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European Conferenceon Heat Treatment 2015&22nd IFHTSE CongressHeat Treatment and Surface EngineeringFrom tradition to innovation

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Programme

With about 160 papers in the programme, we expect ECHT 2015 & 22 IFTHSE in Venice to be a vibrant and successful event. The 90 oral presentations and over 70 posters will give life to parallel sessions covering:

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Exhibition & SponsorshipThe Conference will feature a technical exhibition that will represent many areas of industry with latest equip-ment, facilities and instruments, products and services in the field of heat treatment and surface engineering. Companies will be able to reinforce their participation and enhance their corporate identification by taking advantage of benefits offered to them as Contributing Sponsors of the Conference.Companies interested in taking part in the exhi-bition or sponsoring the Conference may contact the ( · tel. +39 02 76021132).

Conference venueThe Conference will be held in Venice, at

(Viale Ancona 2, Venezia Mestre, Venice - Italy).

Conference organisersASSOCIAZIONE ITALIANA DI METALLURGIAp.le Rodolfo Morandi 2 · 20121 Milano · Italyphone: +39 02.7602.1132 or +39 02.7639.7770fax +39 02.7602.0551 · e-mail: [email protected] conference website: www.aimnet.it/ht2015.htm

ScopeThe long-standing co-operation between AIM - Associazione Italiana di Metallurgia - and the IFHTSE - Inter-national Federation for Heat Treatment and Surface Engineering - has already led to the joint organisation of significant events, such as the Congresses held in Florence in 1982 and in 1998 and the Congress on automotive applications organized in Riva del Garda in 2005.Now, our combined resources will be focused on the organization of a Conference that will deal with tradition and innovation in heat treatment and surface engineering. The Conference will join the European Conference on Heat Treatment 2015 and the 22nd IFHTSE Congress.

Conference website: www.aimnet.it/ht2015.htm

VENICE ITALY, 2022 MAY 2015

La Metallurgia Italiana

Direttore Responsabile:Gianangelo Camona

Comitato scientifico - Editorial Panel:Livio Battezzati, Riccardo Carli, Mario Conserva, Augusto Di Gianfrancesco, Franco Dinucci, Carla Gambaro, Gian Luca

Garagnani, Bevis Hutchinson, Chong Soo Lee, Alberto Molinari, Roberto Montanari, Elena Pereloma, Giorgio Poli, Emilio

Ramous, Claudia Rinaldi, Roberto Roberti, Hans J. Roven, Dieter Senk, Piotr R. Scheller, Pierre Soulignac, Jean-Marc Steiler, Stefano Trasatti, George F. Vander Voort, Maurizio Vedani

Segreteria di redazione:Antonella Donzelli

Comitato di redazione:Federica Bassani, Gianangelo Camona,

Antonella Donzelli, Ottavio Lecis, Carlo Mapelli

Direzione e redazione:AIM - P.le R. Morandi 2 - 20121 Milano

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Controllo e caratterizzazione

Comparative analysis on phase quantification methods in duplex stainless steels weldmentsM. Breda, J. Basoni, F. Toldo, C. Bastianello, S. A. Ontiveros Vidal, I. Calliari ....................................................................3

Storia della metallurgia

Fakes in African art: study of a reliquary figure (Mbulu-Ngulu) from GabonC. Soffritti, E. Fabbri, A. Fortini, M. Merlin, G.L. Garagnani ..9

Trattamenti termici

Effetti della diluizione sulla microstruttura e comportamento ad usura di una lega Fe-C-B-Cr-MR. Giovanardi, G. Poli, P. Veronesi, G. Parigi, N. Raffaelli ........15

Nanomateriali

Validity of Wulff construction used for size-dependent melting point of nanoparticlesS. Zhang, L. Zhang, L. Chen .................................................. 25

Acciai

Microstructural characterization and production of high yield strength rebarE. Mansutti, G. Luvarà, C. Fabbro, N. Redolfi ...................29

Metalli non ferrosi

On the ageing of a hyper-eutectic Zn-Al alloyA. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano ...........37

Forgiatura

Implementation of an open-die forging process for large hollow shafts for wind power plants with respect to an optimized microstructureM. Wolfgarten, D. Rosenstock, L. Schaeffer, G. Hirt .........43

International Journal of the Italian Association for Metallurgy

Mensile dell’Associazione Italiana di Metallurgia fondata nel 1946

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CONSED T V a e uropa Un ta 29 34073 G ado (GO) el 0431 876070 Fax 0431 88 507 www consed t com in o@conse it com

La Metallurgia Italiana O gano uf cia e dell Assoc az one I al ana di Me al urg aRiv sta fondata nel 1909

International Journal of the Italian Association for Meta lurgy

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N 4 ap i e 2015 - Anno 107

N. 4/Aprile 2015

Anno 107 - ISSN 0026-0843

THE HEARTOF HEAT

QualityQuality

SafetySafety

ReliabilityReliability

Efficiency

www.cieffe-forni.com - [email protected]

MICHAEL JORDAN

La Metallurgia Italiana - n. 4/2015 3


Comparative analysis on phase quantification methods in duplex stainless steels weldments

M. Breda, J. Basoni, F. Toldo, C. Bastianello, S. A. Ontiveros Vidal, I. Calliari

Duplex Stainless Steels (DSS) are biphasic steels of increasing interest, employment as structural materials in aggressive environments. In these steels, the austenite-to-ferrite phase ratio is maintained at about one – even

if a slightly wider range between 40/60 and 60/40 is in any case accepted – giving the best combination of mechanical and corrosion-resistance properties. However, DSS must be handled with extreme care, especially

if thermal cycles are involved, owing to the possible formation of dangerous secondary compounds that can highly worsen their excellent features. In industry, the production of big pipes requires manufacturing welding

operations on steel plates or sheets and the end products must satisfy specific requirements. Therefore, since DSS properties depend on phase ratio, ferrite quantification at an industrial scale represents a topic of great interest, which must be as reliable as possible and, at the same time, of fast execution. In the present

paper, different methods currently employed for ferrite estimation in DSS weldments are compared, in order to understand the limits deriving from each technique.

Keywords: Stainless Steel - Welding - Metallography - Material analysis

J. BasoniDe Pretto Industrie S.r.l.,

Via Fogazzaro 5, 36015 Schio (VI) - Italy

M. Breda, I. CalliariIndustrial Engineering Department (DII),

University of PadovaVia Gradenigo 6A, 35131 Padova – Italy

F. Toldo, C. BastianelloLaboratorio Prove Materiali San Marco

Via Lago di Alleghe, 30/32 - 36015 Schio (VI)- Italy

S. A. Ontiveros VidalInstituto Tecnológico de Saltillo,

Venustiano Carranza 24000, Tecnologico,25280 Saltillo, Coahuila de Zaragoza – Mexico

INTRODUCTION

In the biphasic austeno-ferritic stainless steels, commonly named Duplex (DSS), the presence of equal volume frac-tions of the phases provides an excellent combination of mechanical properties and corrosion resistance, especially when compared with conventional stainless steels grades [1,2]. In Off-Shore engineering, their usage permits the design of components having smaller thicknesses – and therefore lighter – without compromising the corrosion re-sistance and avoiding the employment of expensive anti-

corrosion coatings. In DSS, the balanced phase ratio can be obtained through an appropriate solution-annealing tre-atment in the temperature range 1050–1100°C followed by water quenching and, even if the 50/50 ratio is the de-sired one, phase amounts ranging from 40/60 to 60/40 are in any case accepted.However, if subjected to improper thermal cycles and especially in the temperature range 800–950°C, DSS are sensitive to secondary phases precipitation (intermetalli-cs, carbides, nitrides), which can determine drastic losses in their advantageous properties [1-8]. Therefore, the solu-tion treatment is usually performed after the forming ope-rations and, besides ensuring the achievement of the de-sired Duplex microstructure, it permits to re-dissolve any dangerous precipitate formed during the manufacturing cycle. However, weldments are not always treatable after components joining, especially when in-service operations are performed, and special care must be adopted for wel-ding purposes. In this regard, DSS must be managed as austenitic grades but using dedicated devices, avoiding the formation of undesired structures and considering that the solution-annealing treatment could not be performed when big parts are joined. After welding, DSS are required to be free from secondary phases and the austenite/ferrite volume fractions must be maintained within the desired forks, in order to guarantee the expected corrosion resistance and mechanical proper-ties. Therefore, the qualification tests require a systematic measurement of the ferrite percentage (%FE) in different parts of the joint – base material (BM), heat affected zone (HAZ) and fusion zone (FZ). The international reference

La Metallurgia Italiana - n. 4/20154

Memorie

standards assess how to evaluate %FE, but customers often necessitate dedicated procedures by adapting the standardized approach to their specific requirements. The most widely accepted international standard for phases quantification is Optical Microscopy (OM) after polishing and chemical etching, by following the provisions of ASTM E562-11 standards (manual point count), which provides information on selecting type of pattern and number of fields to be analyzed that ensure a well-defined relative accuracy. On the other hand, points counting through au-tomatic image analysis is otherwise possible, in accordan-ce with the provisions of ASTM E1245-03 standards (auto-matic image analysis). In both cases, the use of Scanning Electron Microscope (SEM) instead of OM is permitted.Phase quantification by image analysis on OM and SEM micrographs is often replaced by other simple – and faster – field-methods, among which the use of ferritoscope is the most popular. This technique is based on measuring the magnetic field generated by an induced-currents pro-be, by placing it in contact with the metal surface; since ferrite is a magnetic phase while austenite is amagnetic, this device provides %FE by measuring the magnetic re-sponse of the material. However, this method is very sen-sitive to the finishing of the contact surface and cannot be applied near edges or corners, owing to the distortion of the magnetic field. Moreover, the use of ferritoscope is li-mited to the investigation of wide areas of welds and base material, whereas HAZ is not readily controllable, due to its small size. In this method, results are given in units cal-led Ferrite Number (FN) and are automatically converted to %FE through an internal correlation.Another method for %FE determination is based on calcu-lations from phase diagrams, according to the AWS 5.4 assessments, by knowing the Cr equivalent (Creq) and Ni equivalent (Nieq) values of both base and filler materials derived from chemical analyses. A disadvantage of this method lies on the low accuracy level of the obtained va-lues, since the lines on diagrams are drawn only for some reference values and all points in between are evaluated graphically, implying many difficulties and a dependence on the adopted interpolation method. Moreover, if on one side BM and FZ compositions can be easily determined, the chemical analysis on HAZ cannot be executed, owing again to its small size.In the present paper, a comparison of the previously de-scribed method for %FE evaluation on SAF 2205 DSS wel-ded joints is reported, in terms of relative reliability and associated accuracy.

EXPERIMENTAL

The comparative study was carried out on two full-pene-tration qualification beads (Bead-1 and Bead-2), obtained by joining two types of SAF 2205 DSS (UNS S31803) base materials and adopting the same welding procedure. The base materials (compositions in Table 1) were of different manufacture, produced by traditional forging (forged ma-terial) and Hot Isostatic Pressure (HIP material), whereas the welding procedure involved the Gas Tungsten Arc Wel-ding (GTAW) process for the first passes (about 6 mm of deposited material), subsequently filled using the Shielded Metal Arc Welding (SMAW) technique.The qualification beads were prepared for metallographic investigation on a Leica DMRE OM by mechanical polishing and electrolytic etching at 5V using a solution composed of 30% NaOH in deionized water. In each sample, diffe-rent zones of the weldments were distinguished: BM (at about half thickness), upper HAZ (cover layers), lower HAZ (root layers), upper FZ and lower FZ. For the comparati-ve analysis, %FE was estimated taking into account four different methods: ASTM E562 manual counting points method using fixed parameters (magnification, grid and number of fields), ASTM E562 manual counting points me-thod using variable parameters, automatic image analysis (ASTM E562 and ASTM E1245) and manual image analysis (variable magnification). For the first three OM methods, a 500x magnification for all the investigated areas was maintained.The same zones of the weldments were observed using a Leica Cambridge Stereoscan 440 SEM operating in back-scattered-electron mode (BSE) at 29 kV; as it is known, the SEM-BSE observation allows distinguishing the micro-structural constituents according to their average atomic number and, therefore, ferrite (the lighter phase) appears darker than austenite (the heavier phase). In this case, the samples were slightly etched, in order to only better define phase boundaries, while the magnification was set as va-riable, according to the observed microstructure. For fer-rite quantification, the micrographs were edited using an image-analysis software by applying proper filters to im-prove phase contrasts and minimize grayscale threshold errors (the filtering procedure was set-up for the specific case).For the aim of the present work, %FE was also measured using a portable Fischer MP30 Ferritscope in the previou-sly defined areas, and calculated using phase diagrams (Schaeffler-Delong, ESPY and WRC-1992), after the eva-luation of the local chemical composition and considering the Creq and Nieq values.

C Si Mn Cr Ni Mo Cu P S N

forged 0.027 0.22 1.12 22.2 5.9 3.3 0.16 0.023 0.002 0.16

HIP 0.021 0.71 1.13 22.4 5.2 3.1 0.17 0.018 0.003 0.17

Table 1 – Chemical composition of the base materials [wt.%].


Fig. 1 – OM images of etched microstructures in the investigated DSS weldments (Bead-1):forged BM-HAZ (left), FZ (middle), HIP HAZ-BM (right) .

RESULTS AND DISCUSSION

In both the qualification beads, the microstructures were similar; the BMs were free from intermetallics and the wel-ded zones exhibited the classical dendritic morphology achieved from melt solidification (Figs. 1 and 2). As expec-ted, the welding processes (GTAW and SMAW) determined some differences in the final FZ microstructure: SMAW, owing to its lower heat input, caused the formation of a finer dendritic structure in the filling passes, but an increa-sed micro-porosity level respect to GTAW was obtained.The results of the investigation are listed in Table 2, where the values relative to Bead-1 are reported over those con-cerning Bead-2. Starting from the OM quantification, it is possible to note that increasing the number of grid points, the accuracy of the evaluation also increased, as the stan-dard deviation was reduced; however, a major number of points lead to an increase in quantification time spent by the operator, which is not suitable at an industrial scale. The variable magnification method did not provide impro-vements on phase quantification, and the results were si-milar to those obtained using the manual method with a low number of grid points. On the contrary, the automatic image analysis software provided the more accurate re-sults, even using a smaller number of fields, because the system is able to automatically delete the “problems” con-cerning the choice of the suitable grating; however, in this case, the evaluation procedure must be properly set-up.

As can be seen from the table, the middle part of FZ was not always taken into account, since it is a transient region and requires particular assessments. Concerning the OM automatic method, the micrographs can be easily edited using image-analysis, because the employed reagent dar-kens ferrite and leaves austenite unaffected; therefore, the images appear in a grayscale pointing toward a black-and-white image, thus having a net phase contrast, allowing for a simple determination of ferrite volume fraction and easing phases discrimination.The analysis of SEM images showed substantial differen-ces respect to OM manual quantification, especially in the base materials, and ferrite was mainly underestimated. In this regard, SEM introduces a further source of uncertain-ty, since the signal from the BSE detector is processed by assigning different levels of grey to the detected ener-gies. Therefore, the micrographs are no more like black-and-white images, but an extended greyscale image is created, and the threshold defining the boundary between the phases cannot be univocally defined at all (Fig. 2). In this case, the operator plays a key role in the accuracy of the estimated phase amount, because the assignation of the threshold became strongly subjective. In addition, austenite and ferrite compositions in FZ are conditioned by the welding processes, since temperature and cooling rate affect elements partitioning. Therefore, changes in average atomic number can occur, leading to difference



Memorie

in the observed microstructure. In the present study, the quantification was slightly facilitated by the preliminary electrochemical etching, but the introduction of a large amount grayscales unavoidably leaded to an increase in estimation errors.Ferritoscope provided values that sometimes coincided and sometimes differed from those obtained through image methods; this was mainly due to the extent of the analysed volume (about 10 mm3), greater than that invol-ved using other techniques, and also to the amplification of the uncertainty in the conversion from FN to %FE. Final-ly, the values calculated from phase diagrams showed the greatest variability, which is related to the chosen method.

Method%FE (st.dev.)

Forged BM Upper HAZ Lower HAZ Upper FZ Middle FZ Lower FZ HIP BM

manual OM 16 pt. grid

(500x)

52 (12)51 (7)

60 (13)58 (8)

55 (9)57 (6)

37 (9)41 (10)

-30 (6)30 (9)

-

manual OM 96 pt. grid

(500x)

52 (5)50 (6)

61 (3)62 (4)

57 (5)58 (5)

32 (3)37 (4)

-30 (3)31 (3)

-

automatic OM

(500x)

54 (1)52 (1)

59 (1)59 (2)

60 (1)56 (2)

34 (2)38 (1)

55 (4)-

32 (2)32 (1)

33 (9)-

manual OM (var. magn.)

50 (7)56 (10)

- -36 (3)43 (5)

46 (3)37(6)

29 (4)34 (8)

56 (2)57 (2)

SEM(var. magn.)

44 (6)47 (7)

58 (7)64 (8)

59 (8)65 (9)

38 (4)43 (6)

40 (4)36 (9)

26 (5)32 (7)

49 (3)52 (2)

Ferritoscope 53 (1)58 (1)

- -37 (1)40 (2)

-32 (1)26 (1)

-

*Schaeffler 49 - -38 34

- - 49

ESPY 82 - -2825

- - 89

*WRC-1992 84 - -5651

- - 98

* Value in FN

Table 2 – Comparison of ferrite quantification methods (Bead-1 over Bead-2).

Fig. 2 – SEM images of etched microstructures (Bead-1): forged HAZ (left), upper FZ (middle), lower FZ (right).

Except for the Schaeffler-Delong diagram, the results were highly misaligned to that of OM, SEM and ferritoscope; this is intrinsically due to such diagrams, which have been developed to perform calculations on areas having a high concentrations of alloying elements, more similar to that of filler materials rather than the welded ones.

CONCLUSIONS

In the present paper, different methods currently emplo-yed for ferrite quantification in DSS weldments were com-pared, in order to understand the limits related to each technique. Analyses involving OM, SEM, ferritoscope and


phase diagrams were considered, and the related ferrite estimations were presented.From the results, it was revealed that OM is preferable for ferrite quantification if compared to SEM, both in base ma-terial and in the welded zone. In manual counting, several grid points must be adopted for a proper phase estimation and the uncertainty can be strongly reduced, but the in-creased counting time makes this method not readily ap-plicable when fast estimation are required. On the other hand, the automatic method provided the best results, but the procedure must be properly set-up; in this case, OM images are nearly black-and-white and, therefore, the subsequent analysis can be easily performed, since the th-reshold values are better defined and the “problems” con-cerning the choice of the suitable grating are automatically deleted by the program. On the contrary, SEM microgra-phs are grayscale images and phases quantification highly suffers from the assigned thresholds by the operator to phase boundaries. Moreover, since phases compositions are conditioned by the previous thermo-mechanical pro-cesses, difference in the observed microstructure may oc-cur by adopting the SEM-BSE observation method, owing to changes in the phase average atomic number.Finally, the investigation confirmed the low accuracy of fields-methods such as those deriving from phase dia-grams, whereas ferritoscope can be considered as an intermediate-accuracy technique, even if it requires large volumes to be measured. These methods, although simple and fast, cannot always assure a reliable ferrite quantifica-tion, owing to uncertainties intrinsic to the methods itself.

REFERENCES

[1] J. O. Nilsson. Super Duplex Stainless Steels. Mater Sci Tech 8 (1992), p. 685.

[2] R. N. Gunn. Duplex Stainless Steels: Microstructure, Properties and Applications. Abington Publishing, Cambridge, England (1997).

[3] J.O. Nilsson, A. Wilson. Influence of isothermal phase transformations on toughness and pitting corrosion of super duplex stainless steel SAF 2507. Mater Sci Tech 9 (1993), p. 545.

[4] I. Calliari, M. Zanesco, E. Ramous. Influence of isother-mal aging on secondary phases precipitation and tou-ghness of a duplex stainless steel SAF 2205. J Mater Sci 41 (2006), p. 7643.

[5] I. Calliari, G. Straffelini and E. Ramous. Investigation of secondary phase effect on 2205 DSS fracture tough-ness. Mater Sci Tech 26 (2010), p. 81.

[6] I. Calliari, M. Pellizzari and E. Ramous. Precipitation of secondary phases in super duplex stainless steel Ze-ron100 isothermally aged. Mater Sci Tech 27 (2011), p. 928.

[7] I. Calliari, M.Breda, E. Ramous, K. Brunelli, M.Pizzo, C. Menapace. Impact Toughness of an Isothermally Trea-ted Zeron®100 SDSS. J Mater Eng Perform 21 (2012), p. 2117.

[8] M. Pohl, O. Storz, Thomas Glogowski. Effect of inter-metallic precipitations on the properties of duplex stainless steel. Mater Charact 58 (2007), p. 65.


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Fakes in African art: study of a reliquary figure (Mbulu-Ngulu) from Gabon

C. Soffritti, E. Fabbri, A. Fortini, M. Merlin, G.L. Garagnani

The aim of the present work is the chemical and microstructural characterisation of a reliquary figure, stylistically consistent with the art of the Kota population, which lived in the eastern part of Gabon (Africa). The artefact was

subjected to preliminary observation by stereomicroscopy, and then Optical Microscopy (OM) and Scanning Electron Microscopy (SEM) analyses are carried out on a fragment and on surface compounds. Lastly, AMS

radiocarbon dating of the wooden support allowed further information about the production period to be obtained.The results show that the artefact was produced by a Cu-Zn alloy and contains non-metallic impurities made up of S and Se. The greenish and whitish surface compounds, which are mainly collected near the nails and in proximity to the overlaid sheets, are probably only partly related to natural corrosive processes. Finally, radiocarbon dating

established that the wooden support certainly dates after 1950.

