La
Metallurgia Italiana
n. 10 Ottobre 2019 Organo ufficiale dell’Associazione Italiana di
International Journal of the Italian Association for Metallurgy
Metallurgia. Rivista fondata nel 1909
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presentazione Il Corso sulle Prove non Distruttive (PND o CND) dell’Associazione Italiana di Metallurgia data agli albori di queste tecniche di indagine su materiali e componenti. Queste tecniche e metodi di prova si sono poi sviluppati con grande rapidità ed ampiezza e sono di grande interesse per moltissimi settori, oltre che per la siderurgia. L’AIM ripropone ciclicamente un corso di informazione indirizzato ai siderurgici ma anche a tutti coloro che sono interessati alle problematiche generali dei CND: innovazioni, efficacia diagnostica, CND non tradizionali, settori di applicazione. Il Corso è volto a fornire strumenti conoscitivi per scegliere le tecniche più adatte a valutare la difettosità di materiali e componenti e stabilire in quali fasi produttive sia opportuno effettuare i controlli. Il Corso è indirizzato ai tecnici coinvolti con le problematiche qualitative, dalla progettazione alla fabbricazione ai controlli in servizio, con particolare riguardo a quelli di sicurezza. Il Corso non ha quindi lo scopo di formare gli addetti alle PND, ma è un utile strumento per fornire una panoramica esaustiva sullo stato dell’arte e sulle tecniche più diffuse ed in questo senso costituisce titolo valido per l’aggiornamento ed il mantenimento delle certificazioni degli addetti alle CND di livello III. Per i motivi suaccennati non è tuttavia necessario ai partecipanti possedere le conoscenze di base nel campo dei CND. Il corpo docenti è composto di esperti del settore in buona parte appartenenti anche all’Associazione Nazionale Prove non Distruttive (AIPND), che ringraziamo sentitamente per la collaborazione, ed è arricchito da una qualificata rappresentanza di costruttori di apparecchiature.
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La
Metallurgia Italiana
n. 10 Otobre 2019 Organo ufficiale dell’Associazione Italiana di
International Journal of the Italian Association for Metallurgy
Metallurgia. Rivista fondata nel 1909
International Journal of the Italian Association for Metallurgy www.soloswiss.it
Organo ufficiale dell’Associazione Italiana di Metallurgia. House organ of AIM Italian Association for Metallurgy.
Forni continui SOLO Swiss 322 Forni SOLO Swiss Profitherm
DEMO CENTER TEST - CONSULENZA - FORMAZIONE - R&S
Rivista fondata nel 1909 Innovative. Reliable. Precise.
n. 10 Ottobre 2019
Anno 111 - ISSN 0026-0843
Direttore responsabile/Chief editor: Mario Cusolito Direttore vicario/Deputy director: Gianangelo Camona Comitato scientifico/Editorial panel: Marco Actis Grande, Paola Bassani, Massimiliano Bestetti, Wolfgang Bleck, Franco Bonollo, Bruno Buchmayr, Irene Calliari, Enrique Mariano Castrodeza, Emanuela Cerri, Lorella Ceschini, Vladislav Deev, Andrea Di Schino, Bernd Kleimt, Carlo Mapelli, Roberto Montanari, Marco Ormellese, Mariapia Pedeferri, Massimo Pellizzari, Pedro Dolabella Portella, Barbara Previtali, Evgeny S. Prusov, Dario Ripamonti, Dieter Senk, Du Sichen, Karl-Hermann Tacke, Stefano Trasatti Segreteria di redazione/Editorial secretary: Valeria Scarano Comitato di redazione/Editorial committee: Federica Bassani, Gianangelo Camona, Mario Cusolito, Carlo Mapelli, Federico Mazzolari, Valeria Scarano Direzione e redazione/Editorial and executive office: AIM - Via F. Turati 8 - 20121 Milano tel. 02 76 02 11 32 - fax 02 76 02 05 51 met@aimnet.it - www.aimnet.it
Trattamenti termici / Heat treatment
