the state task in the field of research for 2017-2019 (Task No. 11.5684.2017 / 6.7).

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

SIGWORTH G.K., KUHN T.A. Grain refinement of aluminum casting alloys // International Journal of Metalcasting (2007), 1: 31-40. LI P., LIU S., ZHANG L., LIU X. Grain refinement of A356 alloy by Al–Ti–B–C master alloy and its effect on mechanical properties // Materials & Design (2013), 7: 522-528. MEENA P.C., SHARMA A., SINGH S. Effect of grain refinement on microstructure and wear behavior of cast Al-7Si alloys // La Metallurgia Italiana (2015), 1: 25-34. WANNASIN J., CANYOOK R., WISUTMETHANGOON S., FLEMINGS M.C. Grain refinement behavior of an aluminum alloy by inoculation and dynamic nucleation // Acta Materialia (2013), 61: 3897-3903. EASTON M.A., QIAN M., PRASAD A., ST. JOHN D.H. Recent advances in grain refinement of light metals and alloys // Current Opinion in Solid State and Materials Science (2016), 20: 13-24. GUAN R.G., TIE D. A Review on grain refinement of aluminum alloys: progresses, challenges and prospects // Acta Metallurgica Sinica (2017), 30: 409-432. VOROZHTSOV S., KUDRYASHOVA O., PROMAKHOV V., DAMMER V., VOROZHTSOV A. Theoretical and experimental investigations of the process of vibration treatment of liquid metals containing nanoparticles // JOM (2016), 68: 3094–3100. ESKIN D.G. Ultrasonic processing of molten and solidifying aluminium alloys: overview and outlook // Materials Science and Technology (2017), 33: 636-645. ZHANG Y., SVYNARENKO K., LI T. Effect of ultrasonic treatment on formation of iron-containing intermetallic compounds in Al-Si alloys // China Foundry (2016), 13: 316–321. S. KOMAROV, Y. ISHIWATA, I. MIKHAILOV. Industrial application of ultrasonic vibrations to improve the structure of Al-Si hypereutectic alloys: Potential and Limitations // Metallurgical and Materials Transactions A (2015), 46: 2876-2883. RABIGER D., ZHANG Y., GALINDO V., FRANKE S., WILLERS B., ECKERT S. The relevance of melt convection to grain refinement in Al-Si alloys solidified under the impact of electric currents // Acta Materialia (2014), 79: 327-338. PRODHAN A. Semi-solid processing by electric current during sand casting of aluminium alloys // IOP Conf. Series: Materials Science and Engineering (2016), 115: 012005. Y. ZHANG, D. RABIGER, B. WILLERS, S. ECKERT. The effect of pulsed electrical currents on the formation of macrosegregation in solidifying Al-Si hypoeutectic phases // International Journal of Cast Metals Research (2017), 30: 13-19. ZHANG, Y., CHENG, X., ZHONG, H., (...), SONG, C., ZHAI, Q. Comparative study on the grain refinement of Al-Si alloy solidified under the impact of pulsed electric current and travelling magnetic field // Metals (2016), 6, 170. BUSTOS O., ORDOÑEZ S., COLÁS R. Rheological and microstructural study of A356 alloy solidified under magnetic stirring // International Journal of Metalcasting (2013), 7: 29-37. WANG X.., LUO X., CONG F., CUI J. Research progress of microstructure control for aluminium solidification process // Chinese Science Bulletin (2013), 58: 468–473. CHEN H., JIE J., FU Y., MA H., LI T. Grain refinement of pure aluminum by direct current pulsed magnetic field and inoculation // Transactions of Nonferrous Metals Society of China (2014), 24: 1295–1300. HAGHAYEGHI R., DE PAULA L.C., ZOQUI E.J. Comparison of Si refinement efficiency of electromagnetic stirring and ultrasonic treatment for a hypereutectic Al-Si alloy // Journal of Materials Engineering and Performance (2017), 26: 1900–1907. TIMOSHKIN I.Y., NIKITIN K.V., NIKITIN V.I., DEEV V.B. // Influence of treatment of melts by electromagnetic acoustic fields on

La Metallurgia Italiana - n. 2 2018

Light metals [20]

[21] [22] [23] [24]

[25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39]

the structure and properties of alloys of the Al–Si system // Russian Journal of Non-Ferrous Metals (2016), 57: 419-423. IVANOV Y.F., ALSARAEVA K.V., GROMOV V.E., POPOVA N.A., KONOVALOV S.V. Fatigue life of silumin treated with a highintensity pulsed electron beam // Journal of Surface Investigation. X-ray, Synchrotron and Neutron Techniques (2015), 9: 1056–1059. IVANOV Y.F., ALSARAEVA K.V., GROMOV V.E., KONOVALOV S.V., SEMINA O.A. Evolution of Al–19.4Si alloy surface structure after electron beam treatment and high cycle fatigue // Materials Science and Technology (2015), 31: 1523–1529. TIMELLI, G., FIORESE, E. Metodi di neutralizzazione del Fe in leghe Al-Si da fonderia // La Metallurgia Italiana (2011), 103: 9-23. LI Q.L., XIA T.D., LAN Y.F., LI P.F. Effects of melt superheat treatment on microstructure and wear behaviours of hypereutectic Al-20Si alloy // Materials Science and Technology (2014), 30: 835-841. PENG J., JINYANG Z., HAORAN G., ZHONGXI Y., XINYING T., DEGANG Z., YAN W., MIN Z., NINGQIANG S. Effect of melt superheating treatment on solidification structures of Al75Bi9Sn16 immiscible alloy // Journal of Molecular Liquids (2017), 232: 457-461. DEEV V.B., SELYANIN I.F., PONOMAREVA K.V., YUDIN A.S., TSETSORINA S.A. Fast cooling of aluminum alloys in casting with a gasifying core // Steel in Translation (2014), 44: 253-254. SELYANIN I. F., DEEV V. B., BELOV N. A., PRIKHODKO O. G. Physical modifying effects and their influence on the crystallization of casting alloys // Russian Journal of Non-Ferrous Metals (2015), 56: 434-436. DEEV V.B., SELYANIN I.F., KUTSENKO A.I., BELOV N.A., PONOMAREVA K.V. Promising Resource Saving Technology for Processing Melts During Production of Cast Aluminum Alloys // Metallurgist (2015), 58: 1123-1127. SELYANIN I.F., DEEV V.B., KUKHARENKO A.V. Resource-saving and environment-saving production technologies of secondary aluminum alloys // Russian Journal of Non-Ferrous Metals (2015), 56: 272-276. FREDRIKSSON H., AKERLIND U. Solidification and Crystallization Processing in Metals and Alloys. John Wiley & Sons, Ltd. (2012). CAMPBELL J. Casting, 2nd ed. Oxford: Elsevier, Butterworth-Heinemann (2003). STEFANESCU D.M. Science and Engineering of Casting Solidification, 3rd ed. Springer International Publishing Switzerland (2015). DANTZIG J.A., RAPPAZ M. Solidification. Taylor & Francis, Lausanne (2009). SUDHEER R., PRABHU K.N. A Computer Aided Cooling Curve Analysis method to study phase change materials for thermal energy storage applications // Materials & Design (2016), 95: 198-203. RAFALSKI I., ARABEY A., LUSHCHIK P., CHAUS A.S. Computer modeling of cast alloys solidification by Computer-Aided Cooling Curve Analysis (CA-CCA) // International Doctoral Seminar, Proceedings. Trnava: Alumni Press, 2009: 291-301. BINCZYK F., CIESLA M., P. GRADON, R. FINDZINSKI. Evaluation of casting shrinkage and liquid metal fluidity of IN-713C Alloy // Archives of Foundry Engineering (2014), 14: 9-12. YANG W., YANG X., JI S. Melt superheating on the microstructure and mechanical properties of diecast Al-Mg-Si-Mn alloy // Metals and Materials International (2015), 21: 382–390. TIMELLI, G., FERRO, P., BONOLLO, F. Compositi a matrice di alluminio solidificati in presenza di vibrazioni meccaniche: Caratteristiche microstrutturali // La Metallurgia Italiana (2010) 102. PRUSOV E.S., PANFILOV A.A. Properties of cast aluminum-based composite alloys reinforced by endogenous and exogenous phases // Metally (2011), 7: 670-674. NORDIN N., ABUBAKAR T., HAMZAH E., FARAHANY S., OURDJINI A. Effect of Superheating Melt Treatment on Mg2Si Particulate Reinforced in Al-Mg2Si-Cu In situ Composite // Procedia Engineering (2017), 184: 595–603.

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Leghe leggere Analisi del comportamento a corrosione di campioni di alluminio AA6012 sottoposti a ECAP e trattamento criogenico A. Viceré, M. Cabibbo, C. Paoletti, G. Roventi, T. Bellezze In questo lavoro è stato studiato il comportamento a corrosione della lega di alluminio AA6012 (Al-Mg-Si-Pb), sottoposta dapprima a un trattamento termico di solubilizzazione (550°C/6h) e successivamente a una deformazione plastica severa, ottenuta attraverso la tecnica dell’Equal Channel Angular Pressing (ECAP). Alcuni campioni delle lega sono stati sottoposti a trattamento criogenico prima dell’ECAP. La presenza di piombo nella lega AA6012 contribuisce in modo significativo alle prerogative di lavorabilità a freddo della lega. L’analisi della resistenza a corrosione dei vari campioni di AA6012 è stata eseguita a temperatura ambiente mediante una caratterizzazione elettrochimica in soluzioni acquose aventi la stessa concentrazione di cloruri (0,1M NaCl), ma a diverso pH (pH 2 e pH 6,5). A questo scopo, sono state effettuate prove elettrochimiche di resistenza di polarizzazione, prove di polarizzazione potenziodinamica e misure di impedenza elettrochimica. L’analisi della microstruttura della lega AA6012 è stata effettuata attraverso osservazioni al microscopio ottico e al TEM. Dai risultati sperimentali è emerso che il trattamento termico di solubilizzazione migliora il comportamento a corrosione dei campioni tal quali, mentre il successivo passaggio ECAP lo peggiora. Infine il trattamento criogenico eseguito prima dell’ECAP fa recuperare, almeno in parte, la resistenza a corrosione persa in seguito alle deformazioni plastiche severe.

PAROLE CHIAVE: LEGHE DI ALLUMINIO, ECAP, TRATTAMENTO CRIOGENICO, RESISTENZA DI POLARIZZAZIONE, POLARIZZAZIONE POTENZIODINAMICA, IMPEDENZA ELETTROCHIMICA INTRODUZIONE Le leghe di alluminio sono utilizzate in diverse applicazioni tecnologiche e industriali, grazie alla loro leggerezza, alle loro proprietà fisiche e meccaniche e al loro basso costo. Negli ultimi anni, i metalli ultrafini, ossia con un grano cristallino al di sotto del micron, hanno destato molto interesse perché presentano un’elevata resistenza meccanica abbinata a una buona tenacità, oltre a proprietà superplastiche a temperature moderate e a velocità di deformazione elevate [1-6]. Attraverso tecniche di deformazione plastica severa è possibile ottenere l’affinamento dei grani nei metalli. Queste tecniche sono in grado di impartire ad un campione di metallo sufficientemente duttile deformazioni plastiche, senza significativi cambiamenti della sua geometria macroscopica. La tecnica Equal Channel Angular Pressing (ECAP) in questo ambito è la più diffusa perché è possibile processare billette anche di elevate dimensioni che possono assumere struttura ultrafine in tutto il loro volume. Tramite questa tecnica, un provino è pressato mediante un punzone e forzato a passare attraverso uno stampo contenente due canali a sezione costante che si intersecano. Il campione mantiene la stessa sezione trasversale dopo la pressatura e così è possibile ripetere il passaggio diverse volte [7].

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Ci sono diversi studi in merito all’effetto della temperatura in fase di esecuzione dell’ECAP sull’evoluzione microstrutturale delle leghe di alluminio [8-15]. La maggior parte di questi lavori sono stati comunque eseguiti a temperatura ambiente o a elevate temperature. D’altra parte, la temperatura criogenica si suppone sopprima il recupero dinamico dei grani durante il processo ECAP [10-12], consentendo probabilmente il mantenimento di

A. Viceré, G. Roventi, T. Bellezze Dipartimento di Scienze e Ingegneria della Materia, dell’Ambiente ed Urbanistica (SIMAU), Università Politecnica delle Marche, Via Brecce Bianche, 12, 60131 Ancona. t.bellezze@univpm.it; tel. 071 2204413

M. Cabibbo, C. Paoletti Dipartimento di Ingegneria Industriale e Scienze Matematiche (DIISM), Università Politecnica delle Marche, Via Brecce Bianche, 12, 60131 Ancona

Light metals un’alta densità di difetti, che possono agire come potenziali siti di ricristallizzazione per la formazione di strutture ultrafini. I processi che portano all’affinamento dei grani di un metallo alterano anche la superficie del materiale, conducendo a cambiamenti nella densità dei bordi di grano, nel loro orientamento e nelle tensioni residue. Di conseguenza, questi cambiamenti possono avere un impatto sul comportamento elettrochimico e quindi sulla suscettibilità alla corrosione di questi materiali. In letteratura c’è un numero relativamente limitato di lavori su come l’affinamento della dimensione dei grani influenzi la resistenza a corrosione di una lega; inoltre, la letteratura esistente è spesso contradditoria, persino quella relativa alla stessa lega [16]. La resistenza a corrosione del materiale metallico avente la microstruttura affinata dall’ECAP, che conduce a un aumento della densità dei bordi di grano attivi, è fortemente legata alla particolare combinazione tra ambiente e materiale [16-17]. Infatti, un ambiente che rende attivo il materiale aumenta la corrosione dei campioni deformati plasticamente tramite ECAP, mentre un ambiente che lo rende passivo riduce la corrosione degli stessi, come riportato da diversi autori [18-20]. La lega AA6012 è stata scelta considerando il suo elevato rapporto resistenza meccanica-peso, la sua spiccata lavorabilità alle macchine utensili dopo estrusione, buona formabilità anche a freddo, buona saldabilità e resistenza a corrosione. Tenendo conto che la deformazione per ECAP è assimilabile, dal punto di vista tecnologico, ad una deformazione plastica per estrusione, si è preferito iniziare questo lavoro di ricerca con una lega di alluminio facilmente estrudibile a freddo, grazie in particolare al fatto che contiene piombo. L’aggiunta di questo elemento, nelle leghe della serie 6000, è attualmente previsto e consentito dalle normative ambientali nell’ambito della Comunità Europea (CE). Attualmente la AA6012 risulta in produzione in diverse nazioni della CE con diverse classificazioni, come, EN AW 6012 (classificazione Europea), P-Al-SiMgMn, in Italia, DIN-1700 AlMgSiPb, in Germania, o A-SGPb, in Francia. In questo lavoro è stata studiata la resistenza a corrosione di provini di lega di alluminio AA6012 (Al-Mg-Si-Pb) sottoposti a deformazione plastica severa tramite la tecnica ECAP, a temperatura ambiente e a temperatura criogenica. La caratterizzazione elettrochimica dei campioni ottenuti con questo metodo di deformazione è stata effettuata attraverso misure potenziodinamiche e di impedenza elettrochimica, in soluzione di cloruri, a temperatura ambiente, sia a pH neutro che a pH acido. I risultati ottenuti hanno permesso di confrontare le prestazioni dei diversi campioni. PARTE SPERIMENTALE La composizione chimica della lega AA6012 studiata è la seguente (% in peso): 0,8 Si, 0,5 Mn, 1,0 Mg, 0,8 Pb, 1.0max (Fe+Cu+Cr+Zn+Ti). Alcune barre cilindriche di questa lega (100 mm di lunghezza e sezione pari a 10 mm) sono state pressate mediante un punzone e forzate a passare attraverso uno stampo ECAP con una forza compresa nell’intervallo 40-80 kN e con una velocità di 100 mm/min. Lo stampo ECAP utilizzato è costituito da un blocco di acciaio per utensili SK3 (Fe-1,1% C) contenen26