Keywords: Copper and alloys – Material characterisation – Metallography – Electron microscopy – Metallurgy

C. SoffrittiTekneHub, Department of Architecture, University of

Ferrara, Via Quartieri 8, 44121 Ferrara, Italy

C. Soffritti, E. Fabbri, A. Fortini, M. Merlin,G.L. Garagnani

Department of Engineering - “A. Daccò” Corrosion and Metallurgy Study Centre, University of Ferrara,

Via Saragat 1, 44122 Ferrara, Italy

Storia della Metallurgia

INTRODUCTION

The proliferation of fakes in African arts has grown enor-mously in recent years, with a particular explosion since the 1950s, due to an increase in demand by collectors, which created new fields of activities for African foundri-es. In fact, in the 1980s the quantity of antiquities on sale increased further and today many replicas of tourist sou-venirs and fanciful copies of traditional forms enrich the art market.The official definition of authenticity for African artefacts consists of two inseparable conditions: any object created for a traditional purpose and by a traditional artist may be considered authentic [1].It is rather difficult to determine if an African artefact is original or a copy because literature is characterised by incomplete information about the African arts and the pro-duction of artefacts by artists [2]. The studies of African artefacts are somewhat incomplete since there is no cor-

relation between the style used in these works, the mate-rials used to produce them and the geological context of the extraction zone.Archaeometric analyses are essential to determine the state of conservation of the objects as well as to evaluate the production period in order to establish the authenticity of the artefacts.The aim of the present work is the characterisation of a sculpture, stylistically consistent with the art of the Kota population, which lived in the eastern part of Gabon (Afri-ca). This community is known for the realisation of metal-lic reliquary figures, which were set on wooden supports and called Mbulu-Ngulu or Bwéte. It should be noted that the first samples of these sculptures arrived in France and Germany during the last quarter of the 19th century. Reli-quary artefacts should, however, be more ancient given that the local copper mines had already been exploited to obtain metal for the coating of artefacts [3].Unfortunately, the majority of researchers have relied so-lely on stylistic analyses of the ornaments, which decorate the surface of the objects. More extensive investigations on the chemical composition of the alloy of the artefacts, in conjunction with a systematic characterisation of origi-nal African metallic objects would allow the evaluation of the provenance and the dating of Kota funeral art [4]. The present paper focuses on the investigation of the symbolic representation of a human abstract figure whose head is bigger than the rest of the body. These abstract figures were used to protect and demarcate the bones of family ancestors, which were preserved in containers made of bark. The artefact consists of a carved piece of wood (42 cm in height, 23 cm wide and 2 cm thick) covered on one side with metal sheets, which were fixed onto the support with


Memorie

small metallic nails. These metal sheets are very thin in order to fix almost perfectly to the carved wood. The fol-lowing morphological elements of the sculpture are detec-table in Fig 1a-c:

• The oval face has stylised eyes and nose but the mouth is not depicted. Two metallic plates are nailed onto the surface to represent a cross. In agreement with E.W. Herbert [5] this element has been found starting from the end of the 15th century as a result of the Congolese population’s conversion to Catholicism;

• Two lateral parts at ear-level which are often conside-red the representation of a hat;

• Two cylindrical pendants placed on the base of the la-teral parts, which are the abstract representation of traditional male and female hairstyles;

• One half-moon shaped sheet is located above the oval face and harmoniously integrated with the lateral parts;

• One rhomboid element that symbolises the body and the legs, placed on a rectangular wooden base;

• Another rhomboid element, on the back of the sculptu-re, stretched along the vertical axis and with a protru-ding “vein”.

The goal of the present work is the chemical and micro-structural characterisation of the sculpture as well as of the products located on its surface. The analyses were carried out by stereomicroscopy, Optical Microscopy (OM) and Scanning Electron Microscopy (SEM) coupled with Energy Dispersion Spectroscopy (EDS). Finally, a wooden fragment of the support was analysed using Accelerator Mass Spectrometry (AMS), which enabled radiocarbon (14C) dating of the artefact.

MATERIALS AND METHODS

The sculpture was observed by stereomicroscopy, equip-ped with a Moticam 2500 – 5 Mp camera, in order to ob-tain information on the manufacturing technique and to check the state of conservation. The investigations have revealed the presence of some compounds, which are

Fig. 1 – Macroscopic images of the sculpture: front (a), back (b) and side view(c).

Fig. 1 – Immagini fotografiche del manufatto: parte anteriore (a), parte posteriore (b) e visione laterale (c).

mainly concentrated near the nails and in proximity to the overlaid sheets.Thereafter, the evaluation of the alloy and the composition of the different colour surface compounds was carried out using a ZEISS EVO MA 15 Scanning Electron Microsco-pe (SEM), coupled with Energy Dispersion Spectroscopy (EDS).Moreover, a metal fragment of a few millimetres was taken from an unobtrusive area of the sheet. The sample was mounted in conductive resin, polished and submitted to conventional metallographic observation using LEICA ME-F4M Optical Microscopy (OM).Lastly, a sample of a few grams was collected from the base of the wooden support and subsequently was dated using Accelerator Mass Spectrometry (AMS) at Centro di Datazione e Diagnostica (CEDAD) – University of Salento.

RESULTS AND DISCUSSION

Macroscopic investigationsPreliminary macroscopic investigations have yielded a gre-at deal of information about the manufacturing technique as well as the nature of the products located on the sur-face. In agreement with some of the literature and private com-munications expressed in the last few years [4], the parts of the face that are not overlaid by the two metallic sheets (positioned in a cross) consist of a single plate. These are decorated with “lamellage”, a technique characterised by various equidistant streaks, which are placed in a slanting or horizontal pattern. These strips are also depicted on the half-moon sheet and on the rhomboid element that symbolises the body and the legs. In the latter two cases, a multitude of pitting embossed using a punch also deco-rates the surface of the plate [4]. Fig. 2 shows representative images of the rare compounds using stereomicroscopy, which were mainly collected near the nails and in proximity to the sheets. The colouring of these products is clearly green or whitish.


Fig. 2 – Macroscopic images of the different colour compounds that appear on the surface.

Fig. 2 – Macrografie rappresentative dei composti che si presentano con cromie differenti.

Radiocarbon measurementThe AMS radiocarbon dating established that the woo-den support is certainly dated after 1950. It is well known that the production of these sculptures ended around the 1930s because of the great number of Catholic missions, which imposed a new social organisation based on We-stern households [4]. Finally, it should be pointed out that the integrity as well as the total absence of signs due to a wooden support substitution is clearly evident.

Microstructural analysisAfter metallographic preparation, the microstructure of the metal fragment taken from the sheet was highlighted.

In particular, in Fig. 3a a shrinkage cavity of remarkable dimensions, formed during the alloy solidification, is vi-sible. Fig. 3b also shows the microstructure of the alloy after chemical etching by FeCl3/HCl. The presence of both non-homogeneous grain size and thermal twin bands would suggest that the artefact was obtained by alternate hammering and annealing steps. It should be noted that the variable grain size is probably due to a heterogeneous plastic deformation induced by manual hammering.

Chemical analysisFig. 4a shows a SEM image of the alloy together with the corresponding EDS spectrum. SEM-EDS analysis highlights that the artefact was produced by a Cu-Zn alloy, without the addition of alloying elements, i.e. Pb. No impurities (i.e. As, Fe, Sb), which are very common in the ancient alloys, were detected. It should be noted that, compara-ble amounts of Cu and Zn (Fig. 4b) could also be found in modern brasses such as the commercial “Yellow Brass” which contains 65 wt.% of Cu and 35 wt.% of Zn [7].Fig. 5a shows a SEM image of rare microscopic inclusions that are visible in the alloy. In particular, SEM-EDS analy-

Fig. 3 – Optical images of: (a) a detail of a shrinkage cavity on the polished surface; (b) the microstructure of grains with the presence of thermal twin bands on the chemically etched surface.

Fig. 3 - Micrografie OM del manufatto: cavità da ritiro (a), in assenza di attacco metallografico; microstruttura a grani e geminati (b); in presenza di attacco con reattivo a base di cloruro ferrico.

sis allows the verification of the presence of non-metallic impurities enriched with S and Se. To our knowledge, only one reference reports some South African (Lowveld) me-tallic artefacts [8], approximately dated from 1000 A.D. to 1980 A.D., which were characterised by many copper-iron sulphide inclusions containing up to 3% Se by weight, resi-dual from incomplete ore reduction.Over the centuries, the Kota reliquary figures went into stylistic decline and they were characterised by more ab-stract and grotesque meanings. Moreover, the demand for these artefacts from Western collectors has grown enormously in recent years, causing the proliferation of sculptures without any “funerary” meaning for the purpo-



Memorie

Element Weight % Atomic %

C K 12.17 40.42

S K 9.38 11.66

Cu K 36.53 22.92

Zn K 36.60 22.32

Se K 5.32 2.69

Totals 100.00

Fig. 4 – SEM backscattered electron image of the alloy indicated by pink square (Spectrum 1), together with the corresponding EDS spectrum; (b) average composition of the area in Fig. 4a (measured by EDS). The contents of Cu and Zn are highlighted in Fig. 4b

Fig. 4 – (a) immagine al SEM della matrice metallica con indicazione (Spectrum 1) della zona analizzata e corrispondente spettro EDS; (b) dati semi-quantitativi relativi allo spettro in Fig. 4a. In Tabella vengono evidenziati i contenuti di Cu e Zn.

Element Weight % Atomic %

C K 2.44 11.59

O K 0.57 2.03

Cu K 62.65 56.35

Zn K 34.35 30.03

Totals 100.00

Fig. 5 – SEM backscattered electron image of an inclusion, indicated by a black arrow (Spectrum 2), together with the corresponding EDS spectrum; (b) average composition of the analysed point in Fig. 5a (measured by EDS). The contents of S and Se are highlighted in Fig. 5b.

Fig. 5 – (a) immagine al SEM di una delle inclusioni visibili all’interno della matrice metallica con indicazione (Spectrum 2) della zona analizzata e corrispondente spettro EDS; (b) dati semi-quantitativi relativi allo spettro in Fig. 5a. In Tabella vengono evidenziati i contenuti di S e Se.

Fig. 7 – Representative SEM image of the morphology of surface whitish compounds, together with corresponding EDS spectrum.

Fig. 7 – Micrografia SEM dei composti di colore biancastro e relativo spettro EDS.


se of enriching the flourishing art market. Starting from the first decade of the 20th century, the practice of reca-sting damaged copper and brass to recover the precious metal was very common. In particular, E. Andersson [9] highlighted that many “Mbulu-Ngulu” were obtained by recasting ancient alloy which was later mounted in more recent wooden supports (second half of the 20th century). In this regard, because of the perfect realisation of the artefact without the addition of alloying elements or com-mon impurities, the rare Se inclusions and the dating of the wooden support (see § Radiocarbon measurement), it is possible to further suppose that the sculpture analysed in this paper was realised by the methods described by Andersson in [9].The SEM image of the greenish surface compounds toge-ther with the corresponding EDS spectrum is reported in Fig. 6. Because of the small amount of products on the surface, it was not possible to take samples and to car-ry out specific analyses like XRD or Raman spectroscopy. First of all, the morphologies in Fig. 6a and 6b are very dif-ferent. It should be noted that the needle-like or lamellar structure shown in Fig. 6a is frequently observed in cop-per carbonate compounds. This evidence is supported by SEM-EDS analysis. On the contrary, the same technique would suggest that the compounds of Fig. 6b are probably zinc oxychloride.Fig. 7 is a representative SEM image of whitish compoun-ds, which are mainly collected in proximity to the overlaid sheets. The EDS spectrum emphasises high concentra-tions of Cl and Pb. In particular, the latter element is totally absent in the alloy and it is possible that it is not produced by natural corrosive processes.

CONCLUDING REMARKS

The present work has proved the usefulness of an interdi-sciplinary approach to clarify some general aspects about the manufacturing process and the state of conservation of metal artefacts.Macroscopic examinations have highlighted a good state of conservation of the sculpture and a manufacturing pro-cess consistent with the reliquary Kota art. Observations by Optical Microscopy (OM) have establi-shed that the sculpture was obtained by a casting and was subsequently subjected to alternate hammering and anne-aling stages.SEM-EDS analysis has highlighted that the artefact was produced by a Cu-Zn alloy, with an amount of the latter elements comparable to those that could can be found in modern brasses (i.e. “Yellow Brasses”). The absence of alloying elements and the presence of rare Se inclusions bear witness to an advanced manufacturing process and this suggests that the artefact was obtained by a rather recent recast. Finally, the chemical analyses of greenish and whitish surface compounds lead to the assumption that they are only in part related to natural corrosive pro-cesses.

BIBLIOGRAPHY

[1] J. CORNET, Afr. Arts, Vol. 9 (1975), no. 1, pp. 52-55.[2] E. BASSANI, Rivista trimestrale di studi e documenta-

zione dell’Istituto italiano per l’Africa e l’Oriente (1980), no. 1 pp. 85-95.

[3] E. BASSANI, Arte Africana, Ed. Skira, Milano (2012).[4] G. DELORME, Arts d’Afrique Noire, (2002), no. 122, pp.

1-14.[5] E.W. HERBERT, Red gold of Africa - Copper in precolo-

nial history and culture, University of Wisconsin Press, Madison (1984).

[6] E. ANDERSSON, Contribution à l’ethnographie des Kuta, Ed. Almqvist & Wiksell, Stockholm (1953).

[7] ASM Specialty Handbook - Copper and Copper Alloys, ASM International, Materials Park, Ohio, (2001).

[8] D. MILLER, D. KILLICK, N.J. VAN DER MERWE, J. Field Archaeol., Vol. 28 (2001), pp. 401-417.

[9] E. ANDERSSON, Contribution to the ethnography of Kuta III, Occasional Papers XV, S. Lagercrantz and A. Loôv (1991).


6th International Conference

STEELSIM2015

BARDOLINO, ITALY23-25 SEPTEMBER 2015

Modelling and Simulation of metallurgical processes cover an important role in optimizing technological processes, decreasing production costs, increasing steel quality and defining the correct design of metallurgical processes in order to improve their sustainability even from the environmental point of view. The fundamentals of metallurgical processes can be investigated through physical and numerical modelling following several numerical approaches. Traditional and new mathematical techniques applied by modern simulation facilities allow to achieve results that are useful to understand physical interaction and to design a profitable metallurgical process. The simulations technique can be applied to the different steps of the metallurgical production route: production and refining of liquid metals, solidification, plastic deformation, thermo-mechanical processes, thermal treatment, verification of structural reliability etc.

CONFERENCE TOPICSState of art and developments in modeling and simulation in steelmaking:• Ironmaking• Primary metallurgy (aluminium alloys, copper alloys

titanium alloys etc.)• Secondary steelmaking• Refining of metal alloys• Thermodynamic and kinetic simulation of the metallurgical

systems • Casting and solidification• Electrochemical processes• Metalforming processes and thermo-mechanical treatment• Heat treatments• Fracture mechanics and safety criteria• Fatigue mechanics• Safety criteria• Reduction of environmental impact

EXHIBITION & SPONSORSHIP OPPORTUNITIESSteelSim 2015 will feature an Exhibition that will enable excellent exposure for company products, technologies, innovative solutions or services. Companies will also be able to become Sponsors of the Conference. Companies interested in taking part in the Exhibition or in sponsoring the event may contact the Organising Secretariat (e-mail: [email protected] / fax: +39 0276020551).

VENUEThe Conference will be staged at the Congress Center of Aqualux Hotel Spa Suite & Terme, in Bardolino (VR), ItalyVia Europa Unita, 24/b 37011

www.aimnet.it/steelsim2015.htm

ORGANIZING SECRETARIAT

AIM - ASSOCIAZIONE ITALIANA DI METALLURGIAP.le R. Morandi, 2 · 20121 Milano · Italy · tel. +39 02 76021132 · fax. +39 02 76020551

e-mail: [email protected] www.aimnet.it

Organised by ASSOCIAZIONE ITALIANA DI METALLURGIA

MODELLING and SIMULATION of METALLURGICAL PROCESSES in STEELMAKING


Trattamenti termici

Trattamenti termochimici di nitrurazione e post-ossidazione su acciai 17-4PH:

ottimizzazione dei parametri di processo per massimizzare la resistenza a corrosione

R. Giovanardi, G. Poli, P. Veronesi, G. Parigi, N. Raffaelli

L’acciaio inossidabile 17-4PH viene solitamente trattato termicamente per incrementarne le proprietà meccaniche. Per migliorare ulteriormente la resistenza ad usura di tale acciaio è possibile sottoporlo a trattamenti termochimici, quali ad esempio la nitrurazione. Trattandosi di un acciaio inossidabile, in grado di presentare allo stato di fornitura una notevole resistenza a corrosione conferita dall’elevato contenuto di cromo presente in lega, viene spontaneo chiedersi se e quanto i trattamenti termici e termochimici possano influenzare questa proprietà.

Lo scopo del presente lavoro è quello di valutare come l’applicazione di trattamenti termici e termochimici, solitamente impiegati per incrementare proprietà meccaniche e anti-usura, influiscano sulla resistenza a

corrosione dell’acciaio e di intervenire sulle variabili di processo degli stessi (oppure mediante trattamenti successivi quali la post-ossidazione) al fine di individuare le condizioni di trattamento ottimali per preservare

una discreta resistenza a corrosione. A tale scopo sono state eseguite prove di corrosione accelerata, mediante acquisizione di curve di polarizzazione in cella elettrochimica, su provini sottoposti a diverse combinazioni di trattamenti termici e termochimici (invecchiamento H1025, nitrurazione, post-ossidazione) eseguiti in diverse

combinazioni di tempi e temperature. Oltre alla caratterizzazione elettrochimica i provini sono stati sottoposti a prove di microdurezza HV superficiale ed in sezione, per valutare l’effettiva efficacia dei trattamenti applicati in

termini di proprietà meccaniche ed antiusura. Nonostante i migliori risultati in termini di incremento della durezza superficiale e di profondità di indurimento siano stati raggiunti con trattamenti che compromettono notevolmente

la resistenza a corrosione dell’acciaio, il lavoro svolto ha permesso di individuare ed ottimizzare sequenze di trattamenti che permettono di preservare quasi completamente la resistenza a corrosione dell’acciaio, pur

incrementando la durezza superficiale fino a valori di oltre 850HV.

Parole chiave: Acciaio inossidabile - Corrosione - Trattamenti termici - Caratterizzazione materiali

R. Giovanardi, G. Poli, P. VeronesiUniversità di Modena e Reggio Emilia,

Dipartimento di Ingegneria ‘Enzo Ferrari’,Via Vignolese 905, 41125 Modena

G. Parigi, N. RaffaelliSTAV srl

Via della Lora 18/I-N,50031 Barberino del Mugello (FI)

INTRODUZIONE

Gli acciai inossidabili PH (precipitation hardening) trovano impiego in una varietà di applicazioni, quali raccordi aerei [1], ingranaggi e fasteners [2, 3], componenti per reattori nucleari [4-7], tuttavia il loro utilizzo è fortemente limitato dalla loro relativamente bassa durezza e soprattutto dal-

le scarse proprietà tribologiche. Per questo motivo risulta estremamente interessante la possibilità di applicare trat-tamenti superficiali o rivestimenti che possano incremen-tare le proprietà antiusura di tali acciai. In questo lavoro verranno sperimentati trattamenti termochimici di nitrura-zione e post-ossidazione sull’acciaio 17-4PH, con lo scopo di incrementare le proprietà tribologiche di tale acciaio senza comprometterne eccessivamente la resistenza a corrosione. In particolare si cercheranno condizioni ope-rative di trattamento che garantiscano:- la formazione di composti ad elevata durezza in zona

superficiale;- il mantenimento di un elevato contenuto di cromo non

legato (come carburo o nitruro) che possa garantire ele-vata resistenza a corrosione dell’acciaio.

In bibliografia sono presenti alcuni studi che riportano le modifiche strutturali subite dall’acciaio 17-4PH quando


Memorie

sottoposto a trattamenti di nitrurazione a diverse tempe-rature [4]. In particolare si ha che:i) per basse temperature di trattamento l’azoto diffonde

nella fase α dell’acciaio, determinandone la saturazio-ne e rendendo più stabile una struttura c.f.c. (vista la sua azione γ-gena); in queste condizioni è possibile ot-tenere una fase α’N (martensite contenente azoto so-vrasatura).

ii) già a temperature di 350°C è possibile la formazione di una fase S metastabile, costituita da austenite espansa (γN).

iii) per temperature pari o superiori ai 420°C la fase meta-stabile S scompare, a causa della formazione di nitruri di cromo (trasformazione fase-S → α’N + CrN).

iv) per temperature superiori ai 450°C si ha completa scomparsa anche della fase α’N a favore dei nitruri di cromo (trasformazione α’N → α + CrN).

Dal punto di vista esclusivamente tribologico, la comparsa di CrN determina un importante miglioramento delle pro-prietà superficiali dell’acciaio, innalzando la microdurezza fino a valori superiori ai 1250 HV.Nel presente lavoro si ricerca tuttavia una condizione di trattamento che possa garantire sì un incremento della durezza superficiale, ma senza compromettere eccessi-vamente la resistenza a corrosione del materiale. Tratta-menti che portano alla formazione di elevati tenori di CrN saranno pertanto da escludere, in quanto determineranno un impoverimento di cromo tale da rendere l’acciaio non più in grado di passivarsi. Saranno pertanto possibili diverse strategie di trattamento:- lavorare a temperature inferiori ai 450°C, cercando la

giusta combinazione di tempi e temperature di tratta-mento tali da incrementare la microdurezza superficiale senza portare ad eccessivo impoverimento di cromo del-la matrice;

- lavorare nelle condizioni che garantiscono la maggior microdurezza superficiale (450°C o superiori) tentando di ripristinare la resistenza a corrosione mediante un post-trattamento di ossidazione.