Cross-Sectional gradients of residual stresses, microstructure and phases in a nitrided steel revealed by 20µm synchrotron x-ray diffraction S. C. Bodner, T. Ziegelwanger, M. Meindlhumer, J. Holcova, J. Todt, N. Schell, J. 6 Keckes Quenching of aluminium alloys in sodium silicate solutions of various silicate concentrations P. Krug, T. Tenostendarp, W. Stips 11 Influence of RT soaking on the stability of retained austenite in 72NiCrMo4 tool steel E. Caldesi, D. Carlevaris, A. Dauriz, S. Lindholm, A. Ometto, M. Zampiccoli, M. 19 Pellizzari Investigation on heat treatment of powder metallurgy carbon free Fe-Co-Mo alloy S. Roggero, D. Franchi - D. Magistroni, A. Rivolta 27 Manifestazioni AIM
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Attualità industriale / Industry news
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Ultra-fine-grained thermo-mechanically treated XTP® bars with high toughness for cold forming and machining edited by: M.D. Bambach, K. Helas - L. Oberli 36
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Nanostructured CVD W/WC coating prevents galling and adhesive wear of mechanisms under dry sliding conditions edited by: Y. Zhuk
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Benefits and properties of laser-hardened tool steel surfaces edited by: S. Nĕmeček - I. Černý, J. Kec - N. Ganev, J. Čapek 55 Experts' Corner / Scenari
Overview on steels, heat treatment and shoot peening in automotive field edited by: Enrico Morgano, Jacopo Tatti 59 Atti e notizie / Aim news Calendario eventi internazionali
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Normativa 69
Forni in vuoto e alto vuoto per il trattamento termico superficiale dei metalli.
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l’editoriale La Metallurgia Italiana The world of heat treatment is looking to the future, and in a complex economic context, full of unknown factors, is facing new challenges. On the one hand, the rapid development of the electric car is causing a crisis in the traditional mechanical industry (e.g. the absence of a gearbox represents a potential business risk for gear manufacturers and heat treaters). On the other hand, there is a growing interest in treatments and surface finishes of the components produced by additive manufacturing. These new products, still in their early stages, could represent an important market opportunity in the next future. The automotive world itself is not new to challenges: think of the advent of
Prof. Massimo Pellizzari
aluminum alloys, which for a long time threatened to replace steel in cars,
Dipartimento di Ingegneria Industriale, University of Trento
before the introduction of second and third-generation high-strength steels. Heat treatments in the intercritical phase field have led to the production of biphasic steels (dual phase) and to those with plasticity induced by phase transformation (TRIP), containing austenite, which have lightened the sections thanks to the improved mechanical strength and excellent ductility. We then arrived at steels with carbide-free acicular structures, containing austenite: in both nanobainitic and Q&P steels (quenching & partitioning), new knowledge on carbon partitioning led to the development of new heat treatments, opening new scenarios for steels. Similarly, surface treatments are also crossing new frontiers, thanks to the nitriding processes of stainless steels and titanium alloys, able to significantly increase hardness and wear resistance without penalizing corrosion resistance. There is also great potential for developments linked to the ever more extensive and advanced use of process modeling. All this and much more was discussed in the recent European Conference on heat treatments "Heat Treatment & Surface Engineering for Automotive", held in Bardolino (Verona) from 5 to 7 June. In the splendid setting of Lake Garda, researchers and technicians from all over Europe (and the world) came together to discuss the state of the art of heat treatments. In a context of great vitality, numerous interesting contributions were presented, some of which are included in this issue of la Metallurgia Italiana. An excellent omen for the 27th national conference to be held in Genoa in June next year, sixty years after the first edition.