te due canali, con sezione circolare di diametro pari a 10 mm, che si intersecano formando un angolo di 90°. I provini sono stati deformati plasticamente a temperatura ambiente fino a 1 passaggio ECAP. I provini sottoposti a trattamento criogenico (cryo) sono stati immersi in azoto liquido per almeno un minuto immediatamente prima di essere processati attraverso lo stampo ECAP. Durante la deformazione di taglio, le billette sottoposte a trattamento cryo non hanno mai raggiunto la temperatura ambiente. I provini esaminati in questo lavoro sono: (i) billette estruse così come fornite dal produttore (tal quali, TQ); (ii) billette sottoposte al solo trattamento termico di solubilizzazione a 550°C/6h (TT); (iii) billette processate una sola volta attraverso l’ECAP dopo la solubilizzazione (TT-ECAP 1); (iv) billette sottoposte in sequenza alla solubilizzazione, al trattamento cryo e infine al passaggio ECAP (TT-cryo ECAP 1). La microstruttura dei provini TT, TT-ECAP 1 e TT-cryo ECAP 1 è stata osservata al microscopio ottico (MO), utilizzando luce polarizzata per meglio identificare la struttura dei grani della lega nelle diverse condizioni sperimentali. Le superfici dei tre provini osservati al MO sono state lucidate meccanicamente e successivamente attaccate elettroliticamente mediante una soluzione Nital al 2% per circa 60 s, a una tensione di 12V. Le analisi quantitative della dimensione media dei grani è stata eseguita mediante un software di analisi di immagini Image Proplus®. I campioni per il TEM sono stati preparati mediante doppio getto con una soluzione di 1/3 HNO3 in 2/3 di alcol etilico, con una tensione di 20 kV e una temperatura di -35°C. Campioni cilindrici di diametro Φ= 1,00-1,14 cm e altezza pari a 1 cm sono stati ottenuti dalle barre e sono stati mantenuti in congelatore (T= -18°C) per impedire l’invecchiamento naturale. Prima di ogni test di corrosione, essi sono stati estratti dal congelatore; una volta raggiunta la temperatura ambiente, i campioni sono stati puliti con carte abrasive fino a 4000 grit, sciacquati con acqua distillata e asciugati con aria calda. La caratterizzazione elettrochimica di questi campioni è stata fatta sulla loro sezione trasversale dopo averli inglobati in una resina epossidica e dopo aver realizzato un opportuno contatto elettrico, che ha permesso di utilizzarli come elettrodi di lavoro. La superficie esposta alle soluzioni corrosive era di 0,50-0,70 cm2. Per le prove di resistenza a corrosione, è stata utilizzata una cella elettrochimica a tre elettrodi, che, oltre all’elettrodo di lavoro, era costituita da un elettrodo di riferimento a calomelano saturo (SCE, +0,241 V vs NHE) e un controelettrodo di platino, tutti connessi al potenziostato Gamry Reference 600. Le soluzioni di prova utilizzate sono state due: una soluzione deareata di NaCl 0,1 M portata a pH 2 con l’aggiunta di HCl concentrato (soluzione 1) e una soluzione areata di NaCl 0,1 M a pH 6,5 (soluzione 2). L’assenza di ossigeno per la soluzione 1 è stata garantita facendo gorgogliare N2 per un’ora con la soluzione mantenuta in agitazione. Trascorso questo tempo, l’agitatore è stato fermato e sono state effettuate le prove elettrochimiche. La resistenza a corrosione dei campioni è stata studiata tramite le misure di resistenza di polarizzazione (Rp) e le curve di Tafel nel caso dei campioni immersi nella soluzione 1 e tramite misure di spettroscopia di impedenza elettrochimica e polarizzazioni La Metallurgia Italiana - n. 2 2018

Leghe leggere anodiche e catodiche nel caso dei campioni immersi nella soluzione 2. Nella soluzione 1 sono state effettuate due scansioni potenziodinamiche (0,166 mV/s): da Ecorr -5 mV a Ecorr +5 mV per le misure di Rp e da Ecorr -150 mV a Ecorr +150 mV per ottenere le curve di Tafel. Queste ultime sono state analizzate con il metodo analitico sviluppato precedentemente, ottenendo così i parametri cinetici utili alla determinazione della densità di corrente di corrosione icorr [21,22]. Nella soluzione 2, le polarizzazioni anodiche e catodiche sono state effettuate separatamente ad una velocità di scansione di 0,166 mV/s, partendo dal potenziale di corrosione, utilizzato come base per le misure di impedenza. Queste ultime sono state effettuate con un’ampiezza del segnale AC di 5 mV rms, 5 punti/decade in un intervallo di frequenza 100 kHz - 10 mHz. I dati ottenuti con le misure di impedenza sono stati analizzati con il software Gamry Echem Analyst. Tutte le prove elettrochimiche sono state condotte a temperatura ambiente dopo aver lasciato stabilizzare Ecorr per almeno 30 minuti dall’immersione di ciascun campione in entrambe le soluzioni utilizzate. I test eseguiti sono stati ripetuti almeno tre volte. RISULTATI E DISCUSSIONE La Fig. 1 riporta le microstrutture rappresentative, acquisite al MO, dei provini TT, TT-ECAP 1 e TT-cryo ECAP 1. La valutazione delle dimensioni medie dei grani è stata fatta su tre piani ortogonali X-Y-Z e i risultati sono espressi in μm3. Il motivo di questa scelta risiede nell’aspetto morfologico dei grani stessi che sono profondamente deformati plasticamente a seguito del processo ECAP, soprattutto lungo i piani Y e Z, mentre le micrografie qui riportate si riferiscono al solo piano X (ortogonale alla direzione di estrusione), ovvero il piano su cui sono state eseguite le prove di resistenza a corrosione. I riquadri nelle Fig. 1b e 1c mostrano la microstruttura dei campioni TT-ECAP 1 e TT-cryo ECAP 1 acquisita nel piano Y, allo stesso ingrandimento relativo al piano X. La dimensione media dei grani, osservata nella condizione TT (Fig. 1a), risulta essere di (XxYxZ): 210x240x275 μm3, con una incertezza del dato sperimentale del 2%. Il confronto diretto tra le due condizioni relative alla lega deformata plasticamente, ovvero TT-ECAP 1 (Fig. 1b) e TT-cryo ECAP 1 (Fig. 1c), mostra chiaramente la natura della deformazione plastica che risulta essere prevalentemente lungo una direzione di circa 45° rispetto a quella di estrusione ECAP. Le dimensioni medie dei grani sono risultate in entrambi i casi sostanzialmente simili: per i provini TT-ECAP 1 e TT-cryo ECAP 1 sono rispettivamente pari a

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110x25x21 μm3 e pari a 105x38x24 μm3. Una differenza significativa tra le due condizioni TT-ECAP 1 e TT-cryo ECAP 1 consiste nella formazione di fasi indurenti β-Mg2Si, osservata nel provino sottoposto ad ECAP (condizione TT-ECAP 1), e nell’inibizione di tale precipitazione nel caso in cui venga eseguito un trattamento criogenico prima della deformazione ECAP (condizione TT-cryo ECAP 1). Questo interessante aspetto microstrutturale è stato documentato mediante osservazioni TEM [23]. Più in dettaglio, le Fig. 1d e 1e riportano due micrografie TEM rappresentative rispettivamente delle condizioni TT-ECAP 1 e TT-cryo ECAP 1. Esse mostrano chiaramente la presenza di particelle nanometriche di fase β nella condizione TT-ECAP 1, mentre queste particelle nanometriche risultano essere assenti nella condizione TT-cryo ECAP 1 [23]. La Fig. 2 mostra le curve di Tafel rappresentative di ogni tipologia di campione di lega AA6012, ottenute nella soluzione 1 deaerata (pH = 2). Dalle curve sono state determinate le costanti di Tafel, i valori di icorr e i potenziali di corrosione Ecorr. Le coppie di valori (log icorr, Ecorr), corrispondenti a ogni curva di Fig. 2, sono indicati nel grafico tramite opportuni simboli colorati. Sulla base dell’andamento dei tratti anodici di Fig. 2, si può osservare che tutti i campioni analizzati della lega AA6012 hanno mostrato una corrosione attiva. In realtà, la curva relativa al campione TT mostra una non trascurabile inflessione del ramo anodico che indica una minore tendenza di questo campione a ossidarsi all’aumentare del potenziale, durante la scansione potenziodinamica. È bene osservare che i valori di densità di corrente registrati in corrispondenza del flesso (circa 10-4,5 A/cm2, Fig. 2) non corrispondono ai valori caratteristici dei materiali passivi (10-5 – 10-7 A/cm2) [24-26]. Per quanto riguarda i tratti catodici delle curve, essi corrispondono alla riduzione dell’idrogeno visto che la loro pendenza (bc) si approssima al valore teorico di 120 mV/decade. Nel dettaglio, i campioni che hanno mostrato una maggiore velocità di corrosione (icorr più elevato) sono il TQ e il TT-ECAP 1, mentre i campioni TT e TT-cryo ECAP 1, al contrario, sono quelli caratterizzati da velocità di corrosione più basse (v. simboli in Fig. 2). In sostanza, il trattamento criogenico prima dell’ECAP 1 (campione TT-cryo ECAP 1) risulta benefico per la lega di alluminio esaminata perché riduce la perdita di resistenza a corrosione dovuta alla successiva deformazione plastica severa (campione TT-ECAP 1), eseguita dopo il trattamento termico (campione TT).

Light metals

Fig. 1 - Immagini MO relative alla microstruttura nel piano X dei provini TT (550Â°C/6h), a), TT-ECAP 1, b), e TT-cryo ECAP 1, c). Nei riquadri sono illustrate le microstrutture nel piano Y. Micrografie TEM (piano X) della lega sottoposta a TT-ECAP 1, d), e a TT-cryo ECAP 1, e). - Optical microscopy images of the microstructures corresponding to the X plane of the samples TT (550 Â°C/6h), a), TT-ECAP 1, b), and TT-cryo ECAP 1, c). Inset: microstructure relative to Y plane. TEM micrographs (X plane) of the alloy subjected to TT-ECAP 1, d), and to TT-cryo ECAP 1, e)

Fig. 2 - Curve di Tafel rappresentative dei campioni analizzati in soluzione 1 deaerata: ( TAL QUALE; TT; TT-ECAP 1; TT-cryo ECAP 1); i simboli rappresentano per ogni curva dello stesso colore i corrispondenti valori (log icorr, Ecorr) - Representative Tafel curves of the samples analyzed in deaerated solution 1: ( AS RECEIVED; TT; TT-ECAP 1; TT-cryo ECAP 1); the symbols represents the (log icorr, Ecorr) values for each corresponding curve, having the same colour 28

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Leghe leggere La Fig. 3 mostra i valori medi di Rp ottenuti dai test di polarizzazione e i corrispondenti valori medi di icorr determinati dalle curve di Tafel per tutti i campioni analizzati. In corrispondenza di tali valori medi, nella stessa figura sono riportate le barre di errore che mostrano una buona ripetibilità dei test effettuati e una buona significatività delle differenze di resistenza a corrosione osservate. In generale, il processo di solubilizzazione a 550°C/6h (campioni TT) migliora la resistenza a corrosione dei campioni TQ della lega AA6012; infatti Rp aumenta da 680 a 1310 Ω cm2 (in corrispondenza, icorr diminuisce da 35 a 28 μA cm-2). Questa variazione significativa di comportamento può essere attribuita alla riduzione delle tensioni residue nel materiale metallico [16]. Le deformazioni plastiche severe prodotte dall’ECAP dopo la solubilizzazione diminuiscono la resistenza a corrosione della lega AA6012 (Fig. 2 e 3); infatti, Rp diminuisce e icorr aumenta dopo il passaggio ECAP (Fig. 3). Questo risultato sperimentale può essere attribuito ad un affinamento dei grani del materiale (Fig. 1b e 1c), che, producendo un aumento della densità dei bordi di grano, determina una maggiore presenza sulla

superficie del campione di siti attivi che aumentano la velocità di corrosione [16], in soluzioni acide come la soluzione 1. Un altro fattore che influisce negativamente sulla resistenza a corrosione della lega in esame sottoposta al passaggio ECAP è la formazione di seconde fasi come β-Mg2Si osservate nei campioni TT-ECAP 1 (Fig. 1d) [23]. In accordo con altri autori [27, 28], questa specifica fase β è elettrochimicamente più attiva della matrice della lega e di conseguenza è soggetta a corrodersi in maniera preferenziale, diminuendo così in generale la resistenza a corrosione del materiale. Il recupero della resistenza a corrosione dei campioni TT-cryo ECAP 1 al livello di quelli TT, e comunque la loro resistenza a corrosione maggiore rispetto ai campioni TT-ECAP 1 (Fig. 3), può essere spiegato dall’assenza delle seconde fasi (Fig. 1e), come la β, determinata proprio dal trattamento criogenico che ha preceduto la deformazione plastica imposta durante il processo ECAP [23]. La Fig. 4 mostra le curve di polarizzazione di alcuni campioni rappresentativi tra tutti quelli studiati, ottenute nella soluzione 2 a pH 6,5.

Fig. 3 - Resistenza alla polarizzazione (Rp) e densità di corrente di corrosione (icorr) dei differenti campioni di AA6012 nella soluzione 1. Le frecce tratteggiate indicano la variazione di queste grandezze nel passaggio dai provini solubilizzati a quelli sottoposti a ECAP - Polarization resistance (Rp) and corrosion current density (icorr) of the different samples of AA6012 in solution 1. The dotted arrows indicate the variation of these characteristic corrosion parameters from the solutioned samples to ECAPed ones

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Light metals

Fig. 4 - Curve potenziodinamiche di polarizzazione rappresentative dei campioni analizzati nella soluzione 2: ( TQ; TT; ECAP 1; TT-cryo ECAP 1) - Representative potentiodynamic polarization curves of the samples analyzed in solution 2: ( RECEVEID; TT; TT-ECAP 1; TT-cryo ECAP 1) Per la lega AA6012, una concentrazione di cloruri pari a 0,1 M, in ambiente quasi neutro, è sufficiente per determinare un fenomeno di pitting diffuso su tutta la superficie del campione. Bisogna però osservare che i campioni TT mostrano nel tratto anodico un flesso prima dell’aumento della densità di corrente (Fig. 4) che porta comunque alla corrosione localizzata diffusa. Il ramo catodico delle curve di Fig. 4 mostra per tutti i campioni testati il tratto caratteristico del logaritmo della densità di corrente limite di scarica dell’ossigeno, che sembra pertanto essere il processo che governa la corrosione nella soluzione 2, determinando così una scarsa possibilità di differenziarli. Bisogna però aggiungere che nel caso dei campioni TT è stata ottenuta una densità di corrente limite più bassa rispetto a quella registrata per gli altri campioni; questo fenomeno è attribuibile alla parziale formazione di un film di passivazione o di uno strato protettivo di specie adsorbite sulla superficie, come indicato anche dal leggero comportamento passivo [24] che questi campioni hanno mostrato nel tratto anodico. Data la scarsa differenziazione del comportamento a corrosione dei vari campioni testati nella soluzione 2, sono state

TTAS

effettuate misure di impedenza elettrochimica, in grado di fornire più informazioni sul sistema in studio. Il fatto che il processo di corrosione dei campioni testati sia governato dal processo di diffusione dell’ossigeno è confermato dalla loro risposta in frequenza, come mostra per esempio il diagramma di Nyquist di Fig. 5, nel quale a basse frequenze si nota un andamento rettilineo con pendenza all’incirca di 45° [29]. Questo diagramma è relativo ad un campione TT, ma tutti gli altri campioni hanno mostrato una risposta analoga. Il fitting dei dati mostrati in Fig. 5 è stato eseguito mediante il software Gamry Echem Analyst, considerando il circuito equivalente mostrato in Fig. 6 in cui: Rs rappresenta la resistenza della soluzione tra elettrodo di lavoro e elettrodo di riferimento; Q dl è l’elemento a fase costante associato con la capacità del doppio strato; Rt è la resistenza di trasferimento di carica del processo di corrosione; W è l’impedenza di Warburg associata al processo di diffusione dell’ossigeno verso l’elettrodo di lavoro. La Fig. 7 mostra i risultati relativi alla resistenza al trasferimento di carica per i vari campioni esaminati con le relative barre di errore.