Un’ulteriore variabile è costituita dal trattamento di in-vecchiamento artificiale. Gli acciai PH sono infatti quasi sempre sottoposti a tale trattamento, che induce la pre-cipitazione di intermetallici estremamente fini, allo scopo di incrementarne la durezza (Precipitation Hardening). Il trattamento di invecchiamento, che può variare a seconda della composizione dell’acciaio, consiste solitamente in un riscaldamento a temperature nel range dei 500-600°C, mantenimento a tali temperature per un periodo di tempo di circa 4 ore e successivo raffreddamento in aria.Per valutare l’influenza che il trattamento di invecchia-mento può avere sui successivi trattamenti termochimici applicati, sono stati previsti trattamenti di nitrurazione (a parità di condizioni) sia su provini invecchiati (selezionan-do come invecchiamento standard l’H10251) che su provini allo stato solubilizzato.

Sigla campione

Tipo di trattamento

A solubilizz. + invecchiamento H1025

B solubilizz. + nitrurazione 520° 12h + post-ossidazione a 470°

Csolubilizz. + invecchiamento H1025 + nitrurazione 520° 12h + post-ossidazione a 470°

D solubilizz. + nitrurazione 470° 4h + post-ossidazione a 470°

Esolubilizz.+ invecchiamento H1025 + nitrurazione 470° 4h + post-ossidazione a 470°

H solubilizz.+ nitrurazione 400° 16h

I solubilizz. + invecchiamento H1025 + nitrurazione 400° 16h

L solubilizzazione (stato di fornitura)

M solubilizz. + nitrurazione 440° 16h + post-ossidazione a 440°

N solubilizz. + nitrurazione 440° 16h

Tab. 1: codifica campioni e specifiche trattamenti termici e termochimici applicati.

Tab. 1: samples identification, in term of thermal and thermochemical treatments applied

In Tabella 1 sono riportati i trattamenti individuati per con-durre la ricerca. Oltre al campione allo stato di fornitura, cioè privo di trattamenti (campione L) sono state previ-ste diverse combinazioni di trattamenti che prevedono appunto: i) invecchiamento H1025, ii) nitrurazione ionica, iii) post-ossidazione. In particolare la nitrurazione è stata condotta introducendo in camera per primo l’idrogeno, gas riducente in grado di agire sugli ossidi passivanti dell’ac-ciaio inossidabile e garantire pertanto un’adeguata prepa-razione delle superfici dei provini, consentendo la giusta diffusione di azoto durante il processo di nitrurazione.

PARTE SPERIMENTALE

Tutti i campioni allo stato di fornitura sono stati lucidati superficialmente prima di eseguire i trattamenti. A seguito del trattamento di invecchiamento H1025 i provini si rico-prono di una patina superficiale bluastra (ossidi iridescen-ti). Come prassi si è deciso di rimuovere meccanicamente (rilucidatura) tale patina prima di eseguire gli ulteriori trat-tamenti previsti (vedi Tabella 1). Per il solo campione allo stato di fornitura (L) è stata ese-guita un’analisi chimica mediante quantometro allo scopo di verificare se la composizione dell’acciaio è in linea con quanto previsto dalla denominazione 17-4PH. Per l’intero set di campioni sono state eseguite prove di

1 invecchiamento alla T di 550° (± 8°C) per un tempo di 4 ore e successivo raffreddamento in aria


microdurezza Vickers superficiale (carico applicato 1kg forza) per valutare gli incrementi di durezza ottenuti a se-guito dei diversi trattamenti. I valori di HV1 sono stati ot-tenuti come media di 10 misure eseguite in diverse zone della superficie del provino.Sono inoltre stati acquisiti profili di microdurezza, ese-guendo indentazioni Vickers con carichi di 100 g forza sulle sezioni dei campioni preventivamente spianate e lu-cidate (sequenza di carte abrasive e panni con sospensioni diamantate).La resistenza a corrosione è stata valutata, per ciascun campione di Tabella 1, mediante due prove di corrosione accelerata secondo normativa ASTM-G5, operando sul lato lucido e variando, nelle due prove, l’ambiente corrosivo: nella prima prova è stato utilizzata una soluzione di cloru-ro di sodio (NaCl) 3.5%m/m (che simula l’azione aggressiva degli ioni cloruro tipica di un’acqua marina), nella seconda è stata utilizzata una soluzione di acido solforico (H2SO4) 0.5M (che simula un ambiente acido tipico da condensa in atmosfera industriale e che rappresenta l’ambiente ti-pico di prova per gli acciai inossidabili secondo normativa ASTM-G5)Di seguito vengono riportate brevemente le specifiche del-la prova di corrosione accelerata:

- area superficiale di campione analizzata: 1cm2;- ambiente: soluzione di NaCl 3.5%m/m oppure soluzione

di H2SO4 0.5M;- polarizzazione eseguita mediante il seguente ciclo:a) polarizzazione catodica dal potenziale di riposo del

2 Nel caso della prova in H2SO4 i potenziali applicati sono diver-si, per assicurarsi di raggiungere la completa transpassivazione dell’acciaio durante la prova: a) polarizzazione fino a (Er – 0.2)V; b) polarizzazione fino a (Er + 1.8)V

campione (Er) fino al potenziale (Er - 0.4)V2; b) polarizzazione catodica dal valore raggiunto precende-

mente, (Er - 0.4)V, fino al valore (Er + 1.6)V2; - velocita di scansione applicata: 0.0004 V/s;- potenziali misurati rispetto ad elettrodo di riferimento

Ag/AgCl/KCl(saturo);

Al termine di ciascuna prova di corrosione sono state ac-

Composition

Carbon 0.07 max.

Manganese 1.00 max.

Phosphorus 0.040 max.

Sulfur 0.030 max.

Silicon 1.00 max.

Chromium 15.00-17.50

Nickel 3.00-5.00

Copper 3.00-5.00

Carboniumplus Tantalum 0.15-0.45

Elemento % (in peso)

Carbonio 0.030

Manganese 0.642

Fosforo 0.024

Zolfo 0.012

Silicio 0.525

Cromo 15.42

Nichel 4.32

Rame 3.56

Tantalio 0.005

Niobio (colombio) 0.259

Tab. 2: composizione chimica tipica di un acciaio 17-4PH (colonna a sinistra) e composizione chimica di un provino L ottenuta a seguito di analisi al quantometro

(colonna di destra).

Tab. 2: typical chemical composition of 17-4PH stainless steel (left column) and chemical composition of sample L

obtained by quantometer analysis (right column)

Siglacampione

Tipo di trattamentoDurezza

superficialeHV1

Asolubilizz. + invecchiamento

H1025375 ± 2

Bsolubilizz. + nitrurazione 520°

12h + post-ossidazione a 470°

751 ± 113a

Csolubilizz. + invecchiamento H1025 + nitrurazione 520°

12h + post-oss. a 470°628 ± 193b

Dsolubilizz. + nitrurazione 470° 4h + post-ossidazione a 470°

477 ± 14

Esolubilizz.+ invecchiamento

H1025 + nitrurazione 470° 4h + post-oss. a 470°

527 ± 23

Hsolubilizz.+ nitrurazione 400°

16h555 ± 19

Isolubilizz. + invecchiamento H1025 + nitrurazione 400°

16h480 ± 17

Lsolubilizzazione (stato di

fornitura)335 ± 3

Msolubilizz. + nitrurazione 440°

16h + post-ossidazione a 440°

870 ± 30

Nsolubilizz. + nitrurazione 440°

16h860 ± 24

Tab. 3: microdurezze superficiali HV1. (a) elevata deviazione standard in quanto la superficie del provino presenta due zone a diversa durezza (una che fornisce

valori di poco superiori ai 600, l’altra con valori superiori a 800) (b) elevata deviazione standard in

quanto la superficie del provino presenta due zone a diversa durezza (una che fornisce valori compresi fra 450 e 500, l’altra con valori di poco inferiori agli 800).

Tab. 3: surface HV1 microhardness. (a) high standard deviation due to the presence, on the sample surface, of two regions with different hardness (a region with values

of almost 600, another region with values higher than 800) (b) high standard deviation due to the presence, on the

sample surface, of two regions with different hardness (a region with values ranging between 450 and 500, another

region with values of almost 800).

Trattamenti termici


Memorie

quisite immagini della zona sottoposta a polarizzazione mediante stereomicroscopio ottico, al fine di valutare la morfologia di corrosione e la severità dell’attacco subito a livello qualitativo. Tale indagine è utile per capire se vi è presenza di corrosione localizzata e se la localizzazione se-gue particolari punti di innesco (ad esempio bordi grano).

RISULTATI E DISCUSSIONE

Composizione chimicaIn Tabella 2 sono mostrate la composizione tipica di un ac-ciaio 17-4PH (colonna a sinistra) e la composizione riscon-trata sul provino L analizzato al quantometro (colonna a destra). I risultati ottenuti al quantometro confermano che la composizione chimica dell’acciaio soddisfa le specifiche dettate dalla designazione 17-4PH.

Microdurezza superficiale HV1In Tabella 3 sono riportati i valori di microdurezza superfi-ciale (HV1) con relativo errore (deviazione standard asso-ciata alla serie di misure eseguite).Osservando i risultati è possibile notare che:

- il trattamento H1025, applicato sul campione solubiliz-zato (L) come unico trattamento incrementa la durezza superficiale di circa 40 unità HV;

- i trattamenti D, E, H, I portano a valori di HV che, con-siderando le deviazioni standard, possono essere ritenuti confrontabili (spaziano in un range centrato intorno ai 500 HV); il valore migliore si ottiene sul provino H, nitrurato a bassa temperatura per tempi lunghi e senza aver prece-dentemente subito il trattamento di invecchiamento.

- i trattamenti B e C, che prevedono l’applicazione della maggior T di nitrurazione, portano ai valori di HV maggiori ma sono caratterizzati da una spiccata disomogeneità del trattamento; si ottengono infatti deviazioni standard molto elevate motivate dal fatto che entrambi i provini presentano zone ad elevata durezza (800 HV) e zone a durezza più con-tenuta. Probabilmente queste sono le condizioni ottimali per massimizzare l’indurimento superficiale derivante dal-la nitrurazione (formazione di CrN in elevata quantità) ma sembra che l’acciaio inossidabile non sia in grado di subire questo trattamento in maniera omogenea: il trattamento di disossidazione della superficie applicato prima di introdur-re in camera l’azoto (quindi in presenza di solo idrogeno) non sembra pertanto adeguato per garantire omogeneità di trattamento quando si nitrura ad elevata temperatura.

- i trattamenti M ed N, eseguiti a temperature intermedie ma per tempi elevati, sembrano essere i più performanti, portando a valori di microdurezza superficiale confronta-bili con quelli ottenuti con i trattamenti a più elevata tem-peratura, ma assicurando una maggiore omogeneità di trattamento.

- il trattamento di invecchiamento H1025, quando seguito da trattamenti termochimici di nitrurazione, sembra osta-colarne l’efficacia. Confrontando le coppie di provini che hanno subito gli stessi trattamenti ma che differiscono per l’applicazione o meno dell’invecchiamento H1025 (B con C, D con E ed H con I) si nota come le durezze superficiali raggiunte siano sempre maggiori in assenza di invecchia-mento (unica eccezione la coppia D con E dove i valori sono comunque molto simili). Questo risultato fa supporre che i fenomeni attivati durante l’invecchiamento influen-zino il successivo trattamento di nitrurazione; evidenze a riguardo emergono in letteratura [8] anche se non sono riportati modelli esaustivi che spieghino il fenomeno.

Profili di microdurezzaIn Figura 1 sono mostrati i profili di microdurezza HV0.1 ot-tenuti sulla sezione lucidata dei campioni trattati secondo Tabella 1. I risultati confermano le evidenze ottenute dalle microdurezze superficiali, in quanto:

- gli unici trattamenti in grado di determinare elevati valori di durezza e una discreta profondità di indurimento sono quelli eseguiti alle temperature maggiori (B) oppure a tem-perature di 440°C per tempi estremamente lunghi (M ed N);

- il trattamento di invecchiamento H1025, quando seguito da trattamenti termochimici di nitrurazione, sembra osta-colarne l’efficacia, mentre in assenza di post-trattamenti di nitrurazione comporta un incremento della durezza del materiale di oltre 50 punti HV (vedi confronto L-A).

- i trattamenti eseguiti a bassa temperatura (400°C, H ed I) non modificano in maniera sostanziale la durezza degli strati superficiali del materiale, risultando di fatto poco ef-ficaci per gli scopi previsti.

PROVE DI CORROSIONE

Provini A ed LIn Figura 2 sono mostrate le curve di polarizzazione otte-nute sui provini A ed L in ambiente salino (NaCl 3.5% in peso). Le curve mostrano un potenziale di pitting notevol-mente differente: circa 0.15 V per il provino L e circa -0.05 V per il provino A. Il processo di invecchiamento sembra quindi compromettere la resistenza a pitting in ambiente clorurato, abbassando il potenziale di pitting dell’acciaio di quasi 200mV. Il peggioramento è confermato anche dalla corrente media nell’intervallo di passivazione che, per il provino A (curva blu in figura) è leggermente superiore.

In Figura 3 sono mostrate le curve di polarizzazione otte-nute sui provini A ed L in ambiente acido (H2SO4 0.5M). In questo ambiente il trattamento di invecchiamento sembra non influenzare minimamente la resistenza a corrosione dell’acciaio, e le due curve appaiono (pur nella loro com-plessità, cioè con almeno due picchi di attivazione ben de-


Fig. 1 - Profili di microdurezza HV0.1 eseguiti sulla sezione lucidata dei campioni trattati secondo Tabella 1.

Fig. 1 - HV0.1 microhardness profile obtained from polished cross section of samples of Table

Fig. 2 - Curve di polarizzazione ottenute in NaCl 3.5%

Fig. 2 - Polarization curves obtained in NaCl 3.5% solution

Trattamenti termici


Memorie

lineati) perfettamente sovrapponibili.I dati facilmente estrapolabili dalle curve di Figura 3, e che permettono di delineare la resistenza a corrosione dell’ac-ciaio 17-4PH in questo ambiente, sono:- il potenziale di transpassivazione (circa 0.95 V per en-trambi i provini);- la corrente media nell’intervallo di passivazione (varia nel range 5÷7x10-6 Acm-2 per entrambi i provini);- la corrente massima raggiunta dal principale picco di atti-vazione (circa 6x10-4 Acm-2 per entrambi i provini);

I risultati emersi dall’analisi dei provini A ed L mostrano pertanto che il trattamento di invecchiamento H1025 de-termina una diminuzione della resistenza a pitting indotto da cloruri; il risultato può essere interpretato consideran-do una precipitazione di carburi di cromo a seguito del trattamento termico (nonostante la presenza di elevato tenore di elementi stabilizzanti quali niobio); questo peg-gioramento della resistenza a corrosione a seguito dell’in-vecchiamento (trattamento eseguito prevalentemente per incrementare le proprietà meccaniche) valorizza il fatto di intervenire mediante successivi trattamenti termochimici superficiali che possano migliorare la resistenza a pitting (nitrurazione ma soprattutto post-ossidazione).

Provini B, C, D, EIn Figura 4 sono mostrate le curve di polarizzazione otte-nute sul lato lucidato dei provini B, C, D, E in ambiente sali-no (NaCl 3.5% in peso); nello stesso grafico sono riportate le curve ottenute nel medesimo ambiente sui provini ana-lizzati in precedenza: A ed L (non sottoposti a trattamenti termochimici). I provini B, C, D, E hanno subito un tratta-mento di nitrurazione a temperature tipicamente utilizzate per massimizzare i risultati in termini di resistenza ad usu-ra; al trattamento è seguito una post-ossidazione. Purtroppo è piuttosto evidente dai grafici che per tutti e quattro questi provini l’intervallo di passivazione viene quasi a scomparire. Osservando le curve nei tratti iniziali (laddove i campioni A ed L mostravano un plateaux di pas-sivazione a valori di corrente prossimi ai 10-6 Acm2) si vede come i provini A, B, C, D evidenzino solamente un accenno di plateaux di passivazione, a correnti molto più elevate (circa 5x10-6 Acm-2) e di lunghezza (in termini di potenzia-le) estremamente ridotta. Dal punto di vista dell’intervallo di passivazione i quattro campioni si classificano in questo modo:

- campioni B e D (curve verde e fucsia): mostrano un in-tervallo di passivazione di lunghezza compresa fra i 100 e i 150mV, emergendo quindi come leggermente migliori rispetto agli altri due;- campioni C ed E (curve nera e azzurra): mostrano un in-tervallo di passivazione quasi inesistente (meno di 50mV di ampiezza, paragonabile quasi ad un flesso);

Osservando attentamente le curve è possibile notare che, a valori di potenziali maggiori (cioè una volta che il feno-

Fig. 3 - Curve di polarizzazione ottenute in H2SO4 0.5M

Fig. 3 - Polarization curves obtained in H2SO4 0.5M solution

Fig. 4 - Curve di polarizzazione ottenute in NaCl 3.5% sui provini A, B, C, D, E, L

Fig. 4 - Polarization curves obtained from samples A, B, C, D, E, L in NaCl 3.5% solution

meno di pitting è iniziato) tutte e quattro presentano una specie di picco dopo il quale la corrente si abbassa note-volmente. Il fenomeno è molto evidente per la curva fucsia (D), mentre è appena accennato per la curva azzurra (E). Questo fenomeno è molto interessante è può essere ri-condotto ad un arresto del processo di pitting una volta che esso procede oltre la superficie dell’acciaio, ad opera dell’azoto presente in soluzione solida nell’acciaio. L’azoto presente nell’acciaio e liberato durante il processo di cor-rosione è in grado di reagire con gli ioni H+ e di abbattere l’acidità all’interno dei pit che si stanno formando:

[N] + 4H+ + 3e- → NH4+ [9-11].

L’azoto agisce anche da stabilizzante per il film di passiva-zione dell’acciaio, rendendolo più resistente all’attacco de-gli ioni cloruro e può produrre ioni nitrato che aumentano


la resistenza al pitting [12-14].Il provino che sembra beneficiare maggiormente di tale fenomeno protettivo è il D (curva fucsia) che sembra mo-strare una vera e propria seconda zona di passivazione, seppur molto limitata.Considerando entrambi i fenomeni (plateaux di passivazio-ne iniziale e ‘arresto del pit’ dovuto all’azoto in soluzione solida) il campione D emerge come il più performante del-la serie dal punto di vista della resistenza a corrosione, anche se rispetto ai provini non trattati (A ed L) questa pro-prietà sembra essere stata notevolmente compromessa. Ad ogni modo dall’analisi emerge che le migliori perfor-mance in termini di resistenza a corrosione sono state ottenute operando a T inferiori (470°C anziché 520°C) e senza invecchiare il materiale.Osservando il comportamento in acido solforico, Figura 5, queste indicazioni vengono ulteriormente confermate:

- il provino C (invecchiato e trattato ad alta T, 520°C) emerge come il peggiore della serie; oltre ad avere valo-ri di corrente media di passivazione molto elevati questo provino presenta un restringimento dell’intervallo di passi-vazione notevole;- confrontando i provini non invecchiati (D con B, curve fucsia e verde) si nota bene come la resistenza a corrosio-ne peggiori con il trattamento ad alta T (520°C, provino B curva verde); il peggioramento si ha soprattutto in termini di corrente sul picco di attivazione e di corrente media di passivazione;- lo stesso risultato emerge confrontando i provini invec-chiati (E con C, curve azzurra e nera): ancora una volta la resistenza a corrosione peggiora con il trattamento ad alta T (520°C, provino C curva nera);

Analizzando i risultati ottenuti fino a questo punto emerge in modo chiaro come la T dei trattamenti influisca forte-mente sulla resistenza a corrosione degli acciai trattati. Le alte temperature compromettono la resistenza a corrosio-ne sia in ambiente clorurato che in ambiente acido; questo accade sia per il trattamento di invecchiamento (che es-sendo condotto a T di 550°C determina un peggioramen-to della resistenza a corrosione) che per il trattamento di nitrurazione (i risultati ottenuti a 470° sono da preferire a quelli ottenuti a 520°C).

Provini H ed IIn Figura 6 sono mostrate le curve di polarizzazione otte-nute dai provini H ed I in ambiente salino (NaCl 3.5% in peso); nello stesso grafico sono riportate le curve ottenute nel medesimo ambiente su tutti i provini analizzati in pre-cedenza. Dal grafico emerge in maniera chiara che:

- entrambi i provini (H ed I) presentano una zona di passiva-zione iniziale (plateaux) più ampia rispetto a tutti i provini trattati visti in precedenza (B, C, D E), ma ancora inferiore, e a correnti maggiori, rispetto a quella dei provini non trat-tati (L ed A). Questo risultato è comunque molto importan-

Fig. 5 – Curve di polarizzazione ottenute in H2SO4 0.5M sui provini A, B, C, D, E, L

Fig. 5 - Polarization curves obtained from samples A, B, C, D, E, L in H2SO4 0.5M solution

Fig. 6 - Curve di polarizzazione ottenute in NaCl 3.5% sui provini A, B, C, D, E, H, I, L

Fig. 6 - Polarization curves obtained from samples A, B, C, D, E, H, I, L in NaCl 3.5% solution

te in quanto mostra come un trattamento di nitrurazione eseguito a bassa T possa garantire una migliore preserva-zione della resistenza a corrosione mostrata inizialmente dall’acciaio. Da questo punto di vista i due provini H ed I sembrano comportarsi in modo molto simile fra di loro.

- entrambi i provini (H ed I) presentano una notevole estensione della seconda zona di passività individuata nei provini precedenti ed attribuita all’arresto del processo di pitting a causa dell’effetto benefico dell’azoto presente in soluzione solida nell’acciaio. Questo effetto è particolar-mente evidente sul provino H e si estende fino a potenziali prossimi a quelli di transpassivazione originali del provino non trattato L.