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Heat treatment
CROSS-SECTIONAL GRADIENTS OF RESIDUAL STRESSES, MICROSTRUCTURE AND PHASES IN A NITRIDED STEEL REVEALED BY 20µm SYNCHROTRON X-RAY DIFFRACTION S. C. Bodner, T. Ziegelwanger, M. Meindlhumer, J. Holcova, J. Todt, N. Schell, J. Keckes
Cross-sectional gradients of residual stresses, phases, microstructure, composition and mechanical properties within the near-surface regions of thermo-chemical treated steels are decisive for the balanced mechanical properties of the final products. In this work, a correlative cross-sectional micro-analytics is introduced to assess those gradients in an exemplary nitrided steel sample to a depth of ~0.8 mm. Cross-sectional synchrotron X-ray microdiffraction with an energy of 87.1 keV and a spatial resolution of 20 µm was performed in transmission diffraction geometry. At first, a methodology to evaluate residual stress magnitudes from two-dimensional X-ray diffraction data is discussed. The data from a ~2.5 mm long sample gauge volume indicate complex residual stresses and near-surface microstructure gradients with a maximal compressive stress of ~-400 MPa, descending diffraction peak broadening and variable crystallographic texture. The results correlate well with the complementary analyses of Vickers micro hardness and sample cross-sectional morphology. In summary, the correlative cross-sectional micro-analytics documents the possibility to determine and correlate a variety of mechanical, structural, morphological and chemical sample parameters obtained using cutting-edge characterization approaches. The complex experimental data can be further used to adjust and verify numerical and technological models.
KEYWORDS: SYNCHORTRON MICRO X-RAY DIFFRACTION - CROSS-SECTIONAL CHARACTERIZATION NITRIDING – RESIDUAL STRESSES.
INTRODUCTION Thermochemical treatments like nitriding, nitrocarburizing or carburizing have been well known to increase the performance of engineering steel components. The reason for the enhancement of wear resistance, mechanical strength and fatigue behavior are gradients in the near surface regions up to a depth of a few hundreds of micrometers. In order to further improve the functional properties of the steels, however, it is necessary to optimize the thermochemical treatments technology, which go hand in hand with the understanding of the processes taking place during the surface treatment at micro and nano scale. Consequently, there is a need to design novel characterization techniques and strategies that can expound the impact of particular processing strategies on material’s functional properties. In this contribution, a correlative cross-sectional microanalytics is introduced to analyze mechanical, structural, morphological and chemical properties of an exemplary nitrided low alloy steel sample. Primarily, a focused high energy X-ray beam with a width of ~20 µm is used to scan a crosssection of the sam-ple in order to assess the complex depth 6
S. C. Bodner, T. Ziegelwanger, M. Meindlhumer, J. Holcova, J. Keckes
Department of Materials Science, Institute of Materials Physics, Montanuniversität Leoben, Austria
J. Todt
Erich Schmid Institute for Materials Science, Austrian Academy of Science, Austria
N. Schell
Institute of Materials Research, Helmholtz-Zentrum Geesthacht, Germany
gradients of phases, microstructure and residual stresses (1). Additionally optical microscopy and micro-hardness profiling were used to obtain complementary data from the sample.
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Trattamenti termici EXPERIMENTAL PROCEDURE Sample Preparation A disc-shaped, gas nitrided cylinder with a diameter of 40 mm and a thickness of 10 mm was exemplarily characterized in this work. The base material is a low alloy steel with a carbon content of less than 0.2%, typically used for case hardening in a pre-heat treaded state and provided in an industrial standard grade. Characterization Techniques To prepare the sample for the synchrotron transmission diffraction experiment, (cf. Fig. 1), a platelet of ~2.5 mm in thickness was extracted from the sample by precision cutting at a Accutom-5 system (Struers, Germany), equipped with a diamond cutting wheel.
The sample cross-section was scanned in transmission at the HEMS beamline P07 at the storage ring PETRA III in Hamburg, Germany (2) using a beam energy of 87.1 keV. Fig. 1 provides a schematic representation of the experimental setup, which will be further denoted as cross-sectional X-ray micro-diffraction (CSmicroXRD). The rectangular beam cross-section was ~(500 x 20) µm² corresponding to the beam width and height (Fig. 1). The scanning increment was set to 20 µm. A twodimensional (2D) amorphous silicon digital X-ray detector (model XRD1621 by PerkinElmer) with a pixel pitch of ~200 µm collected the diffraction signals. LaB6 standard was used to determine the sample-to-detector distance of 1329 mm.