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Leghe leggere

Fig. 5 - Diagramma di Nyquist relativo al campione TT testato nella soluzione 2 â&#x20AC;&#x201C; Nyquist diagram obtained on a sample of TT examined in solution 2

Fig. 6 - Circuito equivalente utilizzato per il fitting dei dati sperimentali mostrati in Fig. 5 â&#x20AC;&#x201C; Equivalent circuit used for fitting the experimental data of Fig. 5

Fig. 7 - Resistenza al trasferimento di carica (Rt) dei differenti campioni di AA6012 in soluzione 2. La freccia tratteggiata indica la variazione di questa grandezza nel passaggio dai provini solubilizzati a quelli sottoposti a ECAP - Charge transfer resistance (Rt) of the different samples of AA6012 in solution 2. The dotted arrow indicates the variation of this characteristic corrosion parameter from the solutioned samples to ECAPed ones La Metallurgia Italiana - n. 2 2018

Light metals Nonostante la maggiore dispersione dei risultati, ottenuta nelle misure di Rt a pH 6,5 (Fig. 7) rispetto a quella ottenuta nelle misure di Rp a pH 2 (Fig. 3), Rt dei campioni TT assume il valore più alto, pari a 4164 Ω cm2. Inoltre, in questa soluzione, i campioni sottoposti ad ECAP hanno mostrato una resistenza al trasferimento di carica (2734 Ω cm2) più bassa rispetto ai campioni TT e comunque più alta rispetto ai TQ (1558 Ω cm2). Infine, il campione sottoposto prima al trattamento criogenico e poi a ECAP ha fatto registrare un nuovo aumento della Rt (3652 Ω cm2). Quindi anche questi risultati ottenuti in ambiente neutro confermano quelli ottenuti per gli stessi campioni nella soluzione 1 avente pH acido e pari contenuto di cloruri: il trattamento termico di solubilizzazione aumenta la Rt dei campioni TQ, mentre le deformazioni provocate dall’ECAP riducono questa resistenza, che però, nel caso della soluzione 2, rimane superiore a quella dei campioni TQ. Inoltre, il trattamento cryo prima dell’ECAP permette un recupero, anche se non in maniera totale, della Rt dei campioni sottoposti solo ad ECAP, facendo riferimento in particolare ai valori elevati di resistenza relativi ai campioni TT. CONCLUSIONI Questo lavoro è consistito nell’analisi del comportamento a

corrosione di campioni di lega di alluminio AA6012 sottoposti a deformazioni plastiche severe attraverso la tecnica ECAP fino a 1 passaggio, dopo trattamento termico di solubilizzazione ed eventuale trattamento criogenico. È stata eseguita una caratterizzazione elettrochimica mediante misure di resistenza di polarizzazione, polarizzazioni anodiche e catodiche e misure elettrochimiche di impedenza, in soluzioni con lo stesso contenuto di cloruri ma a differenti pH. Dai risultati sperimentali ottenuti è emerso che in ambiente acido il trattamento termico di solubilizzazione migliora il comportamento a corrosione dei campioni di lega AA6012 rispetto a quelli senza trattamento, mentre il processo ECAP diminuisce la resistenza a corrosione acquisita con la solubilizzazione. D’altra parte effettuando il trattamento criogenico prima della deformazione plastica severa, è stato osservato un miglioramento della resistenza a corrosione rispetto ai campioni sottoposti solamente a ECAP. I risultati ottenuti in ambiente neutro confermano quasi completamente i risultati ottenuti in ambiente acido: il trattamento ECAP risulta peggiorativo nei confronti dei campioni solubilizzati, mentre il trattamento criogenico eseguito prima dell’ECAP permette di evitare la perdita della resistenza a corrosione dovuta alla deformazione plastica severa.

BIBLIOGRAFIA [1] R.Z. VALIEV, R.K. ISLAMGALIEV, I.V. ALEXANDROV, Progress in Materials Science 45 (2000) 103. [2] M. FURUKAWA, Y. IWAHASHI, Z. HORITA, M. NEMOTO, N.K. TSENEV, R.Z. VALIEV, T.G. LANGDON, Acta Materialia 45 (1997) 4751. [3] R.Z. VALIEV, Materials Science and Engineering A 59 (1997) 234. [4] W.J. KIM, J.K. KIM, T.Y. PARK, S.I. HONG, D.I. KIM, J.D. LEE, Metallurgical and Materials Transactions A 33 (2002) 3155. [5] Z. HORITA, T. FUJINAMI, M. NEMOTO, T.G. LANGDON, Journal of Materials Research and Technology 117 (2001) 288. [6] C.S. CHUNG, J.K. KIM, H.K. KIM, W.J. KIM, Materials Science and Engineering A 337 (2002) 39. [7] M. VEDANI, G. ANGELLA, P. BASSANI, A. TUISSI, Metallurgia Italiana, 5 (2006). [8] A. GOLOBORODKO, O. SITDIKOV, R. KAIBYSHEV, H. MIURA, T. SAKAI, Materials Science and Engineering A 381 (2004) 121. [9] A. YAMASHITA, D. YAMAGUCHI, Z. HORITA, T.G. LANGDON, Materials Science and Engineering A 287 (2000) 100. [10] Y.M. WANG, E. MA, Applied Physics Letters 83 (2003) 3165. [11] Y.M. WANG, M.W. CHEN, H.W. SHEN, E. MA, Journal of Materials Research 17 (2002) 3004. [12] Y.B. CHEN, Y.L. LI, L.Z. HE, C. LU, H. DING, Q.Y. LI, Materials Letters 62 (2008) 2821. [13] W.H. HUANG, C.Y. YU, P.W. KAO, C.P. CHANG, Materials Science and Engineering A 366 (2004) 221. [14] F.S. SUN, P. ROJAS, A. ZUNIGA, E.J. LAVERNIA, Materials Science and Engineering A 430 (2006) 90. [15] F. ZHOU, X.Z. LIAO, Y.T. ZHU, S. DALLEK, E.J. LAVERNIA, Acta Materialia 51 (2003) 2777. [16] K.D. RALSTON, N. BIRBILIS, Corrosion 66 (2010) 075005-1. [17] K.D. RALSTON, D. FABIJANIC, N. BIRBILIS, Electrochimica Acta 56 (2011) 1729. [18] M. HOCKAUF, L.W. MEYER, D. NICKEL, G. ALISCH, T. LAMPKE, B. WIELAGE, L. KRUEGER, Journal of Material Science 43 (2008) 7409. [19] M.-K. CHUNG, Y.-S. CHOI, J.-G. KIM, Y.-M. KIM, J.-C. LEE, Materials Science and Engineering A 366 (2004) 282. [20] D. SONG, A.-B. MA, J.-H. JIANG, P.-H. LIN, D.-H. YANG, Transactions of Nonferrous Metals Society of China 19 (2009) 1065. [21] T. BELLEZZE, G. GIULIANI, G. ROVENTI, Corrosion Science 130 (2018) 113. [22] T. BELLEZZE, G. GIULIANI, A. VICERÉ, G. ROVENTI, Corrosion Science 130 (2018) 12. [23] M. CABIBBO, E. SANTECCHIA, P. MENGUCCI, T. BELLEZZE, A. VICERE’, Materials Science and Engineering A 716 (2018) 107. [24] K. EL-MENSHAWY, A.-W. A. EL-SAYED, M. E. EL-BEDAWY, H. A. AHMED, S. M. EL-RAGHY, Corrosion Science 54 (2012) 167.

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Leghe leggere [25] [26] [27] [28] [29]

T. BELLEZZE, G. ROVENTI, R. FRATESI, Corrosion Engineering Science and Technology 48 (2013) 340. T. BELLEZZE, G. GIULIANI, G. ROVENTI, R. FRATESI, F. ANDREATTA, L. FEDRIZZI, Materials and Corrosion 67 (2016) 831. L. TAN, T. R. ALLEN, Corrosion Science 52 (2010) 548. K. A. YASAKAU, M. L. ZHELUDKEVICH, S. V. LAMAKA, M. G. S. FERREIRA, Electrochimica Acta 52 (2007) 7651. C. GABRIELLI, Technical Report Number 004/83, Solartron, UK 1984.

Analysis of corrosion behaviour of aluminium alloy AA6012 samples processed by ECAP and cryogenic treatment KEYWORDS: ALUMINUM ALLOYS, ECAP, CRYOGENIC TREATMENT, POLARIZATION RESISTANCE, POTENTIODYNAMIC POLARIZATION, ELECTROCHEMICAL IMPEDANCE SPECTROSCOPY

Severe Plastic Deformations (SPD) have been used in order to obtain ultrafine-grained alloys improved in strength and wear resistance with an adequate ductility. Equal Channel Angular Pressing (ECAP) is one of these techniques that consists in pressing through a die, with two intersecting channels, equal in cross section, a sample that is forced to deform by shearing at the intersection of these channels. The sample retains the same cross-sectional area after pressing, so that it is possible to repeat the pressing several times. Modifications of the alloys microstructure produced by ECAP significantly influence their corrosion behaviour. In this work, the corrosion behaviour of the aluminium alloy AA6012 samples solutioned at 550°C/6h and then pressed through ECAP technique was investigated; some of these samples were submitted to a cryogenic treatment before being processed by ECAP. With more details, the samples examined in this work are as received extruded billets (TQ), solutioned at 550 °C/6h (TT), processed only one pass through ECAP (TT-ECAP 1) and submitted to cryo treatment before ECAP (TT-cryo ECAP 1). The samples TT-ECAP 1 were pressed into the ECAP die at room temperature until one pass. The cryogenic treated billets (TT-cryo ECAP 1) were dipped into liquid nitrogen for at least 1 minute immediately before being introduced into the ECAP die. During shear deformation, the cryotreated billets never reached room temperature. The evolution of microstructure was studied by means of optical microscope and TEM observations. The analysis of corrosion behaviour has been carried out at room temperature by means of electrochemical characterization in two different aqueous solutions at the same chloride concentrations (0.1M NaCl): a deareated solution at pH= 2 (solution 1) and an aerated solution at pH= 6.5 (solution 2). For this purpose, polarization resistance, potentiodynamic polarization and electrochemical impedance spectroscopy tests were performed. A three-electrode cell was used with a saturated calomel electrode (SCE, +0.241 V vs NHE) as a reference and two short-circuited platinum sheets as counter, all connected to a Gamry Reference 600 potentiostat. Polarization resistance measurements were carried out from Ecorr – 5 mV to Ecorr + 5 mV, Tafel plot from Ecorr -150 mV to Ecorr +150 mV and potentiodynamic polarizations, anodic and cathodic separated tests, from the corrosion potential (scan rate 0.166 mV/s). Electrochemical Impedance Spectroscopy (EIS) measurements were performed with an AC signal amplitude of 5 mV rms, 5 points/decade and a frequency range of 100 kHz - 10 mHz. Fig. 1 shows grain refinement of the AA6012 samples processed by ECAP (TT-ECAP 1 and TT-cryo ECAP 1) with respect to TT samples. Fig. 2 illustrates representative Tafel curves for each sample examined in deareated solution 1 (pH=2). All the samples showed active corrosion. TT and TT-cryo ECAP 1 samples present lower corrosion rate (icorr) than TQ and TT-ECAP 1 (symbols in Fig. 2). Polarization resistance results (Fig. 3) confirm the better corrosion behaviour of TT sample and the improvement in corrosion resistance of TT-cryo ECAP 1 samples with respect to TT-ECAP 1 ones in solution 1 (pH=2). Corrosion behavior in solution 2 (pH= 6.5) is mainly governed by oxygen diffusion and only limited differences are visible in Fig. 4 between the analyzed samples. Therefore, EIS measurements were carried out in order to obtain more information about the system under examination, which typically presented the impedance response shown in Fig. 5. Through the analysis of this response by the equivalent circuit of Fig. 6, the charge transfer resistance has been determined for all samples (Fig. 7). These results confirm the better corrosion behaviour of TT samples, the worsening of ECAPed samples and the improvement of the corrosion resistance given by cryogenic treatment before ECAP. In conclusion, solutioned samples (TT) show a better corrosion behaviour than as received (TQ) and ECAPed ones. Although severe plastic deformations worsen the corrosion behaviour of solutioned samples, cryogenic treatment before ECAP avoids the loss of the corrosion resistance of AA6012 alloy due to ECAP. La Metallurgia Italiana - n. 2 2018

Light metals Electrochemical corrosion behaviour of binary magnesium - heavy rare earth alloys F. Rosalbino, S. De Negri, G. Scavino, A. Saccone The corrosion properties of magnesium-heavy rare earth (RE) based alloys have been studied. Binary additions of gadolinium (Gd), dysprosium (Dy) and erbium (Er) to pure magnesium were made to a nominal 1 at.%. The corrosion resistance of Mg99Gd1, Mg99Dy1 and Mg99Er1 alloys has been assessed by using open circuit potential measurements, potentiodynamic polarization curves and electrochemical impedance spectroscopy (EIS) carried out in 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4. Electrochemical results showed that heavy RE alloying additions significantly improves the corrosion behaviour of magnesium. This improvement can be attributed to enhanced barrier properties of the corrosion products layer and additional active corrosion protection originated from the inhibiting action of the lanthanide cations entrapped as oxides/hydroxides in this surface layer.