In Figura 7 sono mostrate le curve di polarizzazione otte-

Trattamenti termici


Memorie

nute dai provini H ed I in ambiente acido (H2SO4 0.5M); nello stesso grafico sono riportate le curve ottenute nel medesimo ambiente su tutti i provini analizzati in prece-denza. Dal grafico è possibile notare che le curve relative ai provini trattati a bassa T possiedono:

- un intervallo di passivazione esteso quanto quello del provino non trattato (L);

- una corrente media nell’intervallo di passivazione più ele-vata di quella del provino non trattato, ma posizionata a valori minimi fra quelli riscontrati sui provini trattati. Per vedere meglio questo comportamento è stato realizzato il grafico di Figura 8 dove le curve caratterizzate dalle oscil-lazioni più ampie sono state trattate con un filtro passa-basso per permettere di individuare meglio il valore di corrente media di passivazione. Dal grafico si nota come il provino I si collochi a valori minimi di corrente media (paragonabili a quelli del provino D visto in precedenza) mentre il provino H a valori lievemente superiori.

- assenza quasi completa del primo picco di attivazione; da questo punto di vista i provini H ed I sembrano compor-tarsi meglio rispetto a tutti i provini analizzati, compresi i non trattati.

Provini M ed NIn Figura 9 sono mostrate le curve di polarizzazione otte-nute dai provini M ed N in ambiente salino (NaCl 3.5% in peso); nello stesso grafico sono riportate le curve ottenute nel medesimo ambiente su i campioni non nitrurati (A ed L) e sui provini che in precedenza avevano mostrato i mi-gliori risultati in tale ambiente (H ed I). Dal grafico emerge che:

- entrambi i provini (M ed N) presentano una zona di pas-sivazione iniziale più ampia e a correnti inferiori rispetto ai migliori risultati ottenuti in precedenza (H ed I) e, per la pri-ma volta, si hanno curve confrontabili con il campione non nitrurato (ma invecchiato) A. Questi risultati sono i migliori ottenuti sull’intera serie di trattamenti testati, e possono essere messi in relazione con l’ottenimento di una micro-struttura che, in accordo con i dati bibliografici [4] presen-ta un elevato tenore di azoto disciolto (struttura α’N).

- minima invece la differenza fra i due provini M ed N, risul-tato che consente di affermare che, ottenendo una nitru-razione performante in termini di resistenza a corrosione, non è necessario applicare un post-trattamento di ossida-zione.

In Figura 10 sono mostrate le curve di polarizzazione otte-nute dai provini M ed N in ambiente acido (H2SO4 0.5M); nello stesso grafico sono riportate le curve ottenute nel medesimo ambiente su tutti i provini analizzati in prece-denza. Dal grafico è possibile notare che le curve relative ai provini M ed N non risultino particolarmente performan-ti in ambiente acido. Mentre per il provino M tale com-

Fig. 7 - Curve di polarizzazione ottenute in H2SO4 0.5M sui provini A, B, C, D, E, H, I, L

Fig. 7 - Polarization curves obtained from samples A, B, C, D, E, H, I, L in H2SO4 0.5M solution

Fig. 8 - Curve di polarizzazione ottenute in H2SO4 0.5M sui provini A, B, C, D, E, H, I, L (smoothing applicato alle curve con oscillazioni più ampie utilizzando un

filtro passa-basso)

Fig. 8 - Polarization curves obtained from samples A, B, C, D, E, H, I, L in H2SO4 0.5M solution (the curves that

showed large current oscillations were smoothed with a low-pass filter)

portamento è probabilmente attribuibile ad un incremento dell’area superficiale attiva a causa del post-trattamento di ossidazione, per il provino N si hanno risultati simili a quelli ottenuti per il provino D (trattato a 470°C).Questi risultati permettono di concludere che le ottime performance di resistenza a corrosione in ambiente clo-rurato ottenute con i trattamenti M ed N derivano da una struttura ad elevato contenuto di azoto disciolto, e non dall’assenza di formazione di precipitati CrN.L’azoto disciolto, come visto in precedenza, è in grado di rendere meno severi gli attacchi di tipo caverniforme (pitting) [9-14], mentre in un ambiente decisamente acido


Fig. 9 - Curve di polarizzazione ottenute in NaCl 3.5% sui provini A, H, I, L, M, N

Fig. 9 - Polarization curves obtained from samples A, H, I, L, M, N in NaCl 3.5% solution

Fig. 10 - Curve di polarizzazione ottenute in H2SO4 0.5M sui provini A, B, C, D, E, H, I, L, M, N

Fig. 10 - Polarization curves obtained from samples A, H, I, L, M, N in H2SO4 0.5M solution

(H2SO4) risulta fondamentale il contenuto di cromo dispo-nibile in lega (non legato sotto forma di carburi e nitruri).I campioni H ed I, ottenuti a bassa temperatura, rispondo-no bene a corrosione anche in ambiente acido in quanto risulteranno privi di precipitati di CrN (come confermato anche dalle scarse proprietà di indurimento superficiale riscontrate dalle misure di HV superficiale ed in sezione), mentre i campioni M ed N rispondono bene a corrosione solo in ambiente clorurato in quanto, essendo stati trattati per tempi lunghi a T superiori (440°C) mantengono sì un elevato contenuto di N disciolto, ma avranno sicuramente dato luogo anche alla precipitazione di nitruri di cromo (fe-nomeno ancora una volta confermato dagli ottimi valori di HV riscontrati per questi provini).

CONCLUSIONI

L’attività di ricerca ha permesso di caratterizzare dal pun-to di vista della resistenza a corrosione in due differenti ambienti (clorurato ed acido) diverse combinazioni di trat-tamenti termici e termochimici applicati all’acciaio 17-4PH allo scopo di migliorarne la resistenza meccanica e le pro-prietà tribologiche. I risultati più importanti emersi dalla ricerca sono di seguito riportati:

1) i trattamenti che permettono di ottenere il miglior com-promesso fra durezza superficiale, profondità di induri-mento e resistenza a corrosione sono quelli condotti per tempi lunghi (16 ore) alla temperatura di 440°C; in tali condizioni è possibile ottenere un elevato tenore di azo-to disciolto e una minima formazione di nitruri di cromo, risultato che porta ad un’ottima resistenza a corrosione in ambiente clorurato e ad una notevole microdurezza su-perficiale (superiore ad 800HV). In tali condizioni non si evita comunque l’impoverimento di cromo della matrice dato dalla precipitazione dei CrN, fenomeno che porta ad avere una resistenza a corrosione in ambiente acido di-scretamente compromessa rispetto a trattamenti analoghi eseguiti a più bassa temperatura (400°C).

2) i trattamenti a bassa temperatura (400°C) permettono di mantenere ottime proprietà di resistenza a corrosione sia in ambiente acido che clorurato, ma determinano un incremento della durezza superficiale molto modesto.

3) il trattamento di invecchiamento H1025, se applicato prima di eseguire la nitrurazione, influenza negativamen-te i risultati ottenibili dal trattamento termochimico. L’in-cremento di durezza a cuore ottenibile con il trattamento H1025 è comunque perseguibile anche durante il normale processo di nitrurazione (che opera a temperature tali da consentire i fenomeni di precipitation hardening), pertanto risulta conveniente trattare termochimicamente provini di acciaio 17-4PH allo stato solubilizzato.

BIBLIOGRAFIA

[1] W.T. Chien, C.S. Tsai, J. Mater. Proc. Technol. 140 (2003) 340.[2] P. Li, Q. Cai, B. Wei, X. Zhang, J. Iron. St. Res. 13 (2006) 73.[3] P. Kochmanski, J. Nowacki, Surf. Coat. Technol. 200 (2006)

6558.[4] G. Li, J. Wang, C. Li, Q. Peng, J. Gao, B. Shen, Nuclear Instruments

and Methods in Physics Research B 266 (2008) 1964.[5] J. Wang, H. Zou, C. Li, Y. Peng, S. Qiu, B. Shen, Nucl. Eng. Design

236 (2006) 2531.[6] J. Wang, H. Zou, C. Li, S. Qiu, B. Shen, Mater. Charact. 57 (2006)

274.[7] F. Christien, R. Le Gall, G. Saindrenan, Scripta Mater. 48 (2003)

11.[8] P. Kochmanski, J. Nowacki, Surface & Coatings Technology 202

(2008) 4834.[9] C.X. Li , T. Bell, Corrosion Science 46 (2004) 1527.[10] S.D. Chyou, H.C. Shih, Corrosion 47 (1991) 31.[11] H.J. Grabke, The role of nitrogen in the corrosion of iron and ste-

els, ISIJ International 36 (1996) 777.[12] I. Olefjord, L. Wegrelius, Corrosion Science 38 (1996) 1203.[13] H. Baba, T. Kodama, Y. Katada, Corrosion Science 44 (2002)

2393.[14] U. Kamachi Mudali, P. Shankar, S. Ningshen, R.K. Dayal, H.S. Kha-

tak, B. Raj, Corrosion Science 44 (2002) 2183.

Trattamenti termici


Memorie

Nitriding and post-oxidation treatmentson 17-4PH stainless steel: optimization of the process to preserve high corrosion resistance

Keywords: Stainless Steel - Corrosion - Thermal Treatment - Material Characterization

17-4PH stainless steel is usually heat treated to increase its mechanical properties. In order to obtain a further im-provement of the wear resistance of this steel, it is possible to apply thermochemical treatments, such as nitriding. The untreated 17-4 PH stainless steel has a remarkable corrosion resistance conferred by the high chromium of this alloy, so it’s important to evaluate how the application of a thermal or thermochemical treatment can affect this property.

The aim of this work is to check how the application of thermochemical treatments, usually used to increase the me-chanical and wear properties of iron alloys, affect the corrosion resistance of this steel and to optimize the process variables (considering also the possibility to add subsequent treatments, such as post-oxidation) in order to identify the best treatment conditions in order to preserve a good corrosion resistance.

For this purpose accelerated corrosion tests were performed, through the acquisition of polarization curves in an electrochemical cell, on specimens subjected to different combinations of heat and thermochemical treatments (H1025 aging, nitriding, post-oxidation). In addition to the electrochemical characterization, the specimens were characterized by surface HV microhardness tests and by HV microhardness profiles along their cross-section, to assess the effectiveness of the applied treatments in terms of mechanical properties and wear resistance.

Despite the best results (in terms of increasing of surface hardness and depth of hardening) have been achieved with treatments that significantly compromise the corrosion resistance of the steel, the work has allowed to identify and optimize sequences of treatments that preserve almost completely the corrosion resistance of the steel, while increasing the surface hardness up to values of 850HV.


Validity of Wulff construction used for size-dependent melting point of nanoparticles

S. Zhang, L. Zhang, L. Chen

An integrated model based on the variant of Ba/Bt, is established to predict size-dependent melting point of nanoparticles by considering the geometric and energetic characteristics of Wulff construction. Ba is the rest bond number and Bt denotes the total bond number without broken bonds in a Wulff construction. Without

any adjustable parameters, this model predicts a decreasing trend of melting point with the size dropping for nanoparticles. The good agreement between theoretical predictions and the evidences in experiments and

molecular dynamic simulation confirms the validity of Wulff construction in describing thermodynamic behaviors of nanoparticles even with no need in considering their crystalline structures.

Keywords: Metals - Nanoparticles - Wulff construction - Melting

S. Zhang, L. ZhangSchool of Mechanical Science and Engineering,

Jilin University,Changchun 130025, China

[email protected]

S. ZhangDepartment of Materials Science and Engineering,

Jilin JianZhu University,Changchun 130000, China

L. ChenDepartment of Municipal and environmental

engineering, Jilin JianZhu University,Changchun 130000, China

[email protected]

Nanomateriali

INTRODUCTION

The thermodynamic behavior of nanocrystals differs from that of the corresponding bulk materials mainly due to the large value of surface-to-volume ratio, which strongly in-fluences both the chemical and physical properties in com-parison with the bulk counterpart [1-4]. This is because the surface/volume ratio depends on both size and shape, and the size and shape or structure strongly influences many fundamental properties of nanoparticles [5]. However, the shape or structure is strongly depending on size of materials [6-8]. It has been predicted that Na [9] and Mo [10] substan-ces with a bulk bcc structure would have fcc or more like ico-sahedron structures for nanoparticles. This is because the fcc or icosahedron structures are more compact than the bcc structure and provide a lower surface energy than the

bcc one. It is also found that Co nanoparticles with radius below 10 nm prefer to form a fcc structure, rather than bulk hcp one [11]. Moreover, many other nanoparticles bound by van der Waals or metallic forces (such as Mg, Ca, Sr, Ni and Ba) exhibit structures with fivefold axes of symmetry, i.e., icosahedron structure, despite the fact that the bulk metals exhibit hcp, fcc or bcc packing [12]. It should be noted that nanoparticles must display the bulk crystalline structure at larger r (r shows the size of nanoparticles). Therefore, we can expect that it is the surface energy controlling the sha-pe or structure of nanoparticles, namely, the structure of nanoparticles is the one with the smallest surface energy. Since the shape or structure affect most properties of nano-particles, it is necessary to be investigated. It is known that Wulff construction, which is developed by minimizing surfa-ce energy for a given enclosed volume, is the standard me-thod for determining the equilibrium shape of crystals at the macroscale level. This requirement for small surface energy is also applied to nanoparticles, since the surface energy is the major energy contribution for them. So, considering the compact packing of nanoparticles mentioned above, the ge-ometrical and energetic properties of Wulff construction for fcc crystal is taken for describing nanoparticles in this work. Through introducing a variant Ba/Bt to describe the geome-tric characteristics of Wulff construction, a model without any adjustable parameter is obtained to estimate size-de-pendent melting point of nanoparticles, where Ba is the rest bond number and Bt denotes the total bond number without broken bonds in a system. The good agreement between our model predictions and experimental results suggests that it is valid that taking Wulff construction as nanoparti-cles’ structure for predicting their melting temperature.


Memorie

MODEL

It is known, cohesive energy that describes the bond strength directly, is an effective variable to determine the thermal stability of nanocrystals. With the size reduction, the decline of melting point is an obvious, which implies the lowered thermal stability of nanocrystals. In fact, there is an empirical correlativity between E0 and Tm0 functions, by defining E0 and Tm0 as bulk cohesive energy and bulk melting point [13, 14],

(1.1)

In Eq. (1.1), kb is the Boltzmann’s constant. According to Eq. (1.1), if applying this relation to the nanoscale a similar treatment for the relationship between E(r) and Tm(r) fun-ctions can be expected as a first approximation, that is, (1.2)

Therefore, combining Eqs (1.1) and (1.2), the ratio of the melting temperature of the nanoparticles versus that of the bulk can be read as,

(2)

For a system, E(r) function has been derived by introducing the variant of Ba/Bt, that is [15, 16],

(3)

The broken bonds of the atoms on surfaces inevitably lead to the instability of materials in nano-scale. Thus, as long as Ba/Bt is known, E(r) or Tm(r) is obtained. However, it is necessary to know nanopaticle’s structure and size, since both of them decide Ba and Bt values. For most metallic nanoparticles, the most proper structure of nanoparticles could be Wulff construction. It is clear that Ba or Bt is strongly dependent on the size and shape, since Ba and Bt actually are the multiplying re-sults between the atom number and the average coordina-tion number [15], namely Ba/Bt = ZsNs/ZbNt, where Zs and Zb are average coordination number for surface atoms and bulk interior, and Ns, Nt are the number of surface atoms and total atoms in a system, respectively. So Eq. (3) indi-cates the size and shape dependences of cohesive energy, and even for melting point of nanoparticles. Wulff construction is a segment of fcc (faced-centered-cubic) crystal. By truncating a octahedron, one can obtain a polyhedron with fourteen facets. There have six square (100) facets and eight hexagonal (111) facets at its surface, in which three edges of the hexagon are in common with square (100) facets, while the remaining three edges in common with hexagonal (111) facets. And each edge has

same atom number. Arriving here, a Wulff construction is established, and the size or diameter of a Wulff construc-tion can be altered by controlling the the atom number on edge. To obtain the Ba/Bt of a Wulff contruction, the he total atoms number (Nt) and surface atoms number (Ns) must be known. Let n denoting the atom number on a edge, Nt and Ns can be resolved, e atom number on a edge, Nt and Ns can be resolved,

(4.1)

The number of surface atoms can be expressed as fol-lowing:

(4.2)

In fact, the value of Ns includes the number of the atoms on (111) facets (N111), the number of atoms on (100) faces (N100), the number of atoms at edges (Ne) and the number of atoms on vertex Nv, that is Ns = N111+N100+Ne+Nv. From mathematic point, N111 = 8(3n2-9n+7), N100 = 6(n-2)2, Ne = 36(n-2) and Nv = 24. In addition, the coordination number should be resolved to obtain Ba/Bt. However, the coordi-nation number for atoms at different sites is also different, that is Z111 = 9, Z100 = 8, Ze = 7 and Zv = 6, respectively. So, the average coordination number of surface atoms can be expressed:

(5)

Then, one can obtain Ba = ZsNs/2. It is also easy to get Bt value, since Bt = NtZb/2. Nt is given by Eq. (1). Zb is the coordination number of bulk interior atoms, and Zb = 12 for fcc structure. Therefore, Ba/Bt can be expressed as the following formulation,

(6)

It is clear that n is related with the size of Wulff construc-tion. So the value of Ba/Bt for Wulff construction is relying on size. It is clear that different shape has different Ba/Bt value. That is to say the value of Ba/Bt is simultaneously related with both size and shape. Assuming the radius D of Wulff construction as the biggest distance from center atoms to surface atoms, D has the following expression,

(7)

with h being atomic distance.Substituting Eqs. (3) and (6) into Eq. (2), size-dependent melting point can be expressed

(8)

RESULTS AND DISCUSSIONS

Fig. 1 shows the comparison between model predictions in light of Eq. (8) and experimental results for melting points


of several metallic nanoparticles. It is clear that a good agreement between them is found. As expected, Tm(r) is a continuous function of r and decreases monotonically as r decreases, leading to the lowered thermal stability. This is because of the decreased Ba/Bt value. The results di-splayed in Fig. 1 confirm the success of Wulff construction for describing the geometric and energetic characteristics of nanoparticles almost throughout the whole size range. This is because the variant Ba/Bt appearing in Eq. (3) is re-lated not only to size, but also to shape or structure. It can effectively change the Ec(r) and Tm(r) value by swaying the thermodynamic stability due to the change in Bt value. As r decreasing, Ns relatively increases, which results in the decrease in the total bond number and the increase in the broken bond number. Moreover, it should be noted that Eq. (8) is still valid for In and Sn nanoparticles with their bulk structures being tetragonal. Therefore, it is expected that taking Wulff construction as a standard shape to describe nanoparticles is reasonable in full size range from micro to macro without taking structure change into account.For larger particles with r > 10 nm, the validity of Wulff construction is clear. This is because, the change in bond energy in comparison with that in bulk interior, is small, and Ba = Bt for larger particles. However, the assumption used in Eq. (10) also results small difference for smaller particles, since only surface bond relaxation is conside-red in Eq. (3). In fact, except surface atoms, interior atoms also become unstable compared to bulk interior, resulting in larger estimation of Eq. (8). In addition, the defect or vacancy in a nanoparticle is not considered in this work, which means the result of ideal crystal by using Eq. (8). This also may lead to small overestimation of Eq. (8), of necessity for small particles, as presented in Fig. 1. De-spite the existing errors, Eq. (8) can still be regarded as a valid and simple way to predict Tm(r) values even in full size range. It should be note that for small nanoparticles with r < 5 nm, the validity of Eq. (8) implies that the na-noparticles possess close-packed structure whatever the bulk structure is. Based on Eq. (8), to determine the Tm(r) or E(r) values of nanoparticles, there is no need to know surface energies or shape, and even other thermodynamic information but atomic distance and the size of nanopar-ticles. In the previous studies of nanoparticles, a spherical shape is usually taken into account, and the reasonability of this action for melting point of nanoparticles is also presented by taking the ratio of surface/volume as the only variant [17-20]. To further confirm the validity of Wulff construc-tion developed in this work, the ratio of surface to volume (δ) may explain this point. For spherical nanoparticle, can be simply determined as,

(9)

For comparison, δ function for Wulff construction is gi-ven,

(10)

Fig. 1 - The comparison of Tm(r) functions between model prediction in terms of Eq. (7) (solid lines) with the help of Eq. (6) and experimental results of Al, Pb, In and Sn nanoparticles, respectively, where ●, ◄, ♦, and ► show experimental results [6]. The h used in Eq. (6) are separately 0.3164 nm, 0.3870 nm, 0.3684

nm, and 0.3724 nm for Al, Pb, In and Sn elements.

To further describe the validity of Wulff construction in de-scribing the shape of nanoparticles, the comparison of δ by Eqs (9) and (10) is made. As shown in Fig. 2, the changes in δ with respect to size are presented. Similar trend for spherical nanoparticles with that of Wulff construction is found. As r increases, δ decreases and δ → 0 with r → ∞. And the difference in δ between spherical shape and Wulff construction decreases with r increasing. When r > 6 nm, the difference between them is almost indistinguishable. Note, with r decreasing, the particles is no longer a sphere

Fig. 2 - The comparison of δ(r) functions between sphere (dotted line) and Wulff construction (solid line) in light of Eqs. (8) and (9) respectively, where h = 0.3

nm is taken for simplicity.

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one, thus the model based on sphere consideration is not reasonable. However, Wulff construction is the truncated octahedron with fcc structure and due to the small surface energy, some small particles perfer to take this shape. As a result, the model established in this work could be used to predict the thermodynamic stability of small particles. In addition, our results also strongly support the assum-ption of spherical shape usually considered for particles.