Fig. 1 - Experimental setup of the CSmicroXRD experiment at DESY in Hamburg. The cross-section of the thermo-chemical treated sample is being scanned in transmission diffraction geometry by moving the sample in the beam along the z-axis with an increment of ~20 µm. At each scanning position, two-dimensional X-ray diffraction data are collected using a 2D detector.
The microstructure of the sample was studied by optical microscopy from the metallurgical cross sections. Therefore, the specimen was hot-embedded at 180°C at a pressure of 15 kN for 900 seconds before it was grinded, polished and etched with a Nital 3% etchant. Vicker’s hardness profiles were determined on the embedded
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sample using a Mitutoyo/Buehler Micromet 5104 testing device. A testing force of 4.90 N – corresponding to HV0.5 – was used to indent the polished, unetched cross-section. The hardness values given in Fig. 2c were calculated by averaging the result of multiple indents at redundant depth positions.
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Heat treatment METHODOLOGICAL APPROACH Debye-Scherrer rings collected using the 2D detector contains information on the lattice parameters, crystallographic texture, crystallite size and strain state of phases within the gauge volume. The 2D data were analyzed using a Python pyFAI software package (3). An integration of the 2D diffraction data over the azimuthal angle δ (Fig. 1) was used to obtain an intensity distribution as a function the Bragg’s angle θ and the obtained I(θ) data were used for the determination of the lattice parameters and the identification of the particular crystalline phases in the sample. Crystallographic texture was determined by analyzing the azimuthal intensity distributions of the Debye-Scherrer rings of the martensite phase in an angular range of 0 < δ < 90°. An information on the size of the coherently scattering domains and strains of II and III order can be obtained by analyzing the full width at half maximum (FWHM) of the diffraction peaks. Since FWHM data can be derived as a function of the azimuthal angle δ, FWHM analysis can be used to obtain direction dependent microstructural properties of the sample (4). Residual strain depth gradients along the scanning direction z (Fig. 1) can be determined by the evalua-tion of the DebyeScherrer ring’s ellipticity. In general, X-ray elastic strain along the diffraction vector specified by the angles δ and θ (7) can be expressed as follows (5)
[1]
where d0 is the unstrained lattice parameter and d(δ,θ) is the direction dependent lattice parameter. Since the initial geometry of the sample was rotationally symmetric, it can be approximated that there was no significant in-plane sample anisotropy and for the in-plane strain components ε11 (z)=ε22 (z) is valid. Furthermore, it was for simplicity supposed that the shear strain and stress components can be neglected with εij≅0 and σij≅0. Consequently, the measured in-plane X-ray elastic strain ε(δ,θ) can be expressed as
X-ray elastic strain components, while the indices i = 1,2 and 3 correspond to the axis x, y and z in Fig. 1. As extensively discussed in reference (6), by inserting the Xray elastic constants S1(hkl) and 1/2S2(hkl) in Eq. 2 and considering the relationship between X-ray elastic strains and unknown stresses σii, the distortion of the Debye-Scherrer rings as a function of the sample’s depth z can be expressed as
[3] where σ22 (z) represents to the magnitude of depth dependent in-plane residual stresses. In contrast to the cross-sectional X-ray nanodiffraction analysis (CSnanoXRD) described in (1,6–8), the out-of-plane residual stress component σ33 (z) cannot be always considered to be negligible at sample depths of several hundreds microns. Consequently, in the case of CSmicroXRD, Eq. 3 must be modified as follows
[4] RESULTS AND DISCUSSION Experimental data obtained using the correlative cross-sectional micro-analytics, including CSmicroXRD, microscopy and hardness profile characterization, are presented in Fig. 2. Due to the insufficient detector resolution, reflections of the martensite and the bainite coincide and appear as one in depthresolved phase plot of the nitrided sample in Fig. 2a. Therefore, the bainite reflections notation was used. The peaks of the thin compound layer directly at the sample’s surface are clearly visible in the colored illustration which is provided in the electronic version of the ECHT2019 conference proceedings. They indicate the presence of ε nitrides. The spatial resolution could be improved by choosing a smaller X-ray beam size and a scanning increment in the near-surface region. Fig. 2b shows the azimuthal intensity distribution of the ⍺ 110 peak. The quantitative analysis of the texture gradient within the sample indicates a presence of random texture in the nearsurface region and emerging texture at depths of ~280 µm and more.