KEYWORDS: MAGNESIUM; HEAVY RARE EARTH; CORROSION; POLARIZATION; ELECTROCHEMICAL IMPEDANCE SPECTROSCOPY INTRODUCTION Magnesium alloys constitute a very interesting alternative to the materials traditionally applied in structural applications, and, especially in the automotive and aircraft industry. The main advantage of magnesium alloys is the reduction of the weight of the components due to the low density and high specific strength of these materials [1, 2]. However, the principal drawback of magnesium alloys is the low corrosion resistance, which is generally, much lower when compared to many other competing materials, like aluminum alloys or steels [2]. This limits the range of technical applications of these materials. The stability of several metals and alloys depends upon the formation of stable surface films. However, in magnesium and its alloys the surface film that forms spontaneously is poorly protective and very unstable in a wide range of pH values. This film becomes protective and stable only at pH values over 11 [3]. Literature reporting the electrochemical behavior of pure magnesium [4 - 7], zirconium free [8 - 14] and zirconium containing alloys [15 - 19], has been published in the last decade. Improvements to the corrosion resistance of Mg alloys have been correlated with the addition of alloying elements such as aluminum [8, 11 - 13], zirconium [16 - 18] and yttrium [19]. Addition of rare earths (RE) as alloying elements to the solid phase is an effective way to enhance the corrosion behavior of magnesium alloys. The scavenger effect, optimized microstructure and formation of more protective corrosion 34

product films were considered as the main key factors to improve the corrosion resistance of RE-containing Mg alloys surface and to the inhibition of further corrosion [20 - 23]. However, it has also been suggested that the oxidation of the RE, followed by an enrichment of these elements in the surface film, decreases the corrosion rate by forming a pseudo passive layer [24]. Moreover, It has been reported [11] that the oxides formed can be more compact and able to trap harmful anions by making the surface charge more positive. The objective of this work is to contribute to a better under-

F. Rosalbino, G. Scavino Dipartimento di Scienza Applicata e Tecnologia, Politecnico di Torino, Corso Duca degli Abruzzi, 24 – 10129 Torino (Italy)

S. De Negri, A. Saccone Università degli Studi di Genova, Dipartimento di Chimica e Chimica Industriale, Via Dodecaneso 31 – 16146 Genova (Italy)

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Leghe leggere standing of the corrosion mechanism of magnesium alloys and especially the role of heavy rare earths in buffer neutral/ alkaline solution. The electrochemical corrosion behaviour of binary Mg-RE alloys containing gadolinium (Gd), dysprosium (Dy) and erbium (Er), was assessed and compared with that of unalloyed magnesium in order to gain information about the influence of the heavy RE on corrosion process. The corrosion behaviour of the RE-containing Mg alloys was investigated by using open circuit potential measurements, potentiodynamic polarization curves and electrochemical impedance spectroscopy (EIS) carried out in borate buffer solution. This electrolyte allows for an insight into the role of the RE in the corrosion mechanism by imposing a pH at which Mg can cover itself with scarcely protective oxide or hydroxide which checks the dissolution reaction [3]. Furthermore, the buffering effect avoids pH fluctuations, which in NaCl solutions, for example, may increase more than 4 units. Experimental Binary alloys with nominal composition Mg99RE1 (at.%, RE = Gd, Dy, Er) were prepared by direct synthesis from the constituent metals (purity > 99.9 mass %, Mg was supplied by MaTecK, Jülich,Germany, the rare earth metals by Newmet Koch, Waltham Abbey, England). Stoichiometric amounts of the elements were enclosed in arc-sealed Ta crucibles, in order to prevent Mg losses due to evaporation, and induction melted under argon flow. Melting was repeated three times in order to ensure homogeneity. Ingots with a diameter of ~0.8 cm and a mass of ~0.4 g were extracted from the crucibles after quenching in cold water. Chemical composition of samples was determined by spectrometric analysis employing an argon microwave plasma torch coupled to spark ablation (Spectro Analytical Instruments). Specimens were ablated by a medium voltage spark (450 V, 370 Hz) in a point-to-plane configuration (spark times: 125 s) and swept into a 100-W, 2.45-GHz argon microwave discharge. The microwave plasma was observed end-on and the radiation analyzed with a polychromator. Microstructure examination was performed by a scanning electron microscope (SEM) Zeiss Evo 40 equipped with a Pentafet Link Energy Dispersive X-ray Spectroscopy (EDXS) system managed by the INCA Energy software (Oxford Instruments, Analytical Ltd., Bucks, U.K.). Smooth surfaces for microscopic observation before the electrochemical tests were prepared by using SiC papers and diamond pastes with grain size down to 1 μm. All the electrochemical tests were carried out in a single compartment cell using a standard three electrode configuration: saturated calomel electrode (SCE) as a reference with a platinum electrode as counter and a sample as the working electrode. The surface area exposed to the test solution was 0.5 cm2. Before each measurement, the samples surface was first ground with 320, 400 and 600 μm SiC abrasive papers in anhydrous ethyl alcohol and then La Metallurgia Italiana - n. 2 2018

automatically polished up to 1 μm using diamond pastes and non aqueous lubricants until a mirror-bright surface was achieved. All the experiments were performed at room temperature (25 ± 0.1 °C) in naturally aerated sodium borate/boric acid buffer solution of pH 8.4. The p.a. reagents (0.075 M Na2B4O7 and 0.05 M H3BO3) were added in Millipore® water. Potentiodynamic polarization curves were recorded starting from -2400 mV/SCE and moving in the electropositive direction at a scan rate of 1 mV/s, after allowing a steadystate potential to develop. Free corrosion potential was recorded with respect to the SCE every minute for a period of 2 h. All stationary measurements were carried out using an Amel System 5000 potentiostat controlled by a personal computer. Electrochemical impedance was measured at the open circuit potential using a Gamry FAS2 Femtostat with a PC4 Controller. The frequency range analyzed went from 100 kHz up to 10 mHz, with the frequency values spaced logarithmically (seven per decade). The width of the sinusoidal voltage signal applied to the system was 10 mV rms (rootmean-square). Impedance measurements were performed at different exposure times in the electrolyte. For comparison, the electrochemical tests were also performed on unalloyed magnesium, supplied by Johnson Matthey, London, UK. SEM-EDXS were used to investigate morphology and chemical composition of specimens surface after the electrochemical tests. Results and discussion Representative SEM micrographs of the Mg – heavy RE alloys are shown in Figure 1. All tested alloys are basically constituted by a Mg-based matrix, where the rare earth is totally dissolved, giving an average concentration near to the nominal one. In all specimens the magnesium solid solution is the only detected phase. The obtained results are in agreement with the solubility trend of the rare earth metals in magnesium, increasing from Gd to Er [18]. Moreover, the presence of small isolated particles is also detected. These particles may have formed during the melting process and typically contains Mg and O, as estimated using EDXS analysis.

Light metals

(a)

(b)

Fig. 1 - Representative SEM micrographs of Mg99RE1 alloys. (a) Mg99Gd1 alloy; (b) Mg99Er1 alloy (single phase samples) Figure 2 reports the open circuit potential, EOC, of unalloyed Mg and Mg99RE1 alloys in naturally aerated sodium borate/ boric acid buffer solution, monitored for 2 h. The open circuit potential provides some indication of the reactivity of the metal surface. As can be seen, the open circuit potential of each Mg99RE1 alloys and for unalloyed Mg is, as expected, relatively negative immediately after the specimen is immersed in the solution, and subsequently the open circuit potential gradually shifts towards more positive values. This

positive shift can be attributed to deposition of corrosion products on the sample surface. By comparing the results reported in Figure 2, it can be observed that the open circuit potential of Mg99RE1 alloys is less negative than that recorded on unalloyed Mg. This behaviour indicates that the corrosion products layer formed at the surface of Mg99RE1 alloys displays better corrosion protection characteristics than the one formed on unalloyed magnesium in 0.075 M Na2B 4O7 + 0.05 M H3BO3 solution.

Fig. 2 - Variation of the open circuit potential, EOC, with exposure time in naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4, for unalloyed Mg and Mg99RE1 alloys

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Leghe leggere Figure 3 shows typical polarization curves of unalloyed Mg and Mg 99RE1 alloys recorded in naturally aerated sodium borate/boric acid buffer solution at 25 °C. As can be seen, the shapes of the polarization curves are similar indicating similar corrosion mechanisms. One important feature is that the presence of heavy RE provokes a significant shift in the corrosion potential to more positive values as compared to unalloyed Mg, thereby confirming the previous EOC measurements. The anodic response for all specimens is characterized by activation controlled kinetics in the vicinity of the corrosion potential. In fact, the anodic current densities initially increase exponentially with increasing potential above Ecorr. However, as the potential becomes more anodic, the behaviour is no longer exponential and all samples reach a maximum current density after which the current density gradually drops with further increasing potential. The potential where the current density peaks to a maximum is known as the passivation potential, Epass, while the current at this potential is referred to as the critical current density, icc. For unalloyed Mg, the E pass value is −1150 mV/SCE and the icc value is 4.7 mA cm −2. Regarding Mg 99RE1 alloys, the passivation potential is shifted towards more noble values, while the critical current density exhibits lower values. On further scanning in the anodic direction the current density remains independent of potential and a well-defined plate-

au appears thereby indicating the onset of a pseudopassivation process attributable to formation of a corrosion product layer at the surface of corroding sample [26 - 30]. This current is known as the pseudopassivation current density, ipp. As can be observed in Figure 3 the ipp value of unalloyed Mg is 2.6 mA cm −2 while the pseudopassive region extends from about −1150 to 1490 mV/SCE when the surface layer begins to break down in the solution employed and the anodic current density rises again. The presence of rare earths improves the pseudopassivation behaviour, resulting in a larger pseudopassive region extending from about −930 to 1780 mV/SCE for Mg99Gd1 alloy, from about −910 to 1840 mV/SCE for Mg99Dy1 alloy and from −880 to 1920 mV/SCE for Mg99Er1 alloy. The electrochemical parameters deduced from the analysis of the anodic part of the polarization curves recorded in 0.075 M Na 2B4O 7 + 0.05 M H 3BO3 solution are summarized in Table 1. As was previously shown, heavy RE alloying additions increase the extent, ∆E, of the pseudopassive region, thereby indicating the formation of a more stable surface layer on Mg 99RE1 electrodes. Moreover, the RE-containing Mg alloys show pseudo-passive current density values lower than that of unalloyed Mg, suggesting that the corrosion products layer formed on these systems exhibits better protective characteristics.

Fig. 3 - Potentiodynamic polarization curves for unalloyed Mg and Mg99RE1 alloys in naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4

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Light metals Tab. 1 - Electrochemical parameters obtained from the potentiodynamic polarization curves of unalloyed Mg and Mg99RE1 alloys recorded in naturally aerated Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4

Sample

Ecorr (mV/SCE)

Epass (mV/SCE)

icc (mA cm−2)

∆E (mV)

ipp (mA cm−2)

-1790

-1150

4.70

2640

2.60

Mg99Gd1

-1590

-930

3.90

2710

2.20

Mg99Dy1

-1530

-880

3.40

2750

1.90

Mg99Er1

-1480

-830

2.90

2800

1.60

Ecorr = Corrosion potential; Epass = passivation potential; icc = critical current density; ∆E = pseudopassive range; ipp = pseudopassivation current density

In order to obtain e deeper understanding of the corrosion processes, electrochemical impedance spectra (EIS) were measured at hourly intervals for Mg 99RE1 and unalloyed Mg at the corrosion potential during 24 h exposure to Na 2B 4O7 + 0.05 M H3BO 3 solution. The surfaces were examined at the end of each experiment using scanning electron microscopy coupled with EDXS analysis. Figure 4 shows the evolution of the impedance diagrams, in the form of Nyquist plots, for unalloyed Mg and Mg99RE1 alloys a function of exposure time to naturally aerated sodium borate/boric acid buffer solution. The EIS spectra are similar in all cases and changes in a similar manner with increasing exposure time. The Nyquist plots are characterized by two well-defined capacitive loops: a high to medium frequency (HF) and a medium to low frequency (MF) capacitive loop, as labelled. The diameter of the HF capacitive loop typically represents the charge transfer resistance (Rct) of an actively corroding electrode, with determined capacitance values consistent with the electrochemical double layer. The MF capacitive loop is attributed to the presence of the surface film influencing the corrosion process. In this case, the charge transfer resistance through this defective (porous) surface layer is not much higher than for the processes occurring on

the bare surface of specimen. The corrosion can be designed as active dissolution in both cases. The difference is in the different dielectric properties of this surface film, inducing higher capacitance values. The polarization resistance, Rp, is evaluated in this case of poor surface stability as the distance from the impedance values at the highest frequency (theoretically at infinite frequency) to the impedance at the lowest frequency (f = 0.01 Hz) determined on the real part of impedance coordinate in the Nyquist plot. The polarization resistance is related to the corrosion resistance and the corrosion rate; high values of polarization resistance relate to high corrosion resistance and to low corrosion rates. The increasing size of the HF loop typically indicates lateral stabilization of the surface layer resulting in a decreasing actively dissolving surface and as a consequence an higher impedance. For unalloyed Mg and Mg 99RE1 alloys, there is an increase of polarization resistance, indicating that the corrosion resistance increases (attributed to an increased stability of the defective surface product layer).

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Leghe leggere

Fig. 4 - Representative Nyquist diagrams of unalloyed Mg and Mg99RE1 alloys for various exposure times to naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4

Figure 5 presents the polarization resistance, R p, versus exposure time as measured from the Nyquist plots. In all cases the polarization resistance (and the corrosion resistance) increases with increasing exposure time, suggesting that a surface corrosion products layer grows on the electrode surface and that this layer effectively slows down active corrosion. As can be seen, Mg 99RE1 alloys exhibit higher values of R p as compared to unalloyed Mg, thereby indicating that the

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magnesium charge transfer process is significantly inhibited by the alloyed heavy rare earth elements because of the formation of a more protective surface film which develops an effective barrier against corrosion, in agreement with the results obtained from open circuit potential and potentiodynamic polarization measurements (Figs. 2, 3).

Light metals

Fig. 5 - Variation of polarization resistance, Rp, for unalloyed Mg and Mg99RE1 alloys as a function of exposure time to naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4

SEM-EDXS characterization of samples surface at the end of each EIS experiment was carried out in order to gain information about the composition and morphology of corrosion products layer which plays an important role in determining the corrosion behaviour of investigated materials. The different solid corrosion products, formed on their surface at the onset of dissolution, may hinder the transportation of the corrosion medium as well as the mass exchange of reagents. Representative SEM micrographs of unalloyed Mg and Mg99RE1 alloys after 24 h exposure to naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution are shown in Figure 6. A corrosion products layer covers the entire surface of all specimens. A thin oxide/hydroxide film forms spontaneously at the surface of Mg on exposition to ambient air [31, 32]. This surface film is not compact and part of the Mg substrate can be easily exposed. In an aqueous solution, the Mg surface film is covered by a thick porous layer of Mg(OH)2. The dissolution reaction mainly occurs at the bare parts of the Mg surface, and can be summarized [4, 33 - 35] as follows. The corrosion of Mg converts metallic Mg to the stable ion, Mg++, in two electrochemical steps, (2) and (3). These anodic reactions are balanced by the cathodic partial reaction (1). The uni-positive ion, Mg+, is so reactive that it has never been detected [33]. The overall reaction is (4). 2H2O + 2e− → 2OH− + H2 Mg → Mg+ + e− Mg+ → Mg++ + e− 40

cathodic reaction (1) anodic reaction (2) anodic reaction (3)

Mg + 2H2O → Mg(OH)2 + H2

overall reaction

(4)

EDXS analysis performed on the surface layer of Mg99RE1 alloys also evidenced a significant amount of rare earth metals (about 5 at.%) in the form of oxide/hydroxide. The lanthanide cations formed during dissolution of alloy (reaction (5)) react with the hydroxyl ions (reaction (1)) giving rise to an insoluble oxide/hydroxide film that precipitates at the alloy surface (reactions (6) and (7)) [36, 37]: RE → RE3+ + 3e− (5) RE3+ + OH− → RE(OH)3 (6) 2RE(OH)3 → RE2O3 + 3H2O (7) Gd2O3, Dy2O3 and Er2O3 are insoluble in water. The solubility product constants, Ksp, for Gd(OH)3, Dy(OH)3 and Er(OH)3 are 1.1 × 10−22, 1.4 × 10−22 and 1.8 × 10−22, respectively, which are much smaller than that of Mg(OH)2 (5.61 × 10−12). Thus Gd2O3, Dy2O3 or Er2O3 and Gd(OH)3, Dy(OH)3 or Er(OH)3 are more likely to be retained in the corrosion products layer than Mg(OH)2, due to the lower solubility. This results in the enrichment of Gd, Dy or Er in the corrosion products layer. The presence of rare earth oxide/hydroxide film may explain the differences between the surface morphology of pure Mg (Fig. 6a), on one hand, and that of rare earth-containing alloys, on the other, evidenced by SEM observations (Fig. 6b-d). Literature reports improved corrosion resistance of magnesium alloys exposed to aggressive environments containing lanthanide ions [38]. This improvement was attributed to the precipitation of a protective film of rare La Metallurgia Italiana - n. 2 2018

Leghe leggere earth oxides/hydroxides on the cathodic sites. Hybrid silica sol-gel coatings containing lanthanide ions formed on pure magnesium and on magnesium alloys were also tested [39, 40]. It was reported that the sol-gel film modified with rare earths behave as conversion coatings

(a)

(c)

on the metallic substrates. The anticorrosive performance of lanthanide ions entrapped in the hybrid silica solgel network occurs by means of the inhibitor effect and self-repairing mechanism (probably associated with rare RE(OH)3 precipitation) [40].