CONCLUSION

By utilizing the geometric characteristic of Wulff construc-tion, the size-dependent melting point of nanoparticles is modeled with the help of the variant Ba/Bt. Similar to other melting models, this model predicts the decreasing trend of melting point when the size is dropping. This mainly arises from the lowered bond number in a nanoparticle if compared with its bulk material. The consistency of the model predictions and experimental results suggests the validity of Wulff construction, one hand to describe the shape or structure, and the other to describe the thermo-dynamic stability of nanoparticles.

ACKNOWLEDGEMENTS

The authors acknowledge the financial supports of the Science and Technology project of Jilin province educa-tion department during the Twelfth Five-year Plan Period ( No.2013232). National Natural Science Foundation of China (grant No. 51101067), Natural Science Founda-tion of Anhui Higher Education Institutions of China (No. KJ2012B159), Open Foundation of Key Laboratory of Auto-mobile Materials of the Ministry of Educations, Jilin Univer-sity (No. 12-450060481289).

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371-423.20] Y.F. Zhu, J.S. Lian, and Q. Jiang, J Phys. Chem. C 113

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Microstructural characterization and production of high yield strength rebar

E. Mansutti, G. Luvarà, C. Fabbro, N. Redolfi

Various technical standards from all over the world set out the mechanical and chemical characteristics for high yield strength rebar. High yield strength rebar - as defined in this study – is applied to all concrete reinforcement

steel grades which require a minimum yield strength of 600MPa. The standards concerning rebar production were reviewed in order to select all the possible grades that come under the above-mentioned definition.

This research project aims to determine if by applying an in-line quenching and self-tempering process, the technological requirements for high yield strength rebar, as specified in the standards, can be met, in order to optimize the chemical composition and save on alloying elements. The work can be divided into two different phases. The preliminary phase took place in the metallurgical laboratory of Danieli’s research center and the second phase in an industrial plant. Tests done in the laboratory set out to evaluate the effect of quenching

and chemical composition on the rebar’s final mechanical properties and microstructure. The purpose of the industrial-scale tests was to evaluate the potential of DANIELI’s in-line quenching and self-tempering process, referred to as QTB (Quenching and Tempering Bar process), applied to high-strength steels. At the end of the

lab tests, three different chemical compositions were selected, deemed suitable for the production of high yield strength rebar. In the industrial-scale tests it was then possible to evaluate the performance of the QTB process in the production of high yield strength rebar in terms of operating flow rates / pressures, optimized chemical

compositions, productivity and process stability.

Keywords: Rebar - Yield Strenght – Quenching - Microstructure

E. Mansutti, G. Luvarà, C. Fabbro, N. RedolfiDanieli & C. Off Meccaniche, Buttrio

INTRODUCTION

The application of high yield strength rebar is provided for in various technical standards from all over the world, such as the US, Russia, Korea and Japan.Russia, for example, introduced the concept of high-yield rebar (980MPa) back in 1982, which was then developed further in GOST 10884 issued in 1994.The mentioned GOST standard takes advantage of the known effect of silicon on enhancing elastic limit, allowing it to be added up to a maximum of 2.3% so that the steel can be included in the At1200 class (or class VI, conside-ring the former standard), which corresponds to a mini-mum yield strength of 1200MPa.Less indicative, however, is the recent Korean standard that for the SD700 class only specifies a limitation regar-ding equivalent carbon (CeqIIW = 0.63).In Japan in 1993 a research project was carried out, refer-red to as “New RC Project”, which was then incorporated into the National Building Code [1] [2].The US has published the most recent ASTM standards on this subject. Both standard A615/A615M and A706/

A706M introduced “Grade 80”, which not only requires mi-nimum yield stress values but also particularly high mini-mum UTS values (725MPa for standard A615 and 690MPa for A706). In addition, the A706 is more demanding in terms of Rm/Rp ratio, maximum carbon content and Ceq; in practice this makes it more complicated to apply on-line heat treatments in rolling mills, requiring greater attention to be placed on chemical composition. It is also important to bear in mind that compared to European standards, US standards are more stringent in terms of statistical relia-bility of technological values, requiring rebar producers to guarantee yield strengths that are significantly higher than the minimum requirements of the standard.Again, in the US market standard A1035 provides for the possibility of producing high-tensile corrosion-resistant re-bar through high chrome content (around 9%) and by con-trolling the final microstructure by taking advantage of the new technologies for in-line heat treatments [3].It is interesting to note that in the US market various rebar producers are pushing for the introduction of high yield-strength grades (such as proposing classes “100” and “125”), even if current market demand for this type of pro-duct is low.In China there are no reference standards for equivalent grades, although some studies refer to the use of V-N mi-croalloyed steels and ultrafine grained steels [4][5][6].

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COUNTRY Ref. Standard Maximum yield strength Remarks

RUSSIA GOST 10884-94 1200 MPaHigh yield strength with ad-dition of silicon up to 2.3%

UKRAINE DSTU 3760-06 1000 MPa -

JAPAN “New RC Project 1993” 980 MPaAlso includes grades @ 1275 MPa but only for transverse reinforcement applications

USA ASTM A1035-14830 MPa(120 ksi)

High yield strength by con-trolling microstructure

KOREA KS D3504-11 700 MPaCeq increase allowed up to

0.63

ENGLAND BS 6744-01+A2:09

650 MPa Stainless steel rebar

INDIA IS 1786-08 600 MPaMicroalloyed steel with maximum Ceq of 0.53

CHINA GB1499.2-07 500 MPa Ceq max 0.55

Tab. 1 – Overview of international standards for high-tensile rebar. Ceq as per IIW standard: (C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5.

Fig. 1 – Diagram of a timed quenching system and example of a sample with a thermocouple attached to it (top). Photo showing the quenching station with heating oven and timed quenching system.


Table 1 summarizes the main international reference stan-dards for high yield strength rebar.Civil engineering applications until now have been rather li-mited even if this type of rebar is promising as it simplifies the reinforcement of concrete [7].

GOAL

This study aims to examine the possibility of meeting the technological requirements specified in various internatio-nal standards for high yield strength rebar, using the in-line quenching and self-tempering process, thereby optimizing the chemical composition with considerable savings in al-loying elements.Tests done in the laboratory set out to evaluate the effect of quenching and chemical composition on the rebar’s fi-nal mechanical properties and microstructure.The industrial-scale tests then evaluated the potential of DANIELI’s in-line quenching and self-tempering process, referred to as QTB (Quenching and Tempering Bar pro-cess), applied to high-strength steels.

LABORATORY TESTS

In order to study the mechanical and microstructural pro-perties of various rebars subjected to a quenching and self-tempering process, a device (suitable for different diameters) was set up to heat-treat samples of rebar that were previously fitted with thermocouples (Figure 1).The system is made up of:- Heating oven with inert atmosphere (Ar)- Brine quenching tank with timed immersion systemFor the experiment it was decided to use a DIA 16 rebar with a length of about 250mm.Shown in Figure 2 is a representative trial heat cycle of a sample subjected to testing.The experiment involved:- Heating the sample to 900 °C- Soaking it for 5 minutes

Fig. 2 – Temperature trend measured at the core of a DIA 16mm rebar during the test

- Air cooling it down to 850°C and then quenching it (from 1 to 5 seconds)- Interrupting the cooling process by self-tempering of the material surface, and final air cooling.DIA 16mm rebars with three different compositions were selected for the experiment (see Table 2) in compliance with specific international standards.The following three compositions were used:- Composition#1 with medium carbon content and high silicon content;- Composition #2 with low carbon content and medium Mn content;- Composition #3 with high Mn content and medium Si content.The aim of the experimental plan was to determine the combined effect of C, Mn and Si on hardenability and per-formance in terms of mechanical and microstructural pro-perties. In particular the effect of a composition with lower Mn and higher C and Si contents (such as composition #3 for example), was compared to the other two chemical compositions with lower carbon and higher Mn contents.

LABORATORY EXPERIMENTAL PROCEDURES AND MA-TERIAL CHARACTERIZATION

Samples of each rebar composition shown in Table 2 were quenched at increasing immersion times from 1 to 5 se-conds while the core temperature was continuously moni-tored. For each test both microstructural properties and mechanical strength were analyzed and measured.To facilitate the comparison of results, before performing the tests, various heat treatments were considered to de-cide which one would be used to determine the same prior austenitic grain size for all the samples. This made it pos-sible to use the same heat treatment (briefly described in the previous chapter) for all three steel grades studied.The quenched rebar pieces were subjected to a tensile test to determine their mechanical properties.Figure 3 shows a growing linear trend of yield strength up to a maximum of 1000 MPa for compositions #2 e #3,

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while composition #1 exceeds 1200 MPa.On Figure 3 it can be noted that elongation diminishes along a linear path as quenching time increases, with all three steels following the same trend, while the decrease in Rm/Rp ratio is less marked.Because of its higher carbon and silicon contents, che-mical composition #1 is able to reach the required yield strength within a shorter quenching time. Moreover, even with higher carbon and silicon contents, ultimate elon-gation is not penalized for up to 3 seconds of quenching (which makes it possible to obtain a product with a yield strength of 980MPa).Composition #2, which has lower C, Mn and Si contents, produces the lowest elongation value, even with relatively short quenching times (1sec). On next page is reported a summary of the results for:- Analysis of microhardness (HV0.3) within the cross-sec-tional area of a rebar subjected to different cooling times;- Description of the microstructure observed at the core of the rebar with temperature measured at the end of immer-sion (thermocouple placed at the core of the rebar).In general, a gradual increase is noted in the presence of rapidly cooled structures at the core of the rebar, and even completely hardened structures resulting from quenching times of between 2.5 and 3.0 seconds.Compared to composition #2, the increase in Mn and Si for chemical composition #3 leads to a slight rise in mar-tensite hardness at the end of the cycle (in both cases cooling time was 4”).Increased material hardenability due to higher Mn and Si, together with the effect of tempering stability provided by the silicon, still leads to increased hardness within the cross-sectional area of a rebar quenched for the same amount of time.A comparison of the above results with those of chemi-cal composition #1 show that the high C and Si contents, which ensure greater hardenability, make it possible to achieve complete hardening with shorter quenching times. The result obtained with a shorter quenching time (2 sec) and higher final temperatures (550 °C) is comparable to the performances of the other steel grades.The strategy of using higher amounts of carbon and silicon while reducing the amount of manganese is only effective if managed properly through controlled cooling.

Composition %C %Mn %Si %P %S %V %Al %Cu %Cr %Ni N ppm Ceq

#1 0.36 0.67 0.96 0.033 0.026 0.008 0.003 0.06 0.05 0.03 84 0.49

#2 0.22 0.98 0.18 0.018 0.015 0.003 0.004 0.28 0.10 0.08 117 0.41

#3 0.19 1.31 0.50 0.033 0.033 0.006 0.003 0.03 0.02 0.01 58 0.41

Tab. 2 – Result of the chemical analyses performed on samples from lots selected for the experiment; elements not shown on the table are present only in trace amounts. Ceq according to standard IIW:

(C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5.

SteelCooling

time [s]

Rp0.2[MpA]

Rm[MpA]

Rm/Rp A%A5d

%

2#

1.0 420 603 1.44 16 162.0 600 888 1.48 6 32.5 861 1053 1.22 10 93.0 880 1005 1.14 8 73.5 920 1118 1.22 5 43.5 920 1132 1.23 6 54.0 1068 1296 1.21 3 34.0 1000 1329 1.33 4 3

3#

1.0 434 589 1.36 17 202.0 660 803 1.22 12 32.5 650 829 1.27 12 143.0 920 1025 1.11 8 103.5 1040 1234 1.19 9 94.0 1114 1282 1.15 7 4

1#

2.0 650 830 1.28 15 163.0 980 1188 1.21 10 113.0 980 1276 1.30 8 94.0 1000 1293 1.29 4 34.0 1120 1527 1,36 7 55.0 1200 1578 1.,32 5 5

Fig. 3 – Change in mechanical properties as quenching time increases (right bottom). The measured

mechanical properties are reported on the top right.

Red n.: breakage outside calibrated lenght


INDUSTRIAL-SCALE TESTS

Following the results of the lab tests and in keeping with specific plant requirements, various tests were performed in a real plant to evaluate the capability of in-line heat tre-atment processes (QTB) in the production of high yield strength rebars.These tests were essential in order to validate the results of the laboratory tests, overcome their limitations and sim-plifications and determine the stability of the in-line pro-cess in a real rolling mill. Based on the results of the laboratory tests, an initial che-mical composition was selected in order to ensure good material weldability.Before running the tests, the technological parameters of the rolling mill were studied using the thermal/metal-lurgical software DLPP (Danieli Long Products Predictor), which was also used to evaluate the test results [8]. Figure 7 shows the heat profile of one of the tests, from reheating furnace exit to the cooling bed.For each test, a suitable number of samples was selected for technological and metallurgical characterization. The bend tests and elongation measurements were done using several methods described in various international stan-dards. Figure 8 shows the effects of bend tests on some samples, according to various standards.The rolling tests in conjunction with metallographic and technological characterization made it possible to deter-

Cooling Time TemperatureMicrostrucutural characteri-

stics for composition 2#

[s] [°C]Sub-surface

areaCore

1.0 640 Martensite + PFBainite in

ferritic pearlitic strufture

2.0 471 Martensite + PFBainite and

Ferrite3.0 430 Martensite Martensite + PF4.0 277 Martensite + PF Martensite + PF

Fig. 4 – Results from microstructural analysis of rebar core with composition #2 and microhardness profiles within the cross-sectional area of the quenched rebar,

at increasing cooling times from 1” to 4”.




areaCore

1.0 635 Martensite + PFBainite in ferritic

pearlitic struf-ture

2.0 526 Martensite + PF Bainite3.0 430 Martensite Martensite + PF

4.0 313 Martensite + PF Martensite + PF

Fig. 5 – Results from microstructural analysis of rebar core with composition #3 and microhardness profiles within the cross-sectional area of quenched rebar, at

increasing cooling times from 1” to 4”.




areaCore

2.0 558Martensite +

BainiteBainite


4.0 359 Martensite + FP Martensite + PF


Fig. 6 – Results from microstructural analysis of rebar core with composition #1 and microhardness profiles within the cross-sectional area of the quenched rebar,

at increasing cooling times from 1” to 5”.

PF = Proeutectoid ferrite in trace amounts (<5%)

PF = Proeutectoid ferrite in trace amounts (<5%) PF = Proeutectoid ferrite in trace amounts (<5%)

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mine the limitations of the QTB process in the production of high yield strength rebar, in terms of cooling method, operating flow rates/pressures, chemical compositions, productivity and process stability.It is important to note that for rebars with the same final tech-nological properties, the processing temperatures necessarily differ depending on the composition. In fact, core hardening is necessary in some cases in order to reach the desired figures. This aspect must be taken into consideration to determine the risk of brittle phases being generated, and the possible creation of cracks (enhanced by the quenching process). Fi-gure 9 shows a series of macrographs (hardening depth) and micrographs (surface, core): one of them highlights a crack generated by a defect that spread within the bar.Just like in the lab tests, the microhardness profiles in va-rious processing conditions were examined. This made it possible to determine the exact hardening depth and the effect of the metallurgical transformations.

CONCLUSIONS

The lab experiments made it possible to assess the beha-vior of 3 different chemical compositions after subjecting DIA 16mm rebars to hardening and self-tempering.

Fig. 7 – Thermal simulation using DLPP: temperature trend of the bar from reheating furnace exit to cooling bed entry.

Fig. 8 – Bend tests done on samples of high-tensile rebar heat-treated in line.

Fig. 10 – Yield strength measured with increasing cooling times of rebar in lab tests.

The industrial-scale tests made it possible to evaluate the performance of the QTB process in the production of high yield strength rebar (greater than 1000MPa) in terms of operating flow rates / pressure, optimized chemical com-positions, productivity and process stability.In some DANIELI plants, the QTB process is already being used for the production of high-strength steels.


Fig. 9 – Examples of macrographs and micrographs (surface, core) for high yield strength rebar treated with QTB at different processing temperatures. Note the sample with a crack generated by a defect.

[4] Y. CAIFU, Development of High Strength Construction Rebars, Proceedings of Int. Seminar on production and Application of High Strength Seismic Grade Rebar Con-taining Vanadium, page 58, Beijing 2010

[5] J. HUAIZOHONG, Y. CAIFU, Z. YONGQUAN, Strengthe-ning effects of nitrogen on 20MnSi rebar containing vanadium, Special Steel, Vol.21,No.5, page 20, 2000

[6] WENG YUQUING, Ultra fine grain steel, Metallurgical Industry Press, Beijing, 2003

[7] ENGINEERING NEWS RECORD, High strength rebar cal-led revolutionary, Aug/Sept, 2007

[8] C. FABBRO, M. CIMOLINO High Carbon Grades for Wire Rod Lines - The Core of Danieli Technology, AISTech 2014 Iron & Steel Technology, Indianapolis IN, USA, May 2014

BIBLIOGRAPHY

[1] M.MIYAJIMA, The Japanese Experience in Design and Application of Seismic Grade Rebar, Proceedings of Int. Seminar on production and Application of High Streng-th Seismic Grade Rebar Containing Vanadium, page 12, Beijing 2010

[2] S. MORITA, S. HITOSHI, Development of high strength mild steel deformed bars for high performance reinfor-ced concrete structural members, paper No. 1742. 11th world conference on Earthquake Engineering, ISBN: 0 08 042822 3, Elsevier 1996

[3] WJE Wiss, Janney, Elstner Associates, Inc., Mechanical Properties of ASTM A1035 High Strength Steel Bar Rein-forcement, Final Report WJE No. 2008.9901.0, 2008

Acciai


On the aging of a hyper-eutectic Zn-Al alloy

A. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano

Keywords: Zinc alloys - Aging - Die-casting - Mechanical properties

A. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano Università degli Studi di Brescia - DIMI

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Zinc alloys are widely used in different fields, like handles and locking, fashion and design as well as automotive or electronics, thanks to their good mechanical and technological properties combined with low cost and easy

formability. A limit to a wider use of these alloys is the aging phenomenon that causes a drop in their mechanical properties in time. In order to improve their use in competition with more expensive copper and aluminum alloys,

in the last years the research has been addressed to develop new Zn-alloys compositions. One of these new alloys, containing 15 wt% of Al and 1 wt% of Cu, appears to be suitable for both foundry and plastic deformation forming processes, as resulted from preliminary laboratory and industrial trials. Being a newly developed alloy,

many properties have still to be investigated, to better understand the effective potentiality for a proper industrial application. In this paper the ageing behavior of die-cast Zn-15Al-1Cu hyper-eutectic alloy was studied by means

of tensile tests and microstructural analyses. It was demonstrated that the alloy suffers from a drop in mechanical properties, in particular at the very beginning of soaking at high temperature. A first analysis of the microstructure by optical and scanning electron microscope was not able to fully point out the causes of the aging phenomenon

INTRODUCTION

Zinc aluminum alloys are used for the production of several functional items like hinge and small gears, and also design parts as handles, taps and fittings. These materials, in fact, guarantee a good compromise between performances and production costs. They offer good corrosion and wear resi-stance. Additionally, zinc alloys can be coated with all the traditional deposition techniques improving their aspect according to the specific application [1].Nowadays zinc aluminum alloys are also good candidates for the all the application fields where moderate mecha-nical resistance is required, for example as a substitutive material for bronze and brass components in home furni-shing and fashion parts. Zn-alloys can be also used for the production of bearing and bushings as they exhibit for such application high hardness, wear resistance similar to that of bronze and many properties comparable to those of cast steel [2-3]. Moreover these materials have a good machinability [3] that makes easy the manufacture of finished components with the correct tolerances.The presence of aluminum, in proper percentage, pro-motes the fluidity of the alloy, increases the mechanical properties and enhances the corrosion resistance in mild aggressive environments. Copper, in percentage between

1 and 3 wt%, improves hardness, tensile strength, wear resistance and creep behavior of the material [4].The main restrictions in using zinc alloys in structural ap-plications are the high density (more than twice of alumi-num and similar to that of steel) and the considerable drop of mechanical properties with temperature [5-6]. For this last reason the threshold temperature for the use of zinc alloys is commonly fixed at 80°C. A further problem affecting zinc alloys is the formation of instable phases during solidification and cooling, able to evolve in time, also at temperature below 100°C [3]. Solid state transformations through much stable configurations can occur, causing a decrease of mechanical properties in time [5]. This drop of zinc alloys properties is normally known as aging. Studies on conventional Zn-Al alloys aging showed that the phenomenon is enhanced by temperature. In parti-cular, mechanical tests performed maintaining sample at 105°C for different times, showed that the maximum drop of properties occurs in the first 24h while for longer soaking times, the decrease is slower and the mechani-cal properties tend to set at constant values [6]. Based on these findings some components are artificially aged (24h at 105°C in the case of automotive parts [5]) before their use, in order to stabilize the properties and avoid further drop during the component life.Fig. 1 shows the Zn-Al binary alloys phase diagram, whe-re the main families of nowadays used Zn-alloys are hi-ghlighted. The most common alloys are the so called Zamak (Zamak2, 3 and 5), characterized by an aluminum content


Memorie

of 4% and a copper percentage in-between 0 and 3wt%. The low melting temperature of Zamak, close to 390°C, allows the use of hot chamber high pressure die-casting (HPDC). The hexagonal close packed lattice of these alloys makes them scarcely deformable and not suitable for hot stamping [7].Other commercial alloys with higher aluminum and copper content are those belonging to the ZA family (ZA8, ZA12, ZA22, and the most widely used ZA27), which are charac-terized by higher mechanical properties than Zamak family [4]. The melting temperature of these alloys is higher than Zamak ones (over 410°C) and, for this reason, they should be cast only in cold chamber HPDC or in gravity or sand casting, with higher cycle time and costs.Recently, a new zinc alloy with an aluminum content of 15wt% has been developed. This alloy is characterized by good mechanical properties and high corrosion resistance; it is also suitable for foundry (hot chamber HPDC) as well as hot stamping processes [6]. However, many properties are still not known or under investigation. The aim of this paper is to study the effect of artificial aging, carried out at different temperature and for different holding times, on the tensile properties of this new alloy.