(b)

(d)

Fig. 6 -Representative SEM micrographs of the corrosion products layer formed at the surface of pure Mg (a); Mg99Gd1 alloy (b); MgDy1 alloy (c); Mg99Er1 alloy (d) after 24 h exposure to naturally aerated 0.075 M Na2B4O7 + 0.05 M H3BO3 solution, pH = 8.4 99

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Light metals The d.c. polarization and EIS results obtained in the present work show significant improvement of the corrosion behaviour of Mg99RE1 alloys in sodium borate/boric acid buffer solution with respect to unalloyed magnesium. This result is linked with the presence of a more stable surface layer responsible for the higher Rp and lower ipp values compared to unalloyed Mg. Since the negative electrochemical potential of the elements Gd, Dy and Er [41], is very similar to that of the element Mg, elemental RE does not cause micro-galvanic corrosion of the Mg matrix. Moreover, RE dissolved in the Mg matrix increases the protective nature of the surface film as evidenced for Mg-Y alloys [42, 43]. The incorporation of oxidized Y in the surface film was identified as enhancing their degradation resistance. Therefore, Gd, Dy and Er, in the oxidized state, can incorporate in the surface layer, increase its protectiveness and thereby decrease the corrosion progress of metallic substrate. Since rare earth oxides/ hydroxides are very insulating they may contribute to enhance the dielectric properties of the corrosion products layer, thus reinforcing its barrier properties. Besides, rare earths alloying addition to pure Mg may impart active corrosion protection properties to the corrosion products layer, further improving its protection ability. Additional active corrosion protection originates from an inhibiting action of the rare earth cations, RE3+, entrapped as oxides/hydroxides in the corrosion products layer. This surface layer hinders the mass exchange of reagents and products between the substrate and the corrosive medium, thus impeding further degradation and consequently increasing the corrosion resistance. Similar explanation was also proposed in literature [40] to interpret the improved corrosion protection of sol-gel film containing La, Ce and Pr.

Further investigation is planned aiming at a deeper understanding of the protective behaviour of these surface layers. Conclusions The corrosion behaviour of Mg99RE1 alloys in naturally aerated sodium borate/boric acid buffer solution has been assessed and compared with that of unalloyed magnesium. The following conclusions can be drawn: 1. Addition of heavy the rare earth elements Gd, Dy and Er significantly improves the corrosion resistance of magnesium. 2. The increased corrosion stability of Mg99RE1 alloys is ascribed to the formation of a “lanthanide-doped” corrosion products layer (magnesium hydroxide, Mg(OH)2), which exhibits higher stability with respect to that formed on unalloyed Mg (higher Rp and lower ipp values). 3. SEM-EDXS characterization of the corrosion products layer present at the surface of Mg99RE1 alloys evidenced significant amounts of rare earth in the form of oxide/ hydroxide. RE elements, in the oxidized state, are incorporated in the surface layer, increasing its protective effectiveness and consequently decreasing the corrosion rate. 4. Owing to the highly insulating character of rare earth oxides/hydroxides the barrier properties of the corrosion products layer are significantly enhanced. Moreover, the presence of lanthanide cations, RE3+, entrapped as oxides/hydroxides in the surface layer is likely to exert an inhibiting action towards the corrosion process, thereby imparting active corrosion protection properties to the corrosion products layer.

REFERENCE [1] K.U. KAINER, Magnesium – Alloys and Technology, Wiley-VCH Verlag Gmbh & Co. KgaA, Weinheim, 2003 [2] H.E. FRIEDRICH, B.L. MORDIKE, Magnesium Technology, Springer-Verlag, Berlin, Heidelberg, 2006 [3] M. POURBAIX, Atlas of Electrochemical Equilibria in Aqueous Solutions, 2nd ed., NACE, Houston, 1974 [4] G. SONG, A. ATRENS, Adv. Eng. Mater. 1 (1999) 11 [5] G. SONG, A. ATRENS, D. ST. JOHN, J. NAIRN, Y. LI, Corros. Sci. 39 (1997) 855 [6] G. SONG, A. ATRENS, D. ST. JOHN, X. WU, J. NAIRN, Corros. Sci. 39 (1997) 1981 [7] G. BARIL, N. PÉBÈRE, Corros. Sci. 43 (2001) 471 [8] G. SONG, A. ATRENS, X. WU, B. ZHANG, Corros. Sci. 40 (1998) 1769 [9] X.W. GUO, J.W. CHANG, S.M. HE, W.J. DING, X. WANG, Electrochim. Acta 52 (2007) 2570 [10] M.C. ZHAO, M. LIU, G. SONG, A. ATRENS, Adv. Eng. Mater. 10 (2008) 104 [11] Y.L. SONG, Y.H. LIU, S.R. YU, X.Y. ZHU, S.H. WANG, J. Mater. Sci. 42 (2007) 4435 [12] Y.L. SONG, Y.H. LIU, S.R. LIU, X.Y. ZHU, Mater. Corros. 58 (2007) 189 [13] A. PARDO, M.C. MERINO, A.E. COY, R. ARRABAL, F. VIEJO, E. MATYKINA, Corros. Sci. 50 (2008) 823 [14] M. ANIK, G. CELITKEN, Corros. Sci. 49 (2007) 1878 [15] A.M. LAFRONT, W. ZHANG, S. JIN, R. TREMBLAY, D. DUBÉ, E. GHALI, Electrochim. Acta 51 (2005) 489 [16] T. ZHANG, Y. SHAO, G. MENG, F. WANG, Electrochim. Acta 53 (2007) 561 [17] G. SONG, D. STJOHN, J. Light Metals 2 (2002) 1 42

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Leghe leggere [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43]

X.W. CHANG, P.H. FU, X.W. GUO, L.M. PENG, W.J. DING, Corros. Sci. 49 (2007) 2612 F. ZUCCHI, V. GRASSI, A. FRIGNANI, C. MONTICELLI, G. TRABANELLI, J. Appl. Electrochem. 36 (2006) 195 W. LIU, F. CAO, L. CHANG, Z. ZHANG, J. ZHANG, Corros. Sci. 51 (2009) 1334 W. LIU, F. CAO, A, CHEN, L. CHANG, J. ZHANG, C. CAO, Corros. Sci. 52 (2010) 627 W. LIU, F. CAO, B. JIA, L. ZHENG, J. ZHANG, C. CAO, X. LI, Corros. Sci. 52 (2010) 639 G.L. SONG, Adv. Eng. Mater. 7 (2005) 563 F. ROSALBINO, E. ANGELINI, S. DE NEGRI, A. SACCONE, S. DELFINO, Intermetallics 13 (2005) 55 A.A. NAYEB-HASHEMI, J.B. CLARK, Phase diagrams of binary magnesium alloys, ASM International, Metals Park, Ohio, 1988 F. EL-TAIB HEAKAL, A. FEKRY, M.Z. FATAYERJI, Electrochim. Acta 54 (2009)1545 H. GAO, Q. LI, F.N. CHEN, Y. DAI, F. LUO, L.Q. LI, Corros. Sci. 53 (2011) 1401 F. EL-TAIB HEAKAL, A.A. GHONEIM, A. FEKRY, J. Appl. Electrochem. 37 (2007) 405 G.T. BURSTEIN, Corros. Sci. 47 (2005) 2858 G. WU, Y. FAN, A. ATRENS, C. ZHAI, W. DING, J. Appl. Electrochem. 38 (2008) 251 M. LIU, P. SCHMUTZ, S. ZANNA, A. SEYEUX, H. ARDELAN, G. SONG, A. ATRENS, P. MARCUS, Corros. Sci. 52 (2010) 562 M. LIU, S. ZANNA, H. ARDELAN, I. FRATEUR, P. SCHMUTZ, G. SONG, A. ATRENS, P. MARCUS, Corros. Sci. 51 (2009) 1115 G. SONG, A. ATRENS, Adv. Eng. Mater. 5 (2003) 837 G. SONG, Adv. Eng. Mater. 7 (2005) 563 G. SONG, A. ATRENS, Adv. Eng. Mater. 9 (2007) 177 M.A. ARENAS, A. CONDE, J.J. DE DAMBORENEA, Corros. Sci. 44 (2002) 511 A.M. CABRAL, W. TRABELSI, W. SERRA, M.F. MONTEMOR, A.L. ZHELUDKEVICH, M.G.S. FERREIRA, Corros. Sci. 48 (2006) 3740 F. EL-TAIB HEAKAL, O.S. SHEATA, N.S. TANTAWY, Corros. Sci. 56 (2012) 86 H. ARDELEAN, I. FRATEUR, P. MARCUS, Corros. Sci. 51 (2009) 3030 A.L. RUDD, C. BRESLIN, F. MANSFELD, Corros. Sci. 42 (2000) 275 G. MILAZZO, S. CAROLI, Tables of Standard Electrode Potentials,Wiley, New York, 1978 P.L. MILLER, B.A. SHAW, R.G. WENDT, W.C. MOSHIER, Corrosion 51 (1995) 922 H.B. YAO, Y. LI, A.T.S. WEE, Electrochim. Acta 48 (2003) 4197

La Metallurgia Italiana - n. 2 2018

premio Aldo Daccò 2018

premio Aldo Daccò 2018 L’AIM è lieta di indire il bando per l’edizione 2018 del prestigioso Premio Aldo Daccò, con l’obbiettivo di stimolare i tecnici del settore e contribuire allo sviluppo e al progresso delle tecniche di fonderia e di solidificazione con memorie e studi originali. L’Associazione invita tutti gli interessati a concorrere al Premio “Aldo Daccò” 2018, inviando a mezzo email (info@aimnet.it), il testo di memorie inerenti le tematiche fonderia e solidificazione, unitamente al curriculum vitae dell’autore concorrente, entro il 31 luglio 2018. Saranno presi in considerazione e valutati i lavori riguardanti le varie tematiche di fonderia e di solidificazione, sia nel campo delle leghe ferrose che in quello delle leghe e dei metalli non ferrosi. Il premio, pari a Euro 3500 lordi, è offerto dalla Fondazione Aldo e Cele Daccò, istituita dalla signora Cele Daccò per onorare la memoria del marito Aldo Daccò, uno dei soci fondatori dell’AIM e suo encomiabile Presidente per molti anni. Le memorie verranno esaminate da una Commissione giudicatrice designata dal Consiglio Direttivo, il cui giudizio sarà insindacabile. Nel giudicare, la Commissione terrà conto, in particolar modo, dell’originalità del lavoro e dell’argomento in relazione alla reale applicabilità dei risultati. Non sono ammesse candidature da chi abbia già ottenuto riconoscimenti, anche per lavori diversi, dalla Fondazione Aldo e Cele Daccò per la ricerca scientifica. Le memorie premiate e quelle considerate meritevoli di segnalazione, potranno essere pubblicate sulla rivista La Metallurgia Italiana. La cerimonia di premiazione avrà luogo a Bologna, in occasione del 37° Convegno Nazionale AIM.

Per informazioni e candidature: Associazione Italiana di Metallurgia Via Filippo Turati 8 · 20121 Milano Tel. 02-76397770 · E-mail: info@aimnet.it

I vincitori del Premio 1975 1979 1981 1982 1983 1984 1986 1987 1989 1990 1991 1992 1993 1994 1995 1996 1997 1998 1999 2001 2002 2003 2004 2005 2006 2007 2008 2009 2010 2011 2012 2013 2014 2016 2017

M. Remondino, F. Pilastro, E. Natale A. Goria, M. Mischiatti E. Borghigiani R. Medana R. Medana E. Borghigiani, F. Belletti R. Medana L. Piras, L. Lazzaro P. Fumagalli F. Cavalleri, G. Tosi, A. Pedaci R. Roberti, A. Bianchi, F. Pedrotti R. Maspero, B. Calzolai E. Gariboldi, G. Caironi G. Zaramella G. P. Marconi, A. Boccardo R. Medana R. Nada C. Bolner A. Gregori, F. Bonollo C. Mapelli L. Battezzati, P. Ferro R. Venturini, S. Baragiola E. Liotti, F. Piasentini, F. Bonollo, A. Tiziani C. Viscardi D. Baldissin, M. Di Sabatino G. Timelli, A. Manente A. Arrighini M. Merlin A. Morri E. Zanini, G. Timelli D. Casari, C. Soffritti M. Alloni, R. Carli D. Gelli F. De Antoni, M. T. Di Giovanni da assegnare

Le manifestazioni AIM AIM meetings and events 2018 SOLIDIFICAZIONE E COLATA CONTINUA Corso itinerante - Centro A 8-9-15-16-22-23 marzo I METALLI PER L’EDILIZIA SOSTENIBILE. ACCIAIO E RAME LA CERTIFICAZIONE DEI FABBRICATI GdS – Centro MTA Milano, 21 marzo

VALUTAZIONE DELL’ESERCIBILITA’ DEI MATERIALI NEI CICLI COMBINATI TRADIZIONALI ED INNOVATIVI GdS - Centro ME Milano, 28 giugno RADDRIZZATURA E TENSIONI RESIDUE DEL GETTO GdS - Centro P Ceregnano (RO) c/o TMB, giugno

DEFORMAZIONE PERMANENTE: DALL’ACCIAIO AL PROCESSO GdS - Centro TTM Piacenza, 22 marzo

MATERIALI METALLICI E PROCESSI PRODUTTIVI INNOVATIVI PER L'AEROSPAZIO Convegno - Centri ML, MFM e MP Napoli, 19-20 luglio

METALLURGIA DELLE POLVERI Scuola di - Centro MP Imola (BO) c/o SACMI, 19-20 aprile

METALLURGY SUMMER SCHOOL - 2a edizione COMET Bertinoro (FC), luglio

COME GARANTIRE LA CONFORMITÀ DELLE MACCHINE ANCHE A SEGUITO DI MODIFICHE: FORMA E SOSTANZA GdS - Centro AS aprile

LA PREVENZIONE E LA GESTIONE DELLE MALATTIE PROFESSIONALI GdS - Centro AS luglio

AUMENTO DELLA PRODUTTIVITA’ DEGLI STAMPI ATTRAVERSO UN CONTROLLO SPECIFICO DELLA FATICA TERMICA GdS - Centro P Bergamo, 9-10 maggio

37° CONVEGNO NAZIONALE AIM Convegno – SEGR Bologna, 12-14 settembre

METALLURGIA DELLE POLVERI Scuola di - Centro MP Maerne di Martellago (VE) c/o POMETON, 10-11 maggio METALLURGIA DI BASE PROPEDEUTICA AI TRATTAMENTI TERMICI Corso - Centro TTM Milano, 16-17-23 maggio PROFILATI ESTRUSI DI ALLUMINIO: PERCHE’? CONVENIENZA E UTILITA’ NEI VARI SETTORI APPLICATIVI GdS - Centri ML e LPM Milano, 7 giugno ICS 2018 - 7TH INTERNATIONAL CONGRESS ON SCIENCE & TECHNOLOGY IN STEELMAKING Convegno Internazionale e 26° CONVEGNO NAZIONALE TRATTAMENTI TERMICI Convegno Venezia, 13-15 giugno

EOSC 2018 - 8TH EUROPEAN OXYGEN STEELMAKING CONFERENCE Convegno Internazionale Taranto, 10-12 ottobre TRATTAMENTI TERMICI DEGLI ACCIAI PER STAMPI A CALDO E A FREDDO PER IL SETTORE AUTOMOTIVE GdS - Centro TTM ottobre UTENSILI DIAMANTATI GdS - Centro MP Vicenza, 15 novembre GETTI STRUTTURALI GdS - Centro P Brescia, 16 novembre CLEAN TECH - 4TH EUROPEAN CONFERENCE ON CLEAN TEHNOLOGIES IN THE STEEL INDUSTRY Convegno Internazionale Bergamo, 29-30 novembre RIVESTIMENTI - 1° modulo Rivestimenti PVD e CVD Corso modulare - Centro R Roma, novembre

Per ulteriori informazioni rivolgersi alla Segreteria AIM e-mail: info@aimnet.it oppure visitare il sito internet www.aimnet.it

La Metallurgia Italiana - n. 2 2018

Industry news A comparative cradle-to gate impact assessment: primary and secondary aluminum automotive components case a cura di: S. Cecchel, M. Collotta, G. Cornacchia, A. Panvini, G. Tomasoni Road transports release a significant percentage of global CO2. In this field, one of the most effective solution is the reduction of vehicles’ mass, which can be obtained through the substitution of heavy metals with light alloys (i.e. aluminum). In addition, in order to maximize the environmental benefits a current trend is to use secondary material (from scrap) in substitution to primary one (from ore). For this purpose, the present case study compared the environmental burden related to the same light-weighted automotive component (suspension cross beam) made in primary aluminum (from ore) or in secondary one (from scrap). In particular, a cradle to grave Life Cycle Assessment has been analyzed through the software SimaPro 7.3 with the Recipe impact method. The study highlighted and confirmed the relevance of the environmental benefits related to recycling and secondary material use.