MATERIALS AND METHODS

Tensile test samples in Zn-15Al-1Cu alloy were produced by hot chamber HPDC with a 500 ton machine. The alloy was injected into the die at a temperature of 490 ± 10°C, while mold temperature was set at 250°C. The obtained samples have a round section with a diameter of 9 mm. Tensile tests were carried out at room temperature accor-ding to UNI EN ISO 6892 standard with a INSTRON 8501 machine, set in displacement control mode, with a cross head speed of 0.5 mm/min. Displacement was monitored by an extensometer.Three samples per condition were tested.Moreover, on one sample for each aging condition, the ela-stic modulus was determined with five load-unload cycles in the elastic field, setting a test speed of 0.1 mm/min. The evolution of the mechanical properties during aging was studied at three different temperatures: 80, 105 and 130°C, and for different holding times up to 240h, as re-ported in Tab. I. In the time between the alloy production and the aging tre-atment, all the samples were maintained at a temperature of -14°C in a freezer, in order to avoid any natural aging. The “as-cast” samples were tested just after production to evaluate the mechanical properties in not aged conditions (0 hours of aging time, as referred in Tab. I).Three groups of samples were considered, as a function of the heat treatment temperature. The temperature of 105°C (group A) was selected according to literature fin-dings; in fact, Leis et al. evidenced that, for conventional zinc alloys, the Ultimate Tensile Stress (UTS) of specimens aged at this temperature is similar to that of samples natu-rally aged for one year [4]. Additionally, as already mentio-

Fig. 1 - Zn-Al phase diagram [8].

ned, many car producers suggest a pre-treatment of zinc components at 105° for 24 hours before their use, in order to stabilize the mechanical properties [5].The other two temperatures investigated in the present research were chosen 25°C above and below the refe-rence temperature of 105°C. The temperature of 130°C (group B) was set to study the drop of properties at higher temperature, that probably happens with a faster kinetics. On the other hand, the temperature of 80°C (group C) is commonly accepted as the upper limit for the use of zinc alloys [9].

Time [h] Time [h] Time [h]

GROUP A:

T = 105°C

0

GROUP B:

T = 130°C

0

GROUP C:

T = 80°C

01 2 12 8 24 24 48 168 815 240 2424 16872 240

168240

Table. I - scheme of the tensile tests.

It is worth noting that the maximum aging time of 240 hours was chosen to be ten times higher than the one con-sidered to be necessary to reach almost stable mechani-cal properties for the traditional zinc alloys [6].In addition to the mechanical characterization, a micro-structural analysis was performed on all samples polished and etched with Nital 2% by means of a Reichert-Jung MeF3 optical microscope, equipped with the Leica QWin image analyzer software. The chemical composition of Zn-Al phases was assessed by means of the Oxford Energy Di-spersive Spectroscopy (EDS) microprobe, coupled to the LEO EVO 40 Scanning Electron Microscope (SEM).


RESULTS

TENSILE TESTS

Table II shows the values of the UTS and Yield Strength (YS) of Zn-Al alloy samples artificially aged at different ti-mes and temperatures.Fig. 2 shows the behavior of UTS as a function of the aging time. Similarly to what reported in literature for other zinc alloys [3], the new composition suffers from a strong drop of the mechanical properties in the first 24 hours of aging. Results from samples aged for longer time, however, reve-al a further, even though slower, reduction of the UTS. The behavior of the YS is similar to that of the UTS.For the same aging times, group B samples, i.e. those tre-ated at higher temperature, show lower UTS than group A samples and the reduction of the mechanical proper-ties occurs faster. Comparing the drop of the UTS, group A samples has a drop of 6% in the first 24 hours and 11% after 240 hours. For the group B samples the reduction is 10% in 24 hours and 21% after 240 hours. Finally, in the case of samples C, aged at 80°C, the maximum drop of the UTS after 240 hours is only 4%.As shown in Fig. 2, the aging at 105°C for 240 hours gives similar UTS and YS values to those obtained at 130°C for 24 hours. From this experimental result, it follows that the same mechanical properties can be achieved lowering the aging time and increasing conveniently the heat treatment temperature. Such aging conditions can be reasonably ac-cepted only if the cast part is not affected by high levels of porosity or other defects that can be modified or worsened by a long stay at high temperature. It should be also taken into account that a soaking at high temperature usually cau-ses a a coarsening and globularizarion of the precipitates.Finally, a last aging test was performed at 105°C for 720 hours (30 days), aimed at evaluating if a steady state con-dition of mechanical properties can be finally reached. The UTS and YS measured after this heat treatment were 245 MPa and 193 MPa respectively, which are values close to those obtained after the aging at 130°C for 240 hours, confirming the achievement of almost stable conditions.

YS [MPa] UTS [MPa]GROUP A B C A B C

t [h] T=105°C T=130°C T=80°C T=105°C T=130°C T=80°C0 271±6.4 313±7.41 261 ± 1,4 251 ± 4,2 300 ± 7,2 300 ± 0,92 248 ± 1,0 242 ± 1,5 255 ± 1,4 310 ± 6,6 306 ± 6,7 312 ± 4,64 248 ± 1,0 246 ± 0,7 310 ± 6,6 304 ± 0,48 245 ± 3,5 239 ± 3,6 246 ± 0,8 304 ± 0,4 299 ± 3,1 295 ± 6,015 246 ± 6,4 296 ± 3,324 236 ± 0,1 225 ± 0,7 290 ± 4,5 295 ± 5,5 281 ± 1,3 290 ± 2,872 238 ± 4,6 210 ± 3,5 290 ± 3,8 271 ± 2,4

168 228 ± 4,7 196 ± 2,1 249 ±5,1 281 ± 2,5 252 ± 3,6 298 ± 4,8240 224 ±3,6 192 ± 1,2 246 ± 1,4 279 ± 3,5 248 ± 0,4 301 ± 6,3

Table. II - tensile tests results.

Fig. 2 - Ultimate tensile strength of Zn-Al samples versus aging time.

A conclusion that can be derived from this result is that a heat treatment at 105°C for 24 hours is not enough to reach a complete settlement of mechanical properties for a Zn-15-Al-1Cu alloy. To foresee the behavior of a component desi-gned with this alloy for long service applications, the aging treatment should be performed at higher temperatures.Additionally, tests carried out on samples aged at 80°C proved that the drop of mechanical properties is strongly slowed down by the reduction of the treatment temperatu-re, confirming that 80°C can be assumed as the threshold temperature also for this alloy

MICROSTRUCTURE

In order to better understand the aging phenomenon, a metallographic characterization of as-cast and aged sam-ples was performed by means of optical microscope. Ac-

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Memorie

cording to the phase diagram (Fig. 1), the microstructure of the Zn-15Al-1Cu alloy should be composed at room temperature by Al-rich primary phase and Zn-rich eutectic, both transformed by the eutectoid decomposition. Figure 3 shows the microstructures of samples in different conditions: as-cast (a), aged for 240 hours at 105°C (b) and at 130°C (c). The Al-rich primary grains (dark areas) are surrounded by the eutectic phase (light gray areas). No particular differences can be noted between samples, in terms of size, shape or amount of the different phases. No recrystallization phenomena are detectable.All samples are almost free from porosity, regardless of the aging conditions, which means that the reduction of the mechanical properties cannot be ascribed to an incre-ase of defects, as can easily occur in die-cast components after long stays at high temperature.Preliminary SEM analyses were carried out on the samples in the as cast condition (Fig. 4a) and heat treated for 240 h at 130°C (Fig. 4b). Both samples show the eutectoid de-composition of Al-rich primary grains with the consequent formation of a thin lamellar microstructure, surrounded by the Zn-rich eutectic. The Al content of primary dendrites is close to the theoretical value of 32.3% expected from Zn-Al phase diagram. This suggests that notwithstanding the ra-

Fig. 3 - Microstructure of Zn-15Al-1Cu alloy die-cast samples a) “as-cast”, b) after 240 h of aging at 105°C and c) after 240 h of aging at 130°C.

pid cooling imposed by the production process (HPDC), the microstructure is not far away from the equilibrium state. Comparing figure 4a and figure 4b some small differences can be appreciated in the microstructure. In particular, the aged sample shows more globular and coarser precipita-tes at the interface between the phases and the eutectoid lamellae seems to be better defined, maybe as a conse-quence of a more distinct separation between Al and Zn elements. Such rearrangement of the microstructure could justify the drop of mechanical properties measured by tensile tests on aged samples. Increasing the temperature, in fact, the atomic mobility is enhanced and solid state diffusion and solute redistribution are promoted. The achievement of a pseudo-equilibrium state at the end of the aging treatment should allow the formation of a structure free from super-saturated solutions, having a lower local lattice deforma-tion. This condition causes a reduction of barriers to the dislocation movement and a modification of the free mean path and, consequently, a lower mechanical resistance. Unfortunately, such modifications are not detectable by optical or electron microscope and also the spatial reso-lution of EDS microanalysis is inadequate to measure a solute redistribution on so small distances. Transmission


Electron Microscope (TEM) measurements and X-ray dif-fraction (XRD) experiments are desirable to overcome such limitations and explain the results of tensile tests in terms of the alloy microstructural changes.

CONCLUSIONS

In this paper the aging behavior of the hypereutectic Zn-15Al-1Cu alloy was investigated.Being a newly developed alloy, no data are up to now avai-lable in literature about natural or artificial aging. Tensile tests were performed on samples aged at three tempera-tures (105, 130 and 80°C) for different soaking times (up to 240 hours and in one case up to 720 hours) to measure the decrease of mechanical properties.This characterization gave useful information about the potential use of this alloy at relatively high temperatures and allowed to assess the mechanical properties close to steady state conditions.The results show that in the first 24 hours a fast decrea-se of mechanical properties occurs at all the investigated temperatures. After this period the reduction of mechani-cal strength slows. As expected, the aging phenomenon is accelerated by increasing the aging temperature: the reduction of the UTS after 240 hours was 4% at 80°C, 11% at 105°C and 21% at 130°C.Microstructural analyses carried out by optical microsco-pe and SEM-EDS revealed only small differences between the aged and the as-cast samples. In particular, the aged samples showed a slight coarsening of precipitates at the phases boundary and the eutectoid lamellae appea-red better defined, suggesting the redistribution of solute elements. Such microstructural modifications could justify the reduction of mechanical properties measured by tensi-

Spectrum Al Cu Zn

1 36.48 0.17 63.34

2 8.98 0.55 90.47

Spectrum Al Cu Zn

1 32.03 0.39 67.58

2 2.11 0.82 97.07

Fig. 4 - SEM analysis performed on the as cast sample (a) and sample treated for 240 h at 130°C (b).

le tests, anyway. To examine more deeply this hypothesis further investigations with other techniques like TEM and XRD are under development.

ACKNOWLEDGMENT

The authors want to thank Mrs. Valentina Ferrari for the support in the execution of tensile tests.

REFERENCES

1. D. APELIAN, M. C. Donald HERRSCHAFT, Casting with zinc alloys, JOM 33 (11), 1981, 12-20.

2. K.J. ALTORFER, Zinc-alloys compete with bronze in bea-rings and bushings, Metal Progress 122 (6), 1982, 29-31.

3. E. J. KUBEL, Expanding horizons for ZA alloys, Adv. Mat. Proc. inc. Material Progress, 1987, 51-57.

4. Y. H. ZHU, W. B. LEE, S. TO, Ageing characteristics of cast Zn-Al based alloy (ZnAl7Cu3), J. Mat. Sc. 38 (9), 2003, 1945-1952.

5. A. BUCCIOL, Aging behaviour of zinc die casting alloy ZP0810 (Master thesis), Univesity of Padova, 2012.

6. W. LEIS, L. KALLIEN, Ageing and creep of Zinc-Diecast alloys, Int. Zinc Diecasting Conference 2013 “Tradition & Innovation”, Praha 13-14 June 2013.

7. A. POLA, L. MONTESANO, R. ROBERTI, Nuove leghe di zinco per l’industria del design, Proceedings 33° Con-vegno Nazionale AIM, Brescia 10-12 November 2010.

8. Alloy Phase Diagrams, ASM Metals HandBook Vol. 3, ASM International 10 ed., 1992.

9. A. POLA, R. ROBERTI, D. ROLLEZ, Primary and steady state creep deformation in Zamak5 die-casting alloy at 80°C, Mat. Charact. 59 (12), 2008, 1747-1752.

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Implementation of an open-die forging process for large hollow shafts for wind

power plants with respect to an optimized microstructure

M. Wolfgarten, D. Rosenstock, L. Schaeffer, G. Hirt

Keywords: Open-die forging - Microstructure - Process optimization

M. Wolfgarten, D. Rosenstock, G. HirtInstitute of Metal Forming,RWTH Aachen University

L. SchaefferLaboratório de Transformação Mecânica, Universidade

Federal do Rio Grande do Sul, Brazil

Forgiatura

To realize large wind power plants in an economically feasible way, it is necessary to identify potential for lightweight design of the generator hollow-shafts, which are commonly produced by casting. The weight of these shafts can significantly be reduced by producing them by open-die forging, since the forming of the

material leads to a higher strength, which allows to reduce the wall thickness noticeable. This paper describes the development and implementation of a forging process for hollow shafts with respect to an optimized microstructure. To numerically investigate this process, a realistic finite element simulation model was

developed in a first step. The kinematic of the tools has been implemented authentically to provide a realistic material flow and process conditions. Additionally, a material model for the steel 42CrMo4 was integrated into

the simulation model to predict the resulting microstructure. Using the implemented FE model, the forging process was optimized manually to achieve a homogeneous and fine-grained microstructure. The optimization

was based upon a variation of different forging parameters and the sequence of forging steps. In the next step, a forging on laboratorial scale was performed to validate the simulation model. For this purpose, after

forging, specimens from the hollow shaft were evaluated by metallography to determine the final grain size. A comparison of the results with the numerical simulation showed a general agreement of the measured grain size with the numerically calculated grain size. Based upon these results, the process model was transferred

to an FE model with an industrial scale. By this it was possible to analyze the transferability of the used FE model regarding the assumptions about the kinematics and the sequence of the forging steps. A numerical

investigation of the industrial process proved the scalability of the process to an industrial relevant geometry.

INTRODUCTIONInitial Situation During the last years the importance of alternative energy sources like wind energy has risen constantly. To meet with the further anticipated demand for renewable energy, the importance of wind energy will be growing even more as this currently is the cheapest and most effective renewable energy source [1]. A higher energy production by wind power plants cannot only be achieved by setting

up additional wind parks, but also by an increase in the performance and by this the size of newly built plants. This requires larger machine parts, like rotors, shafts, gears and generators, leading to higher requirements for the tower’s construction. To cope with these aspects, the reduction of the nacelle’s weight at the top of the tower offers a good opportunity. One approach is to replace the commonly cast generator shaft by a forged hollow shaft with excellent mechanical properties. In comparison to a cast hollow shaft, a forged shaft could offer a higher strength and therefore would allow reducing the wall thickness significantly. Recker et al. [2] estimated that producing the hollow shaft by open-die forging could allow a weight reduction of up to 60% compared to a cast hollow shaft.

State of the ArtOpen-die forging is mainly used for the production of high quality parts in low quantities. These parts are applied for


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highly loaded purposes in large machines like generator shafts or rolls. Hollow parts like rings are usually forged by applying an upper-die and a mandrel. Usually, this kinematic is used for forging rings and intends to increase the diameter. For the production of a hollow shaft, which has a much higher length and smaller diameter, the main challenge is to ensure a sufficient axial material stretching. The longitudinal material flow is predominantly achieved by improving the tool’s shape and the process kinematics [3]. The best axial stretching can be realized using concave or v-shaped dies. However, two concave dies are lacking flexibility during the process according to the forgeable diameter as the shape of the dies limits the possible workpiece geometries. Two v-shaped dies are disadvantageous for the process handling, since this tool combination impedes the loosening and extracting of the mandrel after a forging heat. Based upon numerical investigations and literature review [4], the combination of a flat and a v-shaped die proved best to realize a sufficient axial stretching.

Motivation and ScopeThe overall objective of the investigation presented in this paper is the optimization of the open-die forging process for hollow shafts with respect to an optimized microstructure and transferring the results to an industrially relevant geometry. As a first step, the main focus was the development of a simulation model for the forging of a 150 kg workpiece under consideration of the kinematic and tool geometry. The second main objective consisted of optimizations of the microstructural properties meant to ensure the final product’s high mechanical strength. This requires a homogeneously distributed, fine grain size along the workpiece, which is likewise supported by a homogeneous strain. So a manual optimization through variation of the bite ratio and height reduction was performed. To verify the results from the numerical simulation, especially in terms of microstructural evolution, forging of a 150 kg workpiece was performed at the IBF. The corresponding results from the experiment were used to validate the numerical simulation model according to the microstructural evolution, tool geometry and kinematics. Based upon the validated simulation model, in a third step, the process was transferred to an industrially relevant geometry by numerical investigation.

METHODS AND PROCEDURE

Numerical Simulation ModelFor correctly predicting the material flow and the microstructural evolution, a simulation model was implemented in Transvalor Forge 2011, whose boundary conditions and kinematics coincide with the forging conditions in reality. The general requirements on the process kinematic for a realistic simulation model are:1. The upper flat die is moving in y-direction and

performs the reduction of the material.2. The lower v-shaped die is fixed in its position and not

moved during the whole process.3. The mandrel supports the hollow shaft during the

process and is held by a manipulator. When the upper die presses in y-direction, the flexibly supported mandrel can move freely in all directions.

The grain size calculation was realized by implementing a microstructure model into the integrated calculation in Forge 2011, based on experiments at IBF and literature [5-8], which showed good accordance with the behavior of the used material.

The numerical simulation of hollow shaft forging is more complex compared to conventional open-die forging due to positioning of the workpiece over a mandrel and the attachment of the mandrel to the manipulator. This setup and the long process time increases the complexity and calculation time of the simulation significantly. The simulation of the reference process described within the paper requires a calculation time of one week on a quad core Intel Xeon workstation. Furthermore, the handling of the workpiece in the simulation is impeded since in hollow shaft forging the workpiece is just indirectly positioned during the process. Sliding of the hollow shaft leads to an inexact positioning. Hence it is not possible to simulate the whole forging process in one simulation. The simulations needs to be interrupted after every pass to control the exact positioning and kinematic.

Chosen Strategy for Hollow Shaft ForgingDifferent parameters influence the forging of a hollow shaft. Firstly, the main influencing parameters are standard parameters for open-die forging, the height reduction εh and the bite ratio sB/h0, which describe the ratio of the contact length between die and workpiece and the initial height (here: initial diameter). The variable parameters are the direction of forging and the combination of forging strokes and rotations. According to the forging direction, two different possibilities exist. At first, the workpiece can be forged from the front of the mandrel towards the manipulator tong. After one rotation has been forged, the manipulator feed is executed and the next rotation is forged.The second possibility consists of forging the opposite direction. As described before, after the forging of one

Fig. 1 - Process setup – hollow shaft forging


rotation, the manipulator feed is executed and the next rotation is forged.Besides the forging direction, the combination of rotation and strokes can be varied as second possibility to mainly influence the forging process. The investigations described within this paper are based on the strategy to forge one rotation at first, translate the workpiece by the manipulator feed and forge the second rotation in the following step. Figure 2 visualizes this process principle. As shown in the top, the whole circumference of the hollow shaft is forged in the first step, which requires 10 strokes or a whole workpiece rotation. After each stroke, the mandrel is rotated by 90°, but just after the fifth stroke of each rotation the mandrel is once rotated by 45°.

Process optimizationGenerally, the bite ratio (quotient of the bite length and initial height of the workpiece – sB/h0) and the height reduction εh can be identified as the forging parameters, which most decisively influence the strain distribution and grain size. Therefore to optimize the process parameters for hollow shaft forging, numerical studies of different bite ratios (0.3, 0.5, 0.7) and height reductions(10%, 20%) were performed and the resulting equivalent plastic strain and average austenitic grain size were analyzed. The optimization was performed for a hollow cylinder with a diameter of 240 mm at an initial temperature of 1200 °C. For the forging on a laboratorial scale, this corresponds to the geometry for the middle steps of the hollow shaft.As due to the long simulation time the process could not be optimized completely, three different process routes were used to investigate the process optimization. The optimization was performed for the forging of two rotations, each for one and two passes. The equivalent strain and the grain size according to the optimization are evaluated along three lines at half of the wall thickness, each 120° distributed over the circumference, see Figure 3.Table 1 summarizes the influence of the bite ratio, height reduction and number of forging steps on the strain distribution in the workpiece. An increase of the bite ratio from 0.3 to 0.7 leads to an average increase of strain of 26% for one pass and 50% for two passes at half of the wall thickness. The standard deviation of the strain distribution can be reduced by 46% for one and by 14% for two passes and thus mainly increases the strain homogeneity. Similar effects can be observed for a higher height reduction.

Fig. 2 - Strategy for the forging of one rotation [2]

Fig. 3 - Points for evaluation of strain and grain size distribution

εh/Number of Passes

ParameterBite ratio (sB/h0)

0.3 0.5 0.7

10% / one pass 0.46 ± 0.13 0.56 ± 0.13 0.58 ± 0.07

20% / one pass 0.54 ± 0.11 0.97 ± 0.14 1.03 ± 0.13

10% / two passes 0.75 ± 0.25 1.03 ± 0.14 1.13 ± 0.14

Table 1 - Influence of bite ratio, height reduction and forging steps on the equivalent plastic strain distribution (average and standard deviation)

So for the optimization of the strain, a higher bite ratio and height reduction should be preferred to achieve the intended results. The grain size as second optimization objective is mainly influenced by the temperature, strain and strain rate. Since the minimum warm forming temperature for 42CrMo4 is 850°C, the workpiece temperature should not drop below this point. So to allow a long enough time frame for forging the initial temperature is set to 1200°C.

Forgiatura


Memorie

Therefore, a variation of the temperature to optimize the microstructure is nearly impossible as always the maximum temperature of 1200°C has to be chosen for the beginning of the forging process. The strain rate results from the tool speed of 40 mm/s and the geometry of the work piece. Table 2 visualizes the influence of the forging parameters on the average austenitic grain size evolution during the forging process. It can be concluded that a large bite ratio in combination with a high height reduction is preferred to yield a fine grained and homogeneous microstructure.From this case study, a process route with a higher bite ratio and a sufficiently height reduction should be preferred to both optimize the strain distribution and microstructure in the hollow shaft. Furthermore, the numerical simulations proved that a height reduction of 20% and more leads to an increase of the inner diameter and deviations from the intended shape. Therefore, the height reduction during the process should be lower than a value of 20%.Besides a small and homogeneous microstructure along the workpiece, likewise a homogeneous distribution over the cross-section is preferred. As example, Figure 4 shows the grain size distribution after the forging of one rotation. It can be concluded that the chosen forging strategy is advantageous to achieve a homogeneous strain distribution over the cross-section. The strong deviation in two points is probably caused by numeric irregularities and disagrees with the general distribution at the cross section.

Properties optimization

ParameterBite ratio

0.3 0.5 0.7

10% / one pass 86 ± 21 µm 56 ± 21 µm 55 ± 20 µm

20% / one pass 82 ± 31 µm 49 ± 16 µm 40 ± 16 µm

10% / two passes 63 ± 17 µm 47 ± 17 µm 30 ± 6 µm

Table 2 - Influence of parameters on average austenitic grain size (average and standard deviation), initial grain size 200 µm

Fig. 4 - Exemplary grain size distribution at the cross section

Fig. 5 - Geometry in different stages of the reference process, all dimensions in mm

Reference forging processThe reference forging process was used to validate the numerical simulation model and in particular the microstructure calculation. Therefore the process was designed such that an inhomogeneous grain size distribution should be obtained in the final workpiece to validate the microstructure calculation on a preferably wide range. For that purpose, sections with different outer diameters were forged using different reductions in diameter and varying reheating conditions. The geometry for the laboratorial forging was limited by the maximum possible length of 490 mm and a weight of 150 kg for the initial workpiece. The development of the shape of the


workpiece during forging is shown in Figure 5. The forging was performed using a 6.3 MN hydraulic press and a 6-axis forging robot. The workpiece was initially heated to 1200 °C and the tool speed was set to 40 mm/s.

RESULTS

Results of the forging processComparing the final shape, material flow and especially the grain size of the experiment to the numerical simulation allows judging the quality of the simulation model. In this context, the validation of a simulated grain size by metallographic analysis is limited since only the final state can be investigated. So it is not possible to analyze the grain size evolution during intermediate steps of the forging process. Figure 6 visualizes the results of the metallographic analysis. Generally, the microstructure is only investigated in one single center point per step and could possibly vary due to different conditions in a step. The metallographic analysis shows that in the largest step (I) the microstructure is just partially recrystallized and shows still similarities to the initial microstructure. This effect results from the low strain, which just has been imposed by the 10 mm diameter reduction during the first pass. In the other three

Fig. 6 - Exemplary results of microstructure investigation for each step

steps of the shaft, the effect of the higher strain clearly becomes visible as the grain size becomes smaller in a fully recrystallized microstructure. Under consideration of the measuring error, the average austenitic grain size for each step is decreasing from 142 µm in step I to 37 µm for step IV with the smallest diameter. The general expectation from the optimization, that a higher strain leads to a finer

Step 1* 2 3 4Metallography 142 ± 89 µm 102 ± 21 µm 110 ± 38 µm 37 ± 14 µm

Simulation 167 µm 163 µm 84 µm 6 µmAbsolute Deviation

in Percent17,6% 59,8% 23,6% 83%

Table 3 - Comparison of numerical and real grain size

grain size and recrystallized microstructure, is fulfilled.According to the numerical calculation, during the reheating process, the grain size increases rapidly up to approx. 500 µm due to grain growth, whereas the dynamic and static recrystallization phenomena during and after a strain increment result in a reduction of grain size. Considering the standard deviation of the metallographic grain size measurements, the grain sizes show roughly the same tendency as summarized in Table 3.A correlation between the strain and the grain size can generally be observed, showing that a higher strain leads to a finer grain size in the workpiece. However, this needs always to be regarded in relation to the process history, since significant grain growth can appear during a necessary reheating process. For this kind of process, an exact prediction of microstructure proves as difficult due to the long process time and the semi-empirical modelling of the microstructure evolution. Nevertheless, the microstructure calculation reproduces the influences on microstructure during forging in a satisfying way.

Transfer to Industrial ScaleThe verification of the simulation model showed in principle a satisfying accordance between the real process and the numerical model. Based upon this, the simulation

* The high standard deviation is resulting from the occurrence of non-recrystallized grains.

Fig. 7 - Industrial hollow shaft geometry, all dimensions in mm

is scaled to a 20 ton geometry (Figure 7), based on [9], in order to prove the process design on a large scale. A general difference between forging on laboratorial and industrial scale consists in the thermal development during the process. Due to the significantly larger volume and a smaller ratio of surface to volume, the workpiece cools down considerably slower. Therefore, the number of necessary heats can be reduced from four heats for the 150 kg part to just two heats for the 20 t part.

Fig. 8 - Strain evolution for a hollow shaft with an industrial geometry

Forgiatura


Memorie

According to the strain distribution in a longitudinal section of the workpiece, as shown in Figure 8, the chosen forging strategy leads to a homogeneously and sufficient equivalent strain in the three smaller steps of the workpiece with values of εV>5 in average. Only in the first step (largest diameter), the average strain reaches a level of εV just above 2. Nevertheless, a homogeneously distributed average equivalent strain of 2 is sufficient to enable good mechanical properties.Figure 9 visualizes the grain size distribution in a longitudinal section. The 2nd, 3rd and 4th step of the workpiece are showing a fine and homogeneously distributed grain size between 5 µm and 40 µm. In contrast, the largest step of the hollow shaft has a final grain size up to 1200 µm. The reason for this behavior is based on the chosen process sequence. The workpiece needs to be reheated to forging temperature of 1200 °C, after the temperature has dropped below the minimum forging temperature of 850°C. As the step I has already been forged in the first heat and is not forged any more after reheating, the resulting microstructure is mainly driven by grain growth, which leads to an enormously high grain size of up to 1200 µm. In contradiction the grain size of step II, III and IV is mainly influenced by the forging during the last heat. As shown in Figure 9, the forging on industrial scale results apart from step I in a more homogeneous grain size than the forging on laboratorial scale.

CONCLUSION

This paper proves by numerical simulation and experiments that open-die forging can generally be used for the process design of hollow shaft forging. The experimental forging of a 150 kg part based on numerical simulation studies showed that the developed kinematic and pass schedule are suitable. The development of the experimental forging process was based on an optimization by variation in the numerical simulation. The optimization showed that high bite ratios of 0.7 and height reductions of 10-20% should be preferred

Fig. 9 - Grain size evolution for a hollow shaft with an industrial geometry (Reheating after 3rd pass)

for an optimized distribution of strain and grain size in the workpiece.In order to validate the numerical simulation model, a 150 kg hollow shaft was forged and compared to an identically simulated numerical process. A comparison of the numerically predicted and the experimentally measured grain size showed a general accordance between both approaches. However, due to the long process time and the semi-empiric microstructure modelling in the numerical simulation, some uncertainties occurred, leading to an average deviation of 46% between the numerical and experimental results. Since in general the trend of the numerical simulation could be observed in the experiment, the results can be considered in principle as sufficient.Based on this, the process design of the 150 kg shaft could be transferred to an industrial scale. A numerical investigation of the industrial process proved the scalability of the process to an industrial relevant geometry. Both in laboratorial and industrial scale, the grain size evolution is mainly influenced by grain growth. In the industrial scale, the resulting grain size distribution is much more homogeneous as less reheating steps are required and by that, the influence of grain growth on the grain size is reduced.

OUTLOOK

The manual optimization presented in this paper only • considers the optimization of few rotations and passes. To perform a more complex optimization, it is necessary to regard the whole process. Possible approaches are a fully automatic optimization through the variation of different process parameters or a pass-wise optimization based upon the main values for the investigation of the microstructure as strain, strain rate and temperature. As a fully automatic process optimization in numerical simulation is too complex regarding simulation time, a fast empirical calculation model could be implemented and used to optimize whole processes. First works to this direction for squared ingots are presented in [10].


A further degree of freedom to influence the workpieces • final properties is the initial geometry. By a variation of the height to diameter ratio, the geometrical and microstructural evolution are influenced significantly.

The investigation of the grain size evolution showed • that grain growth during reheating has the dominant influence on the final grain size. Therefore an optimization approach could be the adaption of the process route, so that in the final heat the whole workpiece is forged in order to reduce the large grain size after grain growth.

While lowering the reheating temperature is not an • option for the 150 kg workpiece, this might be a solution for the 20 ton ingot. As the numerical simulation is validated, it could be used to find the optimal reheating temperature.

ACKNOWLEDGEMENTS

The authors thank the Deutsche Forschungsgemeinschaft and CAPES for the financial support within the project “Forged hollow shafts for power plants” in the “Brazilian German Collaborative Research Initiative in Manufacturing Technology” (BRAGECRIM).

REFERENCES

“Renewables Global Status Report”, Ren21 (Renewable 1] Energy Policy Network for the 21st century) UNEP (United Nations Enegry Program) , Paris, 2012

D. Recker, M. Franzke, G. Hirt, “Forged hollow shafts 2] for wind power drives”; 1. Conference for Wind Power Drives, Aachen (2013), pp. 199-214G. Hirt, D. Schäfer, M. Franzke, “Hohlwellen für 3] Windkraftanlagen – Prozessauslegung anhand von FEM-SImulationen”, in: Industriemanagement, 2/2011G. Spur (Editor): Titel: Handbuch Umformen, 1st Edition, 4] 2nd Volume, p. 230 – 232, Hanser Verlag, 2012Y.C. Lin, M. Cheng , J. Zhong , “Study of static 5] recrystallization kinetics in a low alloy steel” in “Computational Materials Science”, Edition 44, Page 316-321, 2008Y.C. Lin, M. Cheng, „Study of microstructural evolution 6] in a low alloy steel“ in “Journal of Material Science” Edition 44, Page 835-842, 2009Y.C. Lin, M.Cheng, “Numerical simulation and 7] experimental verification of microstructure evolution in a three-dimensional hot upsetting process” in “Journal of Materials Processing Technology” Edition 209, Page 4578-4583, 2009D. Qian, J.Guo, Y. Pan, “Austenite Grain Growth Behavior 8] of AISI 4140 Alloy Steel”, Wuhan, 2010T. Noack and S. Nelle:, Titel: Fehlerfreier Riese: Die 9] Rotorhohlwelle einer 5 MW-Windkraftanlage. in: Giesserei, 2008. 95(9): p. 30-34. 2008D. Rosenstock, D. Recker, M. Franzke, G. Hirt, D. 10] Sommler, K.-J. Steingießer, A. Tewes, R. Rech, B. Gehrmann, S. Kirchhof, R.Lamm, “Online-Visualisation during Open Die Forging and Optimisation of Pass Schedules” in “Steel Research International”, 2013

Forgiatura


Corso itinerante

SOLIDIFICAZIONE E COLATA CONTINUA

7-8-14-15-22 maggio 2015

Organizzato da:Centro di Studio Acciaieria,ASSOCIAZIONE ITALIANA

DI METALLURGIA

L’Associazione Italiana di Metallur-gia ha deciso, come da tradizione, di organizzare la nuova edizione del Corso sulla colata continua degli acciai per continuare a sostenere le imprese nell’azione di formazione del proprio personale, uno dei fattori per mantenere e migliorare la propria competitività. Data l’importanza dei semilavorati destinati alle operazioni di forgia, durante il Corso si affronte-ranno anche le problematiche relati-ve al colaggio dei lingotti e dei blu-mi di grande dimensione. Secondo la formula ampiamente collaudata e gradita ai partecipanti, la formazione viene svolta con due modalità: le-zioni 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 - in modo da poter osservare sul campo gli aspetti più significativi e peculiari riguardanti i sistemi per la colata continua.Per rispondere a queste esigenze il Corso è itinerante e le lezioni si svol-geranno presso alcune interessanti realtà produttive: Acciaieria Arvedi, Prosimet, Acciaieria di Calvisano, Du-ferdofin Travi e Profilati di Pallanze e Calderys.Dopo una sintetica introduzione di carattere storico il Corso abbraccerà i temi della solidificazione, le proble-matiche relative alla struttura della macchina, il colaggio dei lingotti, i componenti refrattari, le polveri di co-pertura, l’applicazione dei campi elet-tromagnetici, la difettologia, i modelli di simulazione, il colaggio di billette, blumi e bramme ecc..., solo per citare i temi salienti. In occasione delle visi-te i tecnici delle società ospitanti pre-senteranno gli impianti con una par-ticolare sottolineatura degli aspetti caratteristici di ogni sistema di colata.

Queste attività saranno affiancate ed integrate da interventi didattici tenuti da docenti del Politecnico di Milano, nonché da esperti di società di inge-gneria di riconosciuta esperienza. Gli interventi didattici e le visite tecni-che si articolano su un arco di cinque giorni e sono organizzati con cadenza tale da evitare ai partecipanti un’as-senza eccessivamente prolungata dalle proprie aziende.Il consistente numero di visite tecni-che, che sono state organizzate per migliorare la qualità del percorso cul-turale, costringono a limitare il nume-ro dei partecipanti, per cui si consi-glia di provvedere all’iscrizione il più presto possibile. Il Corso, coordinato da Ottavio Le-cis, Silvia Barella e Serena Fasolini, si svolgerà secondo il seguente pro-gramma:

giovedì 7 maggio (Cremona)

Gli aspetti principali del processo •di solidificazione: strutture di soli-dificazione e segregazioniModalità di colaggio•Visita agli impianti di Acciaieria Ar-•vedi

venerdì 8 maggio (Bottanuco BG)Polveri di colata continua: che cosa •sono, a cosa servono e come fun-zionano

SOMMARIO

VITA ASSOCIATIVA

Solidificazione e colata continua ..50

Scuola Metallurgia delle polveri .. 51

Giornata di studio

“La Metallografia passando dalla

preparativa metallografica” ...... 51

Failures nei refrattari................52

Prossime manifestazioni AIM ...53

Attività dei Comitati Tecnici ...................................54

Notizie da Unisider ....... 56

Refrattari nei sistemi di colata con-•tinuaMagnesita Navarras, Zubiri, Spa-•gnaRefrattari per la colata in sorgente•Presentazione di Prosimet•Visita agli impianti Prosimet•

giovedì 14 maggio (Calvisano BS)

Applicazioni di sistemi elettroma-•gneticiErgolines Lab, Trieste•Controllo di parametri di processo •in colata continuaPresentazione di Acciaierie di Cal-•visano (Feralpi Group)Visita agli impianti di Acciaierie di •Calvisano (Feralpi Group)

venerdì 15 maggio (Pallanzeno VB)

Il colaggio in blummi di gradi di-•mensioniPrincipi base per il colaggio dei lin-•gottiPresentazione di Travi e Profilati di •Pallanzeno (Duferdofin Nucor)

Atti e notizie


Vita associativa

Visita agli impianti di Travi e Profila-•ti di Pallanzeno

venerdì 22 maggio (Fiorano Mode-nese MO)

I difetti del processo colata conti-•nuaSoluzioni impiantistiche per elimi-•nare i difetti per i tipi di acciaioDanieli & C. Officine Meccaniche, •Buttrio

SCUOLA METALLURGIA

DELLE POLVERI

Imola, 9-10 giugno 2015

Nei principali paesi industrializzati la metallurgia delle polveri è ormai una tecnologia consolidata. Questa posi-zione riconosciuta favorisce l’organiz-zazione di corsi di formazione tecni-co-scientifica sui processi e sui requi-siti richiesti per una corretta scelta e una valida progettazione dei materiali e dei componenti sinterizzati.Le grandi organizzazioni internaziona-li di categoria (APMA - Asian Powder Metallurgy Association, APMI - Ame-rican Powder Metallurgy Association, MPIF - Metal Powder Industries Fede-ration, EPMA - European Powder Me-tallurgy Association) sono i maggiori enti che, per statuto, sono impegnati in queste attività, finalizzate a favori-re la diffusione delle conoscenze spe-cifiche e gli incrementi delle possibili utilizzazioni.In questo panorama, l’AIM, per qua-si trent’anni, dal 1960, organizzò dei corsi di formazione sulle possibilità e sugli impieghi della metallurgia delle polveri. Poi, una decina d’anni prima della fine del secolo, per mo-tivi sostanzialmente legati a forme di concorrenza esasperata fra gruppi di aziende del settore, quella bella tradi-zione fu interrotta. Grazie all’interesse e alla disponibilità dimostrati dalla SACMI, è stato possi-bile ripartire con questa bella iniziati-va. Dopo il successo di una prima edi-zione, nel 2012, l’Azienda di Imola ha proposto all’AIM di ripetere l’evento, allo scopo di favorire la diffusione di

conoscenze rigorose e aggiornate su una tecnologia competitiva, la cui af-fermazione può contribuire “nel suo piccolo” alla ripresa dello sviluppo economico nazionale.Il Centro Metallurgia delle Polveri dell’AIM e la SACMI si augurano che l’impegno profuso per l’organizzazio-ne del Corso – anche da parte dei do-centi, tutti specialisti del settore – sia adeguatamente riconosciuto attra-verso un’ampia partecipazione all’ini-ziativa di tecnici e studiosi interessati alla metallurgia delle polveri. La scuola si svolgerà secondo il se-guente programma:

9 GIUGNO 2015

Presentazione della SACMI • Presentazione del Centro Metallur-• gia Polveri dell’AIM Introduzione alla metallurgia delle • polveri Le polveri metalliche: produzione e • proprietàParticolari sinterizzati: fasi della • produzione Indicazioni sulla progettazione del-• le forme Nuove tendenze della pressatura – • eco & energy

Sinterizzazione e sinterotempra• Nuovi prodotti e futuri sviluppi in • campo automobilisticoLa precisione dimensionale dei par-• ticolari sinterizzatiRequisiti geometrici dei prodotti • sinterizzati multilivello in accordo con le nuove tecnologie del proces-so di calibraturaI materiali sinterizzati: influenza • della porosità sulle caratteristicheLe operazioni post-sinterizzazione• I cuscinetti autolubrificanti • Visita degli impianti SACMI • La normativa sui materiali sinteriz-• zati e gli acciai sinterizzatiTrattamenti termici e termochimici • dei sinterizzati ferrosi

10 GIUGNO 2015

Non destructive testing and com-• puterized vision in PM parts pro-ductionI materiali a base rame sinterizzati• Metal Injection Molding e/o addi-• tive manufacturingGli acciai inossidabili sinterizzati• Il controllo della durezza e gli esami • microstrutturali sui sinterizzatiGli acciai speciali da polveri• I metalli duri (carburi cementati)• La pressatura isostatica a caldo • Evoluzione delle polveri per la pro-• duzione di componenti sinterizzati. Ottimizzazione delle miscele ed elementi di lega

Per ulteriori informazioni rivolgersi alla Segreteria dell’Associazione Ita-liana di Metallurgia, oppure visitare il sito www.aimnet.it

Giornata di studio

LA METALLOGRAFIA ATTRAVERSO UNA

CORRETTA PREPARATIVA

Milano, 26 maggio 2015

Le proprietà dei manufatti metallici ed il loro comportamento in opera è influenzato da molteplici fattori, chi-mici, fisici e meccanici. La metallografia è la tecnica di analisi della microstruttura dei metalli che


Atti e notizie

permette di valutarne le caratteristi-che e il comportamento. Mediante l’approfondimento me-tallografico dell’evoluzione delle microstrutture - dal momento della solidificazione delle leghe fino all’ot-tenimento del manufatto finale - è possibile evidenziare i punti di forza e/o le debolezze di determinati ma-teriali, nonché comprendere come ot-tenere morfologie ottimali ed idonee a garantire le migliori prestazioni dei materiali in funzione dell’applicazione a cui saranno assoggettati in opera.Si può quindi affermare che compe-titività e qualità passino attraverso l ottimizzazione delle microstrutture. Molteplici campi di applicazione dei metalli sarebbero impensabili senza l’impiego di una corretta tecnica me-tallografica che va dall’identificazione di fasi ed evoluzioni microstrutturali, alla caratterizzazione di prodotti dopo le lavorazioni, allo studio del danneg-giamento meccanico o corrosionisti-co, e a mille altri aspetti. Il Centro di Studio Trattamenti Ter-mici e Metallografia dell’AIM, in col-laborazione con STRUERS - azienda leader nella produzione di strumen-tazioni e consumabili per le prepa-

razioni metallografiche - promuove una Giornata di studio, che vedrà la partecipazione straordinaria di Geor-ge Vander Voort, eminente esperto a livello mondiale sulla Metallografia ed analisi correlate, principale autore di oltre 370 articoli tecnici e di edizioni riconosciute a livello mondiale.Il programma della manifestazione tratterà i seguenti argomenti:

Come ridurre/ottimizzare i costi • di preparazione metallografica in laboratorio Ispezione metallografica dopo trat-• tamenti termochimiciThe Microstructure of Iron-Based • Alloys Fractography. Studying fracture • surface characteristics for better failure diagnosisStudio sulla preparativa di un cam-• pione metallografico di qualità Etching experiences•

Per ulteriori informazioni rivolgersi alla Segreteria dell’Associazione Ita-liana di Metallurgia, oppure visitare il sito www.aimnet.it.