KEYWORDS: LIFE CYCLE ASSESSMENT (LCA); ALUMINUM PROCESS; AUTOMOTIVE COMPONENTS; PRIMARY ALUMINUM; SECONDARY ALUMINUM; RECYCLING

S. Cecchel, M. Collotta, G. Cornacchia, A. Panvini, G. Tomasoni DIMI, Department of Industrial and Mechanical Engineering, University of Brescia

INTRODUCTION Automotive components weight reduction is a matter of great importance and emerging interest, in order to reach a better world sustainability. In fact, as is well known, the lightweighting has a direct influence on the improvement of vehicle’s performances and is strictly connected to the reduction of fuel consumption and emissions [1-4]. Present EU legislation sets mandatory emissions reduction targets for new vehicles by 2020 [5], highlighting that this is a topic of primary interest. At this aim, the use of low density materials (i.e. aluminum alloys) in substitution of “heavy” metals (i.e. steel and cast iron) has exponentially increased during the last few years [6, 7]. In particular, High Pressure Die Casting is the most used by metalworking industry for high-volume of aluminum alloy components, thanks to the cost advantages derived from its

high productivity rates [8]. Currently, for the manufacturing of structural components is mainly used primary aluminum (made from ore). It is clear that a major use of secondary materials would give several environmental and economic advantages such as: a reduction of the energy consumption (-95% than the production from mineral) [9, 10], a decrease of non-renewable resource consumption and a lower socio-economic impact related to the bauxite mining [11]. Anyway, the application of secondary aluminum in this class of components is limited by the difficulty in the removal of some impurities (iron and copper) during the recycling that can lead to a decrease of the mechanical properties [12]. The most common solution used in industry today is dilution with primary [12]. In addition, recently some new technologies (i.e. accurate separation scrap process) have been developed to obtain aluminum scrap with

La Metallurgia Italiana - n. 2 2018

Attualità industriale higher chemical purity that would allow the extension of recycled material use to heavier missions [12]. Taking into account the environmental benefits arising from automotive components weight reduction, many studies have been published, almost all based on the life cycle assessment methodology. The advantages of substituting iron and steel with lighter materials, including aluminum have been confirmed [13, 14]. In particular, some of these studies focus on the entire automotive sector, discussing the benefits of vehicles light weighting at a global level [2, 15-18] or at regional [18]. Anyway, the lightening of vehicles may have other environmental effects than the reduction of emissions; such impacts are not necessarily positive and thus should be carefully taken into account. Moreover, it is important to highlight that very few works which investigate in detail the foundry production step by step have been found [8, 19, 20]. In this context, a previous work [21] presented the application of a model for the environmental analysis of HPDC at the production of a safety relevant primary aluminum component (cross beam suspension) [22], that confirmed the high benefit of aluminum recycling at the end of life. In the present research, the model has been extended and implemented in order to analyze the real environmental benefit given by the use of components made of secondary alloys. In particular, a cross beam suspension produced in primary aluminum (scenario 1), in secondary aluminum (scenario 2) and in a mix of 50% primary aluminum and 50% secondary one (scenario 3) has been analyzed and compared through a Life Cycle Assessment (LCA) and a sensitivity analysis. In particular, these aspects have been investigated through a LCA software SimaPro 7.3 using the Recipe impact method that calculates the environmental burdens for different impact categories. MATERIALS AND METHODS Life cycle assessment The LCA analysis is an innovative technique for assessing the environmental aspects and potential impacts associated with a product or process throughout its life cycle. The main reference for the LCA methodology is the ISO 14040 standard [23], which describes the principles and framework for LCA, including: the definition of the goal and scope, the Life Cycle Inventory analysis (LCI) phase, the Life Cycle Impact Assessment (LCIA) phase, and the life cycle interpretation phase. In the following section, each phase is discussed and the results are presented. Since the environmental evaluation of a product is characterized by a very large number of data and

La Metallurgia Italiana - n. 2 2018

complex analyses, particularly during LCI and LCIA phases, a “LCA software” (in this case SimaPro 7.1 was used) can be a very useful tool to support the calculations and to present the results obtained. Moreover, the software may access directly specific databases containing secondary inventory data, i.e. inventory data coming from literature (in this case Ecoinvent database [24]). Finally, a software allows to apply several impact assessment methods (in this case, the Recipe method was adopted) that calculate the environmental severity for different impact categories through the transformation of the long list of LCI results, into a limited number of indicator scores. Goal and scope definition The aim of this LCA is to identify the potential environmental benefits arising from the production with secondary vs primary aluminum of an HPDC component for light commercial vehicles, using real data. In particular, a comparison of three different scenarios in which the component is made of (1) 100% primary aluminum, (2) 100% secondary aluminum and (3) 50% primary and 50% secondary aluminum is investigated. In detail, with reference to scenario 1, the first life cycle stage for primary aluminum components production consists in the extraction of the ore and its transformation into a primary aluminum ingot, through the following operations: bauxite mining, alumina production, electrolysis and cast house. It’s worthwhile to note that the electrical energy required for the primary smelting process constitutes the major part of energy consumption in primary aluminum production [9, 10]. In particular, the total consumption has to take into account different elements: rectifying loss, DC power usage, pollution control equipment, auxiliary power (general plant use), electric transmission losses from power stations to primary smelters. Specific consumption data collected from smelters and the related true weighted average have been elaborated from European Aluminum Association and are available in literature [7]. With reference to scenario 2, secondary aluminum ingot production involves in an accurate physically separating solid scrap stream to prevent co-mingling and elements [12, 13], melting and alloying in order to reach the required composition. It’s important to underline that this allow to completely avoid the highly energy transformation process needed for ore reduction above introduced. Finally, in scenario 3 a dilution of secondary aluminum with 50% primary aluminum is investigated. The subsequent phases are the same either from primary or

Industry news secondary ingot production. In detail, the component is realized through HPDC with the following main important foundry phases: aluminum ingot melting, molten metal holding, and casting. Then, the foundry system (i.e. sprue, air venting, and heat flow), that has to be trimmed off and recycled by remelting [25, 26]. Subsequently, the castings are subjected to machining in order to obtain the final design required by the component features. Finally, components are recycled at their end of life. Life cycle inventory The functional unit adopted for all the scenarios analyzed is a typical production batch of 250 aluminum components for light commercial vehicles. The component weight is about 15 kg. The use phase of the component is not reported; in fact, having the components the same geometry and the same weight, input and output of mass and energy for the use phase does not differ among scenarios. Whenever possible data taken from the field were used. In particular, foundry and machining phases have been evaluated through the use of the model already mentioned and with the most relevant assumption reported in the authors’ previous work [21]. In this context, it is important to highlight

that most of the parameters related to the foundry operation reported in this section are related to primary data achieved thank to the cooperation of some automotive Italian foundries. For the shake of clarity, the aluminum scraps due to the feeder system are reported as “yielding ratio”. Secondary data regards the remaining phases and has been gathered from literature [9, 10] and Ecoinvent database [24]. In the current case study the following parameters are considered: • Components weight 15 kg, • Foundry time cycle 3 min, • Foundry yielding ratio 40%, • Foundry melting loss 5%, • HPDC cold chamber machine with a clamping force of 3,000 t (about 30,000 kN). • 5 axis-machining, • Machining scrap ratio 1% • Machining time cycle 8 min, • Lost (not recycled) 6% • Shredder loss 5% • Re-melting loss 5% The detail of input/output flows for each scenario are shown and in Fig. 1 and Tab. 1- Tab. 3.

Fig. 1 - Flow charts for scenario 1 (S1), scenario 2 (S2) and scenario 3 (S3) 48

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AttualitÃ industriale Tab. 1 - Input/output flows and process unit for scenario 1 (S1)

INPUT Fresh Water (m3) Heavy fuel oil (kg) Diesel (l) Electricity (kWh) Sodium hydroxide (kg) Lime (kg) Heat, natural gas (MJ) Heat, heavy fuel (MJ) Heat, light fuel (MJ) Steam (MJ) Anode & paste production (kg) Aluminum fluoride (kg) Cathode (kg) Steel (kg) Cryolite (kg) Refractory material (kg) Transport by ship (tkm) Transport by barge (tkm) Transport by lorry (tkm) Transport by rail (tkm) Chlorine (kg) Argon (kg) Nitrogen (kg) Aluminum scrap (kg) Silicon (kg) OUTPUT Carbon Dioxide (kg) Particulates (kg) Sulfur dioxide (kg)

Bauxite Mining (13,24 ton)

Alumina Production (5,89 ton)

Aluminum liquid electrolysis (3,07 ton)

Primary ingot production (4,09 ton)

5,96

2,95

4,30

-1,63

1066,63

45,67

Foundry (263 pcs)

Finishing (250 pcs)

400,53

1458

800

25334,01

5529,71

37907

34314,94

760,18

2,65 4,67 11,92

312,33 247,51

188,00 1467,36 1350,51 48,50 21,18 11,66 4,91 24,55 11672,26 705,73 99,91 4992,22 0,20 8,62 0,90 997,23 98,09 26,48

4914,76

4831,14

461,83

2,25

0,83

2,58

0,16

15,79

22,71

0,61

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Industry news Alumina Production (5,89 ton)

Aluminum liquid electrolysis (3,07 ton)

Primary ingot production (4,09 ton)

Nitrogen oxides (kg)

6,54

1,35

0,86

Mercury (kg)

0,35

Suspended solids (kg)

1,36

Oils (kg)

0,59

Waste (kg)

282,86

Bauxite Mining (13,24 ton)

Finishing (250 pcs)

1,39

11,85

Flouride (kg)

2,67

PAH (kg)

0,04

Benzo(a)pyrene (kg)

0,80

Methane (kg)

0,12

Ethane (kg)

0,01

PFC (kg)

0,12

Disposal (kg)

31,00

8,99

Hydrogen Chloride (kg)

0,08

Dross (kg)

72,75

Dust (kg)

3,68

Refractory material (kg)

2,86

New scrap (ton)

Foundry (263 pcs)

271

1526

Chlorine (kg) TOC (kg) COD (kg)

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Attualità industriale Tab. 2 - Input/output flows and process unit for scenario 2 (S2) Secondary ingot production (4,09 ton) INPUT Fresh Water (m3) Electricity (kWh) Heat, natural gas (MJ) Heat, heavy fuel (MJ) Heat, light fuel (MJ) Chlorine (kg) Argon (kg) Nitrogen (kg) Aluminum scrap (kg) Gas (m3) OUTPUT Carbon Dioxide (kg)

Foundry (263 pcs)

Finishing (250 pcs)

506,79

1458

800

15052,42

37907

23,70

314,70 273,83 1,23 6,95 2,04 4254,57 3,27

2,67 0,04 0,80

1083,06

0,01 0,12

Sulfur dioxide (kg) Nitrogen oxides (kg) Waste (kg) Disposal (kg) Hydrogen Chloride (kg) Dross (kg) Dust (kg) New scrap (ton) Chlorine (kg)

264,43 1443,94 22,07 53,95 63,76 204,35 214,57 1526 5,31

Life Cycle Impact Assessment The LCIA phase consists in the quantification of the impacts on the environment caused by the consumption of the natural resources and by the releases to the pollutant identified and quantified during the LCI. For this analysis, the Recipe model has been employed, which adopt a damage oriented approach and leads to the evaluation of environmental impacts with reLa Metallurgia Italiana - n. 2 2018

271

spect to several categories that declines the impacts affecting the human health, the ecosystem quality and resources availability. It is an “endpoint” method, i.e. it allows to perform characterization, normalization, grouping and weighting, [27] and it allows to choose among three “cultural perspectives” for the calculation of the environmental impact; in this case, the egalitarian perspective was adopted, which is based on 51

Industry news precautionary principle thinking. Table 4 summarizes the main results obtained through the analysis, in particular, the environmental impacts after normalization and the single score. The graph in Figure 2 reports the single score impacts. The normalization is obtained by dividing the impacts in each category by the shares of an average personâ&#x20AC;&#x2122;s emission and

resource use in the world during one year. In practice, normalisation converts complicated units into fractions of the average person's scores per impact category and allows to identify the categories where the impact is more relevant. The calculation of the single score allows to obtain a unique adimensional indicator measured in Points (Pt).

Tab. 4 - Comparison between the impact values of the Recipe method related to the different scenarios: after normalization (a) and single score (b) S1

Normalization

Single Score

Normalization

Single Score

Normalization

Single Score

Climate change Human Health

90,24

27072,80

14,79

4436,54

55,80

16738,54

Ozone depletion

0,00

1,30

0,00

0,37

0,00

0,89

Human toxicity

280,75

84224,45

22,01

6603,91

160,76

48228,14

Photochemical oxidant formation

0,00

0,61

0,00

0,08

0,00

0,37

Particulate matter formation

11,66

3498,33

0,73

218,37

6,58

1973,75

Ionising radiation

0,10

28,73

0,01

1,76

0,05

16,20

Climate change Ecosystems

4,95

2474,43

0,81

405,41

3,06

1529,84

Terrestrial acidification

0,02

8,46

0,00

0,67

0,01

4,85

Freshwater eutrophication

0,00

2,23

0,00

0,16

0,00

1,27

Terrestrial ecotoxicity

0,04

21,03

0,00

1,16

0,02

11,76

Freshwater ecotoxicity

0,00

0,33

0,00

0,02

0,00

0,19

Marine ecotoxicity

0,00

1,64

0,00

0,09

0,00

0,92

Agricultural land occupation

0,03

15,34

0,01

3,01

0,02

9,73

Urban land occupation

0,02

8,36

0,00

1,76

0,01

5,36

Natural land transformation

0,73

367,30

0,21

106,86

0,50

250,60

Metal depletion

0,07

14,79

0,01

2,77

0,05

9,30

Fossil depletion

117,81

23561,53

24,74

4948,95

75,72

15144,39

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Attualità industriale

Fig. 2 - Comparison between the impact values of the Recipe method related to the different scenarios single score

It can be noted that scenario 1 has the highest impact with 141’302 points, followed by scenario 3 with 83’962 points, while the lowest environmental impact is obtained by scenario 2, which has a total score of 16’723 points. The most relevant impact categories are human toxicity, fossil depletion and climate change human health for all the scenarios. These parameters are mainly related respectively to the air emissions during the electrolysis (for scenarios with primary aluminum) and during the aluminum melting in the foundry operations (for all the scenarios), with a strongly higher impact for the former phase. In particular, taking as baseline S1, human toxicity is respectively 92% and 43% lower for S2 and S3, fossil depletion is respectively 79% and 36% lower for S2 and S3 and climate change human health is respectively 84% and 38% lower for S2 and S3. Therefore, these elevate difference between the environment indicators among different scenarios highlight the relevance of secondary aluminum use in order to reduce the transports pollution. In order to further investigate the sensitivity of the results to the uncertainty in the inventory data, a Monte Carlo analysis has been developed. In particular, through the Monte Carlo

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analysis the inventory parameters are transformed into stochastic variables with a log-normal distribution. The log normal distribution was adopted as it is widely used in literature [28] for Monte Carlo analysis in the context of LCA case studies. The average value of each distribution coincides with the value of the correspondent deterministic inventory parameter, while the standard deviation has been taken from the Ecoinvent database. The values of the inventory parameters are then randomly sampled and the impact assessment method is applied for a high number of combinations of such values. Figure 3 shows, for each environmental impact category of the Recipe method, the probability for scenario 1 to have a lower environmental impact with respect to scenario 3 (and viceversa). As can be seen, the probability for scenario 3 to have a lower environmental impact than scenario 1 is very high for all the impact categories. The same result was obtained with the comparisons between scenario 1 and scenario 2 and between scenario 2 and scenario 3. These evidences confirm the robustness of the results obtained.

Industry news

Fig. 3 - Results of the Monte Carlo analysis applied to the comparison of scenario S1 with scenario S3

CONCLUSION The present study analyzed the environmental benefits connected with the production of lightweight commercial vehicles components. In particular, a suspension cross beam for LCVs made in primary aluminum (from ore) or in secondary one (from scrap) was studied through a cradle to grave LCA filled with real field data. The impact of each phase to the overall life cycle has been evaluated. In particular, it was confirmed that the higher impact assessment derives from electricity and heat demands in the liquid aluminum electrolysis phase included in primary aluminumâ&#x20AC;&#x2122;s scenario (scenario 1 and 3). Regarding secondary aluminumâ&#x20AC;&#x2122;s scenario (scenario 2 and 3) the greater impact was related to the production of ingot from scrap, although it naturally appeared clearly lower than that from ore. It is a common point for all scenario that energy demand for the production of aluminum ingot became the primary factor responsible for the corresponding environmental impact, directly followed by the significant impact related to energy and heat demand for the componentâ&#x20AC;&#x2122;s casting. This last contribution demonstrates the importance

of the new approach of a deep and detailed foundry phase investigation to avoid underestimated results. Furthermore, it is worthwhile to note that in all three different scenarios recycling of aluminum provides a positive benefit in terms of energy savings. This result confirmed the relevance of the present trend in the development of new technologies to obtain aluminum scrap with higher chemical purity that would minimizes the current difference of mechanical properties between primary and secondary aluminum alloys. Finally, it was emphasized that the well-known environmental benefit given by automotive lightweighting can be improved and maximized by the substitution of aluminum primary alloys with secondary one. Acknowledgements The component object of the study is developed by Streparava SpA. The Authors are grateful to Streparava SpA for the support in providing data.

La Metallurgia Italiana - n. 2 2018

Attualità industriale REFERENCES [1] [2] [3] [4]

[5]

[6] [7] [8] [9]

[10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]

[21] [22] [23] [24] [25] [26] [27] [28]

H. Helms and U. Lambrecht, Int. J. LCA (2006) doi:10.1065/lca2006.07258. R. Modaresi, S. Pauliuk, A. N. Lovik, D. B. Muller,: Environ. Sci. Technol. (2014), doi: 10.1021/es502930w. S. Cecchel, D. Chindamo, E. Turrini, C. Carnevale, G. Cornacchia, M. Gadola, A. Panvini, M. Volta, D. Ferrario, R. Golimbioschi.; Science of the Total Environment, 613-14, 409, (2018). http://dx.doi.org/10.1016/j.scitotenv.2017.09.081 S. Cecchel, D. Chindamo, M. Collotta, G. Cornacchia, A. Panvini, G. Tomasoni, M. Gadola; Lightweighting in light commercial vehicles: cradle-to-grave life cycle assessment of a safety relevant component, IJLCA, p. 1-12 (2018), https:// doi.org/10.1007/s11367-017-1433-5 European commission, “Proposal for a regulation of the European Parliament and of the Council amending Regulation (EC) No 443/2009 to define the modalities for reaching the 2020 target to reduce CO2 emissions from new passenger cars” (IOP Publishing PhysicsWeb, 2012), http://eur-lex.europa.eu/resource.html?uri=cellar:70f469933c49-4b61-ba2f-77319c424cbd.0001.02/DOC_1&format=PDF. Accessed 11 July 20012. J. Hirsch,: Mater. Forum 28, 15 (2004) J. Kasai,: JSAE Rev. 20, 387 (1999) D.R. Gunasegaram and A. Tharumarajah, Metal. Trans. B 40, 605 (2009). European Aluminium Association. “Environmental Profile Report for the European Aluminium Industry - Life Cycle Inventory data for aluminium production and tranformation processes in Europe”. (IOP Publishing PhysicsWeb, 2013) http://european-aluminium.eu/media/1329/environmental-profile-report-for-the-european-aluminium-industry.pdf. Accessed 2013. European Aluminium Association. “Aluminium Recycling in Europe” (IOP Publishing PhysicsWeb,2006) http:// recycling.world-aluminium.org/fileadmin/_migrated/content_uploads/fl0000217.pdf. Accessed 2006. F. Bonollo, I. Carturan, G. Cupitò, R. Molina,: Metal. Sci. Tech 24, 3 (2006). G. Gaustad, E. Olivetti, R. Kirchain,: Resour. Conserv. Recy. (2012), doi:10.1016/j.resconrec.2011.10.010. R. Modaresi, D.B. Muller,: Environ. Sci. Technol. 46, 8587 (2012). H.C. Kim and T.J. Wallington, Environ. Sci. Technol. 47, 6089 (2013). H. J. Kim, G. A. Keoleian, S. J. Skerlos, J. Ind. Ecol. (2011), doi:10.1111/j.1530-9290.2010.00288.x H. Kim, C. McMillan, G.A. Keoleian and S.J. Skerlos, J. Ind. Ecol. 14, 929 (2010). G. A. Keoleian, J. L. Sullivan,: MRS Bullettin (2012), doi:10.1557/mrs.2012.52. P. Puri, P. Compston, V. Pantano,: Int. J. Life Cycle Assess. (2009), doi: 10.1007/s11367-009-0103-7. R. Singh, C.D. Singh and S. Singh Sidhu, IJMDEBM 1, 21 (2013). S. Dalquist and T. Gutowski, Life Cycle Analysis of Conventional Manufacturing Techniques: Die Casting (Massachusetts Institute of Technology, Paper no. LMP-MIT-TGG-03-12-09-2004). http://web.mit.edu/ebm/Publications/Die%20 Casting%20Dalquist%20Gutowski.pdf. Accessed 8 June 2014. S. Cecchel, G. Cornacchia, A. Panvini,: JOM 68, 2443 (2016). S. Cecchel, D. Ferrario; La Metallurgia Italiana 108, p. 41-44, (2016). International Standard Organisation (ISO), ISO14040: Environmental Management—Life Cycle Assessment: Principles and Framework (Geneva: ISO, 2006), pp. 1–20. R. Frischknecht, et al.,: Int. J. Life Cycle Assess. 10, 3 (2005). S. Nagendra Parashar and R.K. Mittal, in Elements of Manufacturing Processes (New Delhi, IN: Prentice-Hall of India Private Limited, 2006), p. 233. H. Bakemeyer, Report No. E-902 (NADCA, Arlington Heights, IL, 2008). M. Goedkoop et al.: “ReCiPe 2008.” (IOP Publishing PhysicsWeb,2009) http://www.leidenuniv.nl/cml/ssp/publications/recipe_characterisation.pdf. Accessed 6 January 2009. Frischknecht, R. et al., 2005. The ecoinvent database: Overview and methodological framework. International Journal of Life Cycle Assessment, 10, pp.3–9.

La Metallurgia Italiana - n. 2 2018

Industry news PLANBLUE – Una visione della crescita sostenibile e della sua realizzazione a cura di: Ufficio Stampa Ronal Group In occasione del salone REIFEN di Essen, RONAL GROUP ha presentato PLANBLUE, una nuova ruota prodotta secondo i principi della sostenibilità e della responsabilità sociale ed ecologica nei confronti di clienti, dipendenti e popolazione. Il RONAL GROUP offre un ottimo esempio nelle questioni di responsabilità ambientale e sociale con il cerchione RONAL R60 blu presentato per la prima volta a Essen. L’innovazione nella produzione viene realizzata in maniera molto efficiente dal punto di vista energetico con il 100% di elettricità verde e in condizioni di lavoro ottimali. Con i cerchioni R60 blu diminuisce sensibilmente l’impatto ambientale grazie alla riduzione di peso, la messa a disposizione di materiali e le ottimizzazioni nel processo produttivo. In questo modo il cerchione è, da un lato, ambasciatore e pietra miliare del progetto di sostenibilità PLANBLUE della ditta, e dall’altro, precursore di una nuova generazione di cerchioni eco-compatibili. “Con PLANBLUE realizziamo la piattaforma per la crescita sostenibile. Insieme a tutti i dipendenti del RONAL GROUP voglio adempiere alla responsabilità ecologia e sociale che abbiamo nei confronti dei nostri clienti e della società e perseguire i nostri obiettivi di sostenibilità”, dichiara Yvo Schnarrenberger, CEO del RONAL GROUP. In vari ambiti sono già stati attuati proficui provvedimenti: Promozione di un ambiente di lavoro ottimale: Tra cui possibilità di sviluppi ulteriori per tutti i dipendenti, un codice per una direzione aziendale responsabile, provvedimenti per la sicurezza sul lavoro quali check-up gratuiti per i dipendenti nel settore produzione o equipaggiamento di protezione. Inoltre il RONAL GROUP aspira a rispettare la conformità normativa della clientela.

Orientamento all’innovazione: Materiali e processi di produzione innovativi per la riduzione della percentuale di alluminio per mezzo di ruote il più leggere possibile che possono contribuire a un minor consumo di carburante dei veicoli. Realizzazione efficiente dal punto di vista energetico: Produzione eco-compatibile in impianti centralizzati ed efficienti dal punto di vista energetico: Grazie a sette nuovi impianti per il trattamento termico annualmente si consuma il 36,5% di energia in meno per kWh per kg di alluminio. Con un impianto pilota per l’utilizzo del calore delle fornaci è già stato possibile risparmiare all’anno circa il 38% del consumo di gas dell’impianto. Adempimento dei modernissimi standard: Tutti gli impianti produttivi del RONAL GROUP sono certificati secondo DIN EN 14001 per la gestione ambientale. Inoltre, le sedi di Härkingen, Landau e Forst dispongono della certificazione energetica DIN EN 50001. Già il 25% dell’alluminio viene ritirato dai membri della «Aluminium Stewardship Initiative» (ASI) che difende rigidi criteri di sostenibilità. RONAL GROUP, azienda con sede principale ad Härkingen, in Svizzera, rientra tra i produttori leader nel settore delle ruote in lega leggera per automobili e veicoli commerciali. Conta su 7.500 collaboratori in tutto il mondo.

Produzione equa e locale: Vie di trasporto il più brevi possibile, modernissimi standard di produzione e compensi giusti sono considerati la base della soddisfazione di clienti, dipendenti e partner commerciali. Riciclo: Ruote riutilizzabili al 100% e, per quanto possibile a livello tecnico, materiali di scarto in alluminio completamente riutilizzati nel processo di produzione. Inoltre, tutti gli imballaggi di cartone sono in 100% materiale riciclato. 56

La Metallurgia Italiana - n. 2 2018

PREMIO di STUDIO ASSOFOND 2018

ASSOFOND Federazione Nazionale Fonderie

ASSOFOND Federazione Nazionale Fonderie, è lieta di istituire in occasione del XXXIV Congresso Nazionale di Fonderia,

5 PREMI DI STUDIO del valore di 3000 Euro cadauno

destinati a 5 studenti universitari di ingegneria che abbiano trattato o tratteranno negli anni accademici 2016/2017 e 2017/2018, per lo svolgimento della propria tesi di Laurea Magistrale il seguente argomento: La Progettazione di nuovi componenti realizzati per fusione di leghe Ferrose o leghe non Ferrose in sostituzione dei corrispondenti particolari realizzati utilizzando altre metodologie di produzione, modificando il disegno del componente ed utilizzando le specifiche caratteristiche del processo fusorio, con l’obiettivo di far conoscere le potenzialità offerte dai componenti realizzati per fusione. Assofond invita tutti gli interessati al concorso ad inviare la domanda, redatta in carta libera, ✓ ad ASSOFOND via Copernico 54 Trezzano sul Naviglio o via mail a info@assofond.it ✓ entro e non oltre il 30 settembre 2018. Nella domanda il candidato dovrà indicare, sotto la propria responsabilità, oltre il cognome e nome: ✓ luogo e data di nascita; ✓ codice fiscale; ✓ domicilio eletto ai fini del concorso e recapito telefonico; ✓ elenco esami sostenuti con relativo voto; ✓ copia della tesi di Laurea Magistrale e voto.

La Commissione Giudicatrice è nominata dal Comitato di Presidenza di Assofond ed è costituita da tre membri, il cui giudizio è insindacabile. (Presidente di Assofond o da persona da Lui designata, che ne è Presidente; dal Presidente del Centro di Studio AIM-ASSOFOND per la fonderia o da un suo delegato e dal Presidente di AIM o da un suo delegato). ✓ Nel giudicare, la Commissione terrà conto, in particolare modo, dell’originalità del lavoro e dell’argomento in relazione alla reale applicabilità dei risultati. ✓ Il premio non è cumulabile con altri premi, borse di studio, assegni. ✓ In caso di rinuncia da parte del vincitore, il premio verrà assegnato al candidato che segue in graduatoria. ✓ La cerimonia di premiazione avrà luogo in occasione del Congresso Tecnico di Assofond che si terrà al Museo 1000 Miglia a Brescia il 15 e 16 novembre 2018. ✓ Verrà data inoltre la possibilità ai vincitori di illustrare brevemente il proprio lavoro.

Experts’ corner Raffmetal: un modello di economia circolare nell’industria dell’alluminio a cura di Roberta Niboli

Roberta Niboli CEO di Raffmetal e Vice Presidente dell’Associazione European Aluminium

affmetal, con una produzione annuale di 250.000 tonnellate di leghe di alluminio da rifusione provenienti da materiali di recupero quali i rottami e 3 insediamenti produttivi con una superficie di 145.000 m3 di cui 90.000 coperti, continua a dimostrare il suo ruolo di leadership come

principale produttore europeo. L’azienda, fondata da Silvestro Niboli nel 1979, fa parte di Silmar Group che impiega oltre 2.800 dipendenti con un fatturato di oltre 820 milioni di € nel 2016. La Mission – La nostra Mission è produrre alluminio recuperando beni che

hanno esaurito il loro ciclo di vita ri-dando loro un nuovo valore. Ci impegniamo a operare riducendo al minimo i consumi energetici e promuoviamo processi circolari nel rispetto dell’ambiente, delle persone e del territorio a favore dello sviluppo sostenibile.