Giornata di Studio

FAILURES NEI REFRATTARI

4 giugno 2015

Le problematiche relative alle failures (cedimenti, malfunzionamenti, cattive rese ecc.) dei prodotti refrattari utiliz-zati nel rivestimento degli impianti si-derurgici da sempre costituiscono un appassionante campo di discussione tra produttori ed utilizzatori. E’ evidente che questo tipo di pro-blematiche possono avvenire non solo per cause direttamente riferibili alla qualità dei refrattari impiegati ma a tutta una serie di possibili e complesse interazioni tra il refrat-tario e le pratiche operative dell’ac-ciaieria, le soluzioni di montaggio, la conduzione dell’impianto, il tipo di metallurgia adottata ecc.Proprio per questo motivo si sta sem-pre più affermando la filosofia del “best to fit” nella ricerca del miglior compromesso possibile tra resa, co-sto e possibili inconvenienti.

Per tentare di dare un quadro del-lo stato dell’arte in questo campo il Centro di Studio Acciaieria dell’AIM ha pensato di organizzare una mani-festazione specificatamente dedicata a questo settore.Dopo una presentazione sullo stato dell’industria refrattaria italiana ed europea presentata da Confindustria Ceramica, è previstauna illustrazione di carattere generale sulle failures dei refrattari, sulle possibili cause e sui metodi di analisi e successivamente una memoria impiantistica sui princi-pi di rivestimento del forno elettrico; nel corso dei successivi interventi verranno analizzate una per una pro-blematiche e cause nei diversi reat-tori metallurgici per la produzione di acciaio dalla fusione al colaggio.In questo modo si pensa di poter offrire a produttori ed utilizzatori un proficuo scambio di discussione e confronto per sempre meglio ottimiz-zare le scelte future.Il programma dettagliato della Gior-nata – coordinata da M. Martino -

prevede interventi su: Andamento e statistiche del setto-• re refrattario Failure analysis sui refrattari • Influenza dei parametri di proces-• so, sicurezza e manutenzione del


Vita associativa

LE PROSSIME MANIFESTAZIONI AIM

SOLIDIFICAZIONE E COLATA CONTINUA Corso itinerante – Centro A

7-8-14-15-22 maggio

Eur. Conf. HEAT TREATMENT & SURFACE ENGINEERING & 22nd IFHTSE Congress

Venezia, 20-22 maggio http:aimnet.it/ht2015.htm

LA METALLOGRAFIA ATTRAVERSO UNA CORRETTA PREPARATIVAGds – Centro TTM

Milano, 26 maggio

TUBI IN ACCIAIO AL CARBONIO: Tubi saldatiGdS – Centro LP

Cremona, 25 giugno

FAILURES NEI REFRATTARIGdS – Centro A

Milano, 4 giugno

SCUOLA DI METALLURGIA DELLE POLVERICorso– Centro MP

Imola, 11-12 giugno

TUBI IN ACCIAIO AL CARBONIO: Tubi senza saldaturaGdS – Centro LP

Dalmine, 11 giugno

XI Giornate Nazionali CORROSIONE E PROTEZIONEFerrara, 15-16-17 giugno

http:aimnet.it/gncorr2015.htm

TUBI IN ACCIAIO AL CARBONIO: Tubi saldatiGdS – Centro LP

Cremona, 25 giugno

METALLI A GRANO ULTRAFINE Gds – Centri o MFM

Parma, 3 luglio

PROGETTAZIONE STAMPI Corso Avanzato – Centro P

Bergamo, 9 – 10; 21 – 22 luglio

STEELSIM 2015 6th Int. Conf. Modelling and Simulation of Metallurgical Processes in

Steelmaking Bardolino, 23-25 settembre

http:aimnet.it/steelsim2015.htm

TRATTAMENTI TERMICI Corso modulare – Centro TTM

Milano, 29-30 settembre

Pillole per Preposti: LA MACCHINA FUSORIACorso – Centro A

Brescia, 14 ottobre

MATERIALI DI CARICA IN ACCIAIERIAGds – Centro A

Milano, 18 novembre

Per ulteriori informazioni rivolgersi alla Segreteria AIM, e-mail: [email protected],oppure visitare il sito Internet www.aimnet.it

tino EAF, criteri di progettazione e principi del rivestimento refrattario Fermate ed incidenti nel forno elet-• trico: ma è solo responsabilità del refrattario?“Failure” del lining in COV Limiti di • un lining non idoneo e analogie con EAF Arresto siviera: cause e soluzioni • Sollecitazioni sui rivestimenti di do-• lomite e soluzioni efficaci Criticità nei refrattari di placca per • il colaggio dell’acciaioPaniera: importanza della corret-• ta scelta dei refrattari e della loro messa in opera

Per ulteriori informazioni rivolgersi alla Segreteria dell’Associazione Ita-liana di Metallurgia, oppure visitare il sito www.aimnet.it

QuoTe SoCiAli AiM 2015

(ANNo SolARe)

Benemeriti (quota minima) 1.750,00 €

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 Ita-

liana. Ai soci viene riservato un prezzo

speciale per la partecipazione alle ma-

nifestazioni AIM e per l’acquisto delle

pubblicazioni edite da AIM.

Per ulteriori informazioni, iscrizioni,

rinnovi:

AIM

Associazione Italiana di Metallurgia

Piazzale R. Morandi, 2

20121 Milano

Tel.: 02 76021132/76397770,

fax: 02 76020551

e-mail: [email protected]

www.aimnet.it


Atti e notizie

CENTRO LAVORAZIONI PLASTICHE (LP ) (riunione del C.T. – 18 settembre 2014)

Manifestazioni in corso di organizzazione

- Ravanelli, Coordinatore della GdS “Fabbricazione dei tubi senza saldatura”, ha comunicato che in linea di principio la manifestazione può essere effettuata; si prevede di tenerla a Dalmine a maggio, distribuita su due giornate, in funzione dei contenuti, con visita agli stabilimenti.- Donini e Fanchini, Coordinatori della Gds “Linee di processo e finitura per prodotti piani” riferiscono della difficoltà di trovare, in questo momento particolare della siderurgia, una sede idonea a consentire anche la visita ad impianti. Si conviene sulla opportunità di attendere che le varie problematiche si risolvano, mantenendo monitorata la situazione.- Il Presidente Capoferri riferisce di quanto discusso in comitato ristretto, con i Coordinatori Donini e.Gabrielli, sugli esiti di una prima indagine conoscitiva in merito all’ ipotesi di promuovere una GdS o un Corso sull’ estrusione dell’alluminio. In particolare si è convenuto di attendere il convegno Alluminium 2000 – ICEB benchmark, programmato a maggio a Firenze, dove confluiscono tutti gli interessati al tema. Viene deciso di posticipare l’iniziativa.

Iniziative future

- Si esamina l’eventualità di promuovere un’iniziativa del tipo “Pillole per Preposti”, sperimentati con successo dal Centro Acciaieria. Si individuano le criticità, tra le quali quella di individuare gli impianti per le visite alle aziende di interesse per il CT LP, che si prevede non abbiano propensione ad ospitare. Concludendo, il Comitato rimanda alle prossime riunioni la valutazione di come proseguire.- Viene proposta una GdS dedicata al settore tiranteria, dove il livello di conoscenze è molto basso. Lo stimolo a partecipare può venire da interventi sulla nuova normativa 1090 sulle costruzioni, di difficile interpretazione e applicazione senza adeguate conoscenze metallurgiche. Si incarica Mariani di fare una prima indagine e di proporre una possibile scaletta di contenuti alla prossima riunione.- Si raccomanda di dedicare attenzione al settore raccorderia oleodinamica, ove sono da poco in uso particolari tipi di acciai con basso tenore di zolfo, che presentano criticità di lavorazione. Viene chiesto a Mariani di stilare una possibile scaletta di argomenti.

Stato dell’arte e notizie

- Si commenta che sulla rivista AIM compaiono articoli scientifici, ma poco di carattere divulgativo. La cosa è stata varie volte evidenziata, ma negli ultimi numeri la tendenza sembra accentuata. Il Presidente ne parlerà in sede di riunione dei Presidenti di Centro.

- Viene sollecitata anche una presa di posizione di AIM per spiegare ai media cosa sia una acciaieria, in modo da evitare preconcetti, che sono stati causa di accuse infondate su personale tecnico innocente.

CENTRO METALLI E TECNOLOGIE APPLICATIVE (MTA) (riunione del C.T. – 9 dicembre 2014)


- Si informa che, in merito alla programmazione della GdS. “L’impiego dell’acciaio nelle costruzioni civili”, è stato deciso di posticipare l’appuntamento al 18 marzo 2015. Infatti i tempi tra l’annuncio e lo svolgimento della Giornata risultavano troppo ristretti nel caso della prima data individuata, col rischio di non riuscire a coinvolgere tutti i possibili interessati. Soprattutto, non si sarebbe potuta istruire la procedura per l’ottenimento dei crediti formativi da parte dell’Ordine degli Ingegneri di Milano per i partecipanti. La nuova data è stata accettata da tutti i relatori, che hanno riconfermato la loro disponibilità.

Iniziative future

- Il Presidente Debernardi ha riassunto i principali temi trattati nell’incontro dei Presidenti dei Centri di Studio AIM con il Presidente AIM Mapelli, svoltasi il 27 novembre 2014. In quella sede è emerso che per lo svolgimento delle attività programmate nel 2015 è richiesta la collaborazione del CT MTA per l’organizzazione di tre eventi. Con il Presidente del Centro Fonderia, Caironi, sono già avvenuti contatti per la GdS “ Fonderia delle leghe di rame”, organizzata da AIM con Assofond. Da parte del Presidente del Centro Metalli Leggeri, Vedani, è stata richiesta collaborazione, in fase di studio, per una Giornata sul titanio. Il Presidente del Centro Rivestimenti, Bestetti, propone poi di collaborare ad una Giornata su ” Rivestimenti ed ingegneria delle superfici per l’industria alimentare”.

CENTRO RIVESTIMENTI (R) (riunione del C.T. – 15 dicembre 2014)


- Per quanto riguarda la preparazione del Corso “Rivestimenti: modulo rivestimenti per via umida”, nel quale verranno trattati anche i rivestimento sol-gel, si propone e si approvano le date del 1 e 2 luglio 2015. Si conferma la sede di Padova. I Coordinatori della manifestazione saranno Bestetti, Bolelli, Brisotto. Viene approvato il programma definitivo.- Per la GdS “Rivestimenti ed ingegneria delle superfici per l’industria alimentare”, si è accolta la disponibilità del Presidente del Centro MTA a individuare presentazioni. Tra i presenti si manifestano adesioni ad effettuare interventi.

ATTiViTà dei CoMiTATi TeCNiCi


dai Centri

Parma si ipotizza come sede. Si discute poi una bozza della locandina.

Iniziative future

- Si decide di prevedere, per l’inizio del 2016, una Giornata sulla caratterizzazione dei rivestimenti e delle superfici.

Stato dell’arte e notizie

- Affrontando il rinnovo delle cariche, vengono eletti all’unanimità come Presidente Bolelli; come Vicepresidente Brisotto e come Segretario Bestetti.- Si commenta il Convegno Nazionale sui trattamenti termici e i rivestimenti, che si terrà a Venezia il 20-22 maggio. Si fa presente che l’80% delle memorie viene da paesi stranieri.- Si ricorda che dal 6 al 10 luglio UNIBS sta organizzando una summer school sull’ALD. Per informazioni ci si può rivolgere al Vicepresidente Brisotto.

CENTRO CONTROLLO E CARATTERIZZAZIONE PRODOTTI (CCP) (riunione del C.T. – 30 ottobre 2014)


- Il Corso “Prove Meccaniche” è stato programmato per i giorni 24-26 febbraio 2015, presso il CSM di Roma che fornisce i docenti, le aule, il materiale didattico e i laboratori per le dimostrazioni, con un contributo spese da parte di AIM. Il Coordinatore Trentini conferma che il programma è ormai definito, con profondi cambiamenti rispetto ai corsi precedenti; rimane confermata la durata di tre giorni consecutivi, con l’ultima giornata riservata alle visite e alle prove in laboratorio. - Stella dà la sua disponibilità a coordinare il Corso “Analisi chimiche” . Si concorda sulla necessità di rivedere profondamente il programma classico, evitando per quanto possibile l’intervento di produttori di strumentazione che spesso forniscono presentazioni “commerciali”, con scarso interesse didattico. Si curerà di evitare la sovrapposizione con altre manifestazioni AIM che potrebbero interessare gli stessi partecipanti. Verranno inseriti nuovi argomenti che comprenderanno tecniche di analisi chimica nel campo ambientale e della sicurezza e sui rivestimenti superficiali.

Iniziative future

- L’organizzazione di una prevedibile GdS su circuiti interlaboratorio è demandata ad ALPI in collaborazione con RTM, che ha recentemente organizzato un Round Robin sul Creep. La manifestazione è prevedibile per l’autunno 2015, con sede a Cormano o presso AIM, in funzione del numero dei partecipanti.- Tenuto conto che un Corso “Failure Analysis” è stato appena proposto dal Politecnico di Milano con una consistente partecipazione, si programma di organizzare la 9° edizione del Corso AIM per l’autunno 2015. Il coordinamento rimane a cura di Fossati. Il programma già collaudato potrà presentare variazioni per quanto riguarda

i docenti. Rimane da definire la tipologia specifica per la presentazione dei casi pratici di failure.

CENTRO TRATTAMENTI TERMICI E METALLO-GRAFIA

(riunione del C.T. – 22 gennaio 2015)


- Il Presidente Petta illustra le decisioni riguardanti il Con-gresso “European Conference on Heat Treatment 2015 & 22°IFHTSE Congress” (20÷22 maggio 2015, Venezia Me-stre). Essendo pervenute circa 170 memorie sono state selezionate 90 di esse per le sessioni orali, cui si aggiun-gono due keynotes e un consistente numero di memorie poster, applicando criteri di selezione basati sui contenuti e sulla provenienza. La sessione poster avrà a disposizione un ampio spazio e verrà conferito un premio al miglior po-ster. Al momento sono ipotizzati chairman singoli, ai quali si chiederà di essere molto rigidi nel rispetto dei tempi di esposizione per consentire ai partecipanti di seguire ses-sioni parallele.- Il programma del Corso Modulare “Trattamenti Termici” viene leggermente compattato per riuscire a contenere l’attività in 4 moduli, ovvero otto giornate di lezioni più una giornata di visita ad uno stabilimento (da accorpare all’ul-timo modulo). Per individuare la sede di quest’ultimo sarà da accertare la disponibilità degli ospitanti a dare l’acces-so indistinto a tutti gli iscritti. - Si discute relativamente al Corso “Metallurgia di Base” e, valutato che l’iniziativa comporta un impegno notevole, si ipotizza di posticiparlo al 2016.- Il Coordinatore della GdS “Trattamento termico dei sinteriz-zati”, Morgano, ha fatto pervenire una bozza di programma che prevede la collaborazione del Centro Metallurgia delle Polveri e l’effettuazione a fine settembre/inizio ottobre, in sede da definire. I contenuti saranno incentrati sulle appli-cazioni automotive e si terrà una tavola rotonda conclusiva, eventualmente con inviti mirati a specialisti del settore. - Bavaro, Coordinatore della GdS “Contributo della metal-lografia alla Failure Analysis” sta contattando i relatori del-la precedente edizione e raccogliendo nuove adesioni per inserire argomenti diversi. Si concorda di puntare a collo-care la manifestazione a novembre. Durante la prossima riunione verrà discussa una prima bozza di programma.

Iniziative future

- Per il 2016, oltre a rispettare le cadenze tradizionali dei Corsi, si ipotizzano le seguenti attività: GdS su trattamenti criogenici / tecniche di tempra innovative; GdS su mate-riali per lo stampaggio (incentrato sull’automotive); GdS “Trattamenti superficiali senza materiale d’apporto” (ad esempio: tempra a induzione, laser, plasma, ecc). Si fa una prima valutazione anche di una GdS dedicata a tecniche di tempra innovative, nella quale inserire l’argomento “trat-tamenti criogenici”, sui quali sono in corso diversi lavori scientifici e che si stanno diffondendo all’interno di molte aziende.


Norme sui metalli non ferrosi pubblicate da UNI nel marzo 2015

UNI EN 1559-4:2015

Fonderia - Condizioni tecniche di fornitura - Parte 4: Requisiti addizionali per getti di leghe di alluminio

Progetti in inchiesta prEN e ISO/DIS - aprile 2015

prEN – progetti di norma europei

prEN ISO 16530-1

Petroleum and natural gas industries. Well integrity. Part 1: Life cycle governance (ISO/DIS 16530-1:2015)

prEN ISO 19901-4

Petroleum and natural gas industries. Specific requirements for offshore structures. Part 4: Geotechnical and foundation design considerations (ISO/DIS 19901-4:2015)

prEN ISO 17781

Petroleum, petrochemical and natural gas industries. Test methods for quality control of microstructure of austenitic/ferritic (duplex) stainless steel (ISO/DIS 17781:2015)

prEN 16808

Petroleum, petrochemical and natural gas industries. Safety of machineries. Manual elevators

ISO/DIS – progetti di norma internazionali

ISO/DIS 16530-1

Petroleum and natural gas industries. Well integrity. Part 1: Life cycle governance

Atti e notizie

NOTIZIE DA UNSIDER Norme pubblicate e progetti in inchiesta (aggiornamento 7 aprile 2015)

Progetti al voto FprEN e ISO/FDIS - aprile 2015

FprEN – progetti di norma europei

FprEN ISO 13702

Petroleum and natural gas industries. Control and mitigation of fires and explosions on offshore production installations. Requirements and guidelines (ISO/DIS 13702:2013)

FprEN ISO 24817

Petroleum, petrochemical and natural gas industries. Composite repairs for pipework. Qualification and design, installation, testing and inspection (ISO/FDIS 24817:2015)

FprEN ISO 16961

Petroleum, petrochemical and natural gas industries. Internal coating and lining of steel storage tanks (ISO/FDIS 16961:2015)

FprEN ISO 19901-8

Petroleum and natural gas industries. Specific requirements for offshore structures. Part 8: Marine soil investigations (ISO 19901-8:2014)

FprCEN ISO/TS 17969

Petroleum, petrochemical and natural gas industries. Guidelines on competency for personnel (ISO/DTS 17969:2015)

FprEN 1754

Magnesium and magnesium alloys. Designation system for anodes, ingots and castings. Material symbols and material numbers

ISO/FDIS – progetti di norma internazionali

ISO/FDIS 4689-2

Iron ores. Determination of sulfur content. Part 2: Combustion/titration method

ISO/FDIS 4689-3

Iron ores. Determination of sulfur content. Part 3: Combustion/infrared method

ISO/FDIS 10203

Iron ores. Determination of calcium. Flame atomic absorption spectrometric method

ISO/FDIS 10204

Iron ores. Determination of magnesium. Flame atomic absorption spectrometric method

ISO/FDIS 11536

Iron ores. Determination of loss on ignition. Gravimetric method

ISO/FDIS 15633

Iron ores. Determination of nickel. Flame atomic absorption spectrometric method

ISO/FDIS 15634

Iron ores. Determination of chromium content. Flame atomic absorption spectrometric method

ISO/FDIS 24817

Petroleum, petrochemical and natural gas industries. Composite repairs for pipework. Qualification and design, installation, testing and inspection

Normative

XI EDIZIONEFerrara - 15-17 giugno 2015

www.aimnet.it/gncorr2015.htm

Con il patrocinio di

Corrosionee Protezione

Giornate Nazionalisulla

Organizzate da

Provincia di Ferrara

ASSOCIAZIONE ITALIANA DI METALLURGIA

Coordinatore delle GiornateProf. Cecilia Monticelli

PresentazioneLe Giornate Nazionali sulla Corrosione e Protezione si terranno a Ferrara dal 15 al 17 giugno 2015. Si tratta di un evento che negli anni si è saputo aff ermare su scala nazionale come punto di incontro per discutere questioni scientifi che, tecnologiche e produttive, nell’ambito della corrosione e protezione dei materiali. Il Convegno prevede la presentazione dei risultati raggiunti da vari gruppi di studio e da numerose aziende del settore. Anche in questa undicesima edizione sono stati istituiti dei premi, destinati a giovani ricercatori che si distin-gueranno, nell’ambito della manifestazione, per l’importanza e l’attualità dei temi proposti nelle loro letture.

Aree tematiche principali

Spazio aziendeÈ previsto uno spazio per l’esposizione di apparecchiature, per la presentazione dei servizi e per la distribuzione di ma-teriale promozionale. Informazioni più dettagliate possono essere richieste alla Segreteria organizzativa del convegno ([email protected]).

Presentazione di memorieGli interessati a presentare memorie scientifi che dovranno inviare entro il 27 febbraio 2015, il titolo della memoria, i nomi degli autori della memoria e la loro affi liazione ed un sommario di circa 500 parole mediante il modulo online presente sul sito www.aimnet.it/gncorr2015.htm.

Date importanti:Invio titolo e riassunti 27 febbraio 2015Notifi ca accettazione 20 marzo 2015Apertura iscrizioni 27 marzo 2015Invio dei testi completi 30 aprile 2015

Atti

SedeLa manifestazione si terrà dal 15 al 17 giugno presso le sale Imbarcadero del Castello Estense di Ferrara in Largo Castello

Segreteria organizzativaAIM - Associazione Italiana di Metallurgia

Provato e testato molto tempo prima dell’installazione

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SMS INNSE S.p.A.

Via Milano, 4 Phone: +39 02 2124-1 E-mail: [email protected] San Donato Milanese (MI), Italy Fax: +39 02 2124-699 Internet: www.innse.com

Visit us at

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