La Metallurgia Italiana - n. 2 2018

Scenari L’approccio produttivo circolare rappresenta oggi la miglior risposta che possiamo dare alle esigenze ambientali, sociali ed economiche di un territorio, in quanto rappresenta la più efficace alternativa al modello economico lineare che considera l’ambiente solo come generatore di risorse, le quali vengono trasformate, utilizzate e poi gettate. Al contrario, in

un sistema economico circolare ciò che di solito è considerato come “scarto” assume il ruolo di risorsa, da riutilizzare e valorizzare, ridefinendone quindi ruolo e significato all’interno del ciclo di vita dei prodotti. L’alluminio rappresenta il materiale del futuro in quanto permette di essere riciclato infinite volte e al 100% dando vita

ogni volta a nuovi prodotti. Le potenzialità di questo straordinario metallo sono cresciute negli anni trovando impiego in numerosi settori quali il packaging, l’edilizia, gli elettrodomestici, l’illuminazione e in particolar modo il settore automotive e dei trasporti.

L’industria dell’alluminio in Europa rappresenta perciò una realtà molto importante ed oggi anche un esempio di eccellenza per l’economia circolare. In un mondo in cui continua ad aumentare la richiesta di risorse naturali e soprattutto in un continente, come l’Europa, dove vi è scarsità di tali risorse diventa prioritario passare da un’economia lineare ad un’economia circolare, che prevede la progettazione di un sistema virtuoso di recupero e riutilizzo delle risorse. L’alluminio, riciclabile infinite volte, dà vita ad un loop virtuoso che dalla materia prima, passando attraverso la progettazione e la realizzazione di manufatti, la loro distribuzione e impiego, si conclude poi con un rifiuto che viene recuperato e riutilizzato dando vita a nuova materia prima.

Adottare un modello di economia circolare non può che garantire un futuro migliore all’economia europea, garantendone maggiore indipendenza e riducendo l’importazione di materie prime. Tale modello sostiene inoltre competitività ed efficienza delle imprese, impegnate nel più grande obiettivo di reindustrializzazione dell’Europa attraverso uno sviluppo sostenibile che tiene conto della storia e delle radici del territorio nel quale si opera. Raffmetal è orgogliosa di essere espressione, tramite il suo agire quotidiano, di un approccio circolare, che diventa parte del grande progetto e percorso di sostenibilità aziendale. Per Raffmetal sostenibilità significa non solo l’adozione di un modello economico circolare, ma ricerca avanzata, sviluppo di processi efficienti,

una politica ambientale evoluta, salute, sicurezza e formazione per la crescita dei collaboratori e del territorio. Gli ambiti in cui si articolano i nostri obiettivi in un rapporto sinergico e sistemico di sostenibilità sono quelli di sostenibilità economica, ambientale e sociale. Con il primo punto intendiamo il costante miglioramento dell’efficienza produttiva e dell’innovazione tecnologica. Dal punto di vista ambientale il nostro obiettivo è la realizzazione di un prodotto ecosostenibile, con la massima efficienza energetica e sempre maggiore integrazione ambientale. Con sostenibilità sociale, infine, intendiamo contribuire attivamente a creare valore, progresso e benessere al territorio e alla comunità a cui apparteniamo.

La Metallurgia Italiana - n. 2 2018

Aim news Calendario degli eventi internazionali International events calendar

QUOTE SOCIALI AIM 2018 (ANNO SOLARE) Benemeriti (quota minima) 1.750,00 €

March 11-15, Phoenix, USA TMS 2018 147th Annual Meeting & Exhibition April 12-13, Friedrichshafen, Germany European Conference on Heat Treatment - Nitriding and Nitrocarburizing April 30 - May 3, Houston, USA Offshore Technology Conference (OTC) 2018 May 19-26, Perth, Australia ALTA 2018 Nickel-Cobalt-Copper, Uranium-REE-Li and Gold-PM Conference & Exhibition

Sostenitori (quota minima)

750,00 €

Ordinari (solo persona)

70,00 €

Seniores

25,00 €

Juniores

15,00 €

La quota dà diritto di ricevere la rivista dell’Associazione, La Metallurgia Italiana (distribuita in formato digitale). Ai Soci viene riservato un prezzo speciale per la

May 23-25, Brno, Czech Republic 27th International Conference on Metallurgy and Materials, METAL 2018

partecipazione alle manifestazioni

May 30-31, Haifa, Israel Technological Innovation in Metals Engineering (TIME 2018)

pubblicazioni edite da AIM.

June 3-6, Pittsburgh, USA 2018 Superalloy 718 & Derivates: Energy, Aerospace, and Industrial Applications

Per ulteriori informazioni, iscrizioni, rinnovi:

June 5-7, Erfurt, Germany International Trade Show + Conference for Additive Manufacturing

20121 Milano

June 5-7, Greenville, USA 4th International Conference on HTSE in Automotive Applications June 10-13, Helsingør, Denmark 4th International Congress on 3D Materials Science (3DMS 2018)

AIM e per l’acquisto delle

AIM, Via F. Turati 8 Tel.: 02 76021132/76397770, fax: 02 76020551 e-mail: amm.aim@aimnet.it www.aimnet.it

June 18-21, Gaithersburg, USA Additive Manufaturing Benchmarks 2018 (AM-Bench 2018) July 8-13, Paris, France THERMEC July 15-21, Paris, France International Conference on Composites or Nano Engineering (ICCE-26) Semptember 5-8, Las Vegas, USA MEI2018 (Mining Expo International) September 9-13, Oxford, United Kingdom Eurosuperalloys 2018 September 12-14, Xi’ An, China 25TH IFHTSE September 13-14, Aachen, Germany Metallurgie im Wandel 4.0 October 14-18, Seattle, USA Furnace Tapping 2018 Conference October 16-19, Stockholm, Sweden 3rd Ingot Casting, Rolling and Forging Conference, ICRF 2018 60

La Metallurgia Italiana - n. 2 2018

Le Rubriche - Centri di studio Attività dei Comitati Tecnici CENTRO MATERIALI PER L’ENERGIA (ME) (riunione del C.T. – 16 novembre 2017) Consuntivo di attività svolte - Le GdS “Life assessment e vita residua di componenti e impianti operanti ad alta temperatura” (Milano, 18 e 19 ottobre 2017) hanno registrato la presenza di quasi 60 partecipanti, provenienti da industria e produzione, che hanno dato luogo ad una discussione tecnica vivace e di buon livello. Il presidente Gavelli esprime ai coordinatori Gualco e Merkling il suo apprezzamento per il risultato. Alcuni temi suggeriti dai partecipanti nei questionari potrebbero essere spunto per nuove iniziative.

per avere il tempo di contattare alcune istituzioni potenzialmente interessate all’evento ed eventualmente aprire la giornata agli sponsor sotto diverse forme. Iniziative future - I temi per possibili manifestazioni sono relativi a stampabilità a freddo di materiali metallici innovativi (TWITTRIP), all’influenza dei processi di saldatura sulle strutture metallurgiche e sulla qualità dei giunti saldati ed ad una eventuale giornata sul titanio e sue leghe CENTRO FORGIATURA (F) (riunione del C.T. – 13 dicembre 2017)

Iniziative future

- La GdS sui materiali per l’eolico dovrà essere organizzata con i produttori di componentistica, non essendoci in Italia costruttori di impianti. Il presidente Gavelli proporrà l’argomento al CT Forgiatura. - È’ stata presentata la bozza di programma per la GdS “Leghe di nickel e superleghe”; resta da fissare la data a la sede. - La GdS “Valutazione dell’esercibilità dei materiali nei cicli combinati tradizionali e innovativi” si terrà a metà giugno 2018, e sono stati nominati due coordinatori con l’incarico di stendere il programma.

- Il presidente Rampinini presenta una bozza del programma della GdS relativa all’analisi del rischio. Viene confermata la presenza di un broker per dettagliare la gestione rischi attraverso tutte le fasi di produzione, trasporto ed utilizzo del forgiato e la presenza di esperti per la sicurezza IT di tutti i dati sensibili. La giornata può svilupparsi in forma di dibattito/relazione sull’analisi di un contratto di fornitura. Partendo quindi da un’analisi del rischio si potrà passare alla sua definizione e miglioramento. La gestione del rischio avverrà all’interno dei limiti definiti dalla ISO 9001:2015, che diverrà obbligatoria entro settembre 2018.

Stato dell’arte e notizie - Il relatore Di Gianfrancesco riferisce sullo stato dei lavori di ECCC (European Creep Collaborative Commitee), sulla possibile riorganizzazione dei gruppi di lavoro e sui contatti tra ECCC e ASME. - Il presidente Gavelli introduce un nuovo membro del CT che si presenta e viene accettato. CENTRO METALLI E TECNOLOGIE APPLICATIVE (MTA ) (riunione del C.T. – 21 novembre 2017) Manifestazioni in corso di organizzazione - La GdS “I metalli per l’edilizia sostenibile – Acciaio e rame – La certificazione dei fabbricati” vien ulteriormente spostata alla seconda metà di marzo La Metallurgia Italiana - n. 2 2018

Stato dell’arte e notizie - Il presidente Rampinini ribadisce il ruolo fondamentale di AIM per l’innovazione e la ricerca, anche nel campo della forgiatura, e individua una possibile sinergia con il CT Acciaieria, pur mantenendo comunque chiara l’identità del CT Forgiatori come entità specializzata nel “fare scuola” per la forgiatura e nel proporre attività e giornate non convenzionali relativamente agli aspetti organizzativi, di sicurezza e di bilancio del mondo industriale. CENTRO TRATTAMENTI TERMICI E METALLOGRAFIA (TTM) (riunione del C.T. – 14 dicembre 2017) Consuntivo di attività svolte

Ricerche Fiat la GdS “Failure analysis in campo automotive”. La manifestazione ha avuto un ottimo successo con circa 90 iscritti. I partecipanti hanno valutato la manifestazione con punteggi compresi tra buono ed ottimo. - Si è tenuta a Ivrea nei giorni 19, 25 e 26 ottobre 2017 la prima edizione del corso “Gli stampi dalla progettazione all’utilizzo”. Il presidente Petta si dichiara soddisfatto per la riuscita e per la partecipazione di circa 50 persone, mentre il coordinatore Rivolta commenta i questionari di soddisfazione con la maggior parte dei giudizi tra il buono e l’ottimo. Manifestazioni in corso di organizzazione - Il presidente Petta conferma che il tradizionale Convegno Nazionale Trattamenti Termici, a cadenza biennale, sarà organizzato a Venezia dal 13 al 15 giugno nel contesto del Convegno internazionale ICS 2018 dedicato alla produzione siderurgica; la prima giornata del Convegno Nazionale sarà pensata e calibrata per l’imprenditoria italiana dei trattamenti termici e si svolgerà in italiano. La dott.ssa Bassani, Segretario Generale AIM, attende proposte dai possibili relatori entro la scadenza del 31 gennaio 2018. Iniziative future - È prevista per fine marzo 2018 la GdS “Tempra e deformazioni”. Il coordinatore Petta intende dare ampio spazio alla tavola rotonda finale dal titolo: “I costi delle deformazioni”. A breve saranno definiti il programma dettagliato e la sede. - Il coordinatore Rivolta presenta la bozza del programma per la GdS “Materiali per lo stampaggio”, durante la quale sarà approfondita la tematica relativa alla progettazione. Data e sede sono in fase di definizione. - Il coordinatore Petta conferma che il programma del secondo corso “Metallurgia di base” ricalcherà sostanzialmente la prima edizione, con alcuni miglioramenti volti ad evitare sovrapposizioni. Il corso si terrà a maggio. Stato dell’arte e notizie - Il presidente presenta due nuovi membri del comitato e segnala l’interesse di un ulteriore nuovo membro.

- Si è svolta a Torino presso il Centro 61

Aim news AIM – UNSIDER Norme pubblicate e progetti in inchiesta (aggiornamento 31 gennaio 2018) NORME UNSIDER PUBBLICATE DA UNI NEL MESE DI GENNAIO 2018 UNI EN 10263-4:2018 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo Parte 4: Condizioni tecniche di fornitura degli acciai da bonifica UNI EN 10263-5:2018 Vergella, barre e filo di acciaio per ricalcatura ed estrusione a freddo - Parte 5: Condizioni tecniche di fornitura degli acciai inossidabili

EN ISO 3887:2018 Steels - Determination of the depth of decarburization (ISO 3887:2017) ISO 6507-3:2018 Metallic materials -- Vickers hardness test -- Part 3: Calibration of reference blocks ISO 6507-2:2018 Metallic materials -- Vickers hardness test -- Part 2: Verification and calibration of testing machines

UNI EN 12681-1:2018 Fonderia - Prove radiografiche - Parte 1: Tecnica delle pellicole

ISO 26203-1:2018 Metallic materials -- Tensile testing at high strain rates -- Part 1: Elastic-bartype systems

UNI 11240-1:2018 Acciaio per cemento armato - Giunzioni meccaniche per barre - Parte 1: Requisiti

ISO 10804:2018 Restrained joint systems for ductile iron pipelines -- Design rules and type testing

UNI 11240-2:2018 Acciaio per cemento armato - Giunzioni meccaniche per barre - Parte 2: Metodi di prova

ISO 15653:2018 Metallic materials -- Method of test for the determination of quasistatic fracture toughness of welds

NORME UNSIDER RITIRATE DA UNI NEL MESE DI GENNAIO 2018

ISO 6507-1:2018 Metallic materials -- Vickers hardness test -- Part 1: Test method

UNI 11240-1:2007 Acciaio per cemento armato - Giunzioni meccaniche per barre - Parte 1: Requisiti

PROGETTI UNSIDER MESSI ALLO STUDIO DAL CEN (STAGE 10.99) – GENNAIO 2018

Petroleum and natural gas industry - Pipeline transportation systems -Pipeline integrity management specification Part 2: Full-life cycle integritymanagement for offshore pipeline (ISO/DIS 19345-2:2017) prEN ISO 10426-3 Petroleum and natural gas industries Cements and materials for well cementing - Part 3: Testing of deepwater well cement formulations (ISO/DIS 104263:2018) prEN ISO 10426-4 Petroleum and natural gas industries Cements and materials for well cementing - Part 4: Preparation and testing of foamed cement slurries at atmospheric pressure (ISO/DIS 10426-4:2018) prEN 488 District heating pipes - Bonded single pipe systems for directly buried hot water networks - Factory made steel valve assembly for steel service pipes, polyurethane thermal insulation and a casing of polyethylene prEN 14419 District heating pipes - Bonded single and twin pipe systems for buried hot water networks - Surveillance systems

UNI 11240-2:2007 Acciaio per cemento armato - Giunzioni meccaniche per barre - Parte 2: Metodi di prova

prEN ISO 13520 Determination of ferrite content in austenitic stainless steel castings

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

UNI EN 12681:2006 Fonderia - Controllo mediante radiografia

PROGETTI UNSIDER IN INCHIESTA PREN E ISO/DIS – FEBBRAIO 2018

EN 13480-1:2017/prA1:2017 Metallic industrial piping - Part 1: General

UNI EN 10263-4:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura degli acciai da bonifica UNI EN 10263-5:2003 Vergella, barre e filo di acciaio per ricalcatura a freddo ed estrusione a freddo - Condizioni tecniche di fornitura degli acciai inossidabili NORME UNSIDER PUBBLICATE DA CEN E ISO NEL MESE DI GENNAIO 2018 62

PREN – PROGETTI DI NORMA EUROPEI prEN ISO 19904-1 Petroleum and natural gas industries - Floating offshore structures - Part 1: Ship-shaped, semi-submersible, spar and shallow-draught cylindrical structures (ISO/DIS 19904-1:2017) prEN ISO 19900 Petroleum and natural gas industries - General requirements for offshore structures (ISO/DIS 19900:2018) prEN ISO 19345-2

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

La Metallurgia Italiana - n. 2 2018

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

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).