[2] where εii represents the unknown sample depth-dependent
8
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Trattamenti termici Results from Vicker’s hardness measurements are presented in Fig. 2c. The values decrease within ~250 µm from 514 HV0.5 at a depth position of ~50 µm to the core hardness of 232 HV0.5. An effective nitriding depth of ~180 µm was determined. The optical micrograph displays zones within different microstructure. The surface is covered by a ~4 µm thin, continuously formed compound layer after which the nitrogen diffusion zone is present up to a depth of ~90 µm. The thickness of the compound layer and the diffusion zone was identified by optical microscopy with ~4 and ~90 µm, respectively. Adjacently, a region richer on carbides in a bainitic matrix follows due to the displacement and rearrangement of interstitial carbon by nitrogen in the diffusion zone. The optical micrograph shows, that the microstructure changes again at the total nitring depth of ~280 µm below the surface to the bainitic core material. Fig. 2d and Fig. 2e show the ⍺b 200 peak in in-plane and out-of-plane diffraction vector orientation, respectively. In the case of in-plane orientation, a peak shift to higher Bragg angles in the reciprocal space corresponds to a compression of the crystal planes and negative strain in the real space. In other words, the peak shift to higher Bragg angles indicate the presence of in-plane compressive stresses. This correlates with the peak shift to smaller Bragg angles in Fig. 2e, indicating lattice parameter increase and positive strains, for outof-plane diffraction vector orientation. The lattice parameter increase in Fig. 2e observed within the first ~90 µm below the surface is caused by the same compressive in-plane residual stress that are formed by the nitriding process.
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In the next step, stresses were evaluated from the ⍺b 211 peak by applying X-ray elastic constants. As indicated by the peak shifts (cf. Figs. 2d,e), the analysis of the data revealed a compressive stress profile along the sample’s depth. The stress maximum of ~-415 MPa was determined ~20 µm below the surface. More than 60% of the stress relaxation takes place, however, within the diffusion zone of the sample. The slope of the relaxation changes and the remaining stress reliefs within the following 160 µm and stays approximately constant after. In summary, correlative cross-sectional micro-analytics and CSmicroXRD represent powerful tools to resolve depth gradients of phases, texture, residual stresses, microstructure and mechanical properties across gradient materials such as nitrided, carburized, quenched and tempered steels. In comparison to standard techniques used to characterize the residual stress state in gradient steels, CSmicroXRD offers the possibility of a rapid scanning of the specimens - after the sample alignment, all experimental data from Figs. 2a,b,d-f were captured within ~250 seconds. A further advantage is the simple sample preparation, even industrial parts with a complex geometry can be characterized position resolved as one-part without cuttig by applying the conical slit system (9). Finally, the experimental data obtained using cross-sectional micro-analytics and CSmicroXRD (Fig. 2) can be used to verify numerical and technological models as well as applied to further predict the effect of modified and adapted thermochemical treatments.
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Heat treatment
Fig. 2 - Depth-resolved experimental results from the gas nitrided martensitic steel sample. (a) Phase plot with the particular reflections. (b) The quantitative texture analysis indicates random near-surface texture induced by the thermo-chemical heat treatment. (c) Hardness values decrease gradually within the diffusion zone of the layer and are slightly elevated down to a depth of ~230 µm (carbide-rich zone). (d,e) Zooms to the in-plane and out-of-plane orientations of the 200 peak show the shift to higher and lower angles, respectively, due to in-plane residual compressive stresses. (f) The depth resolved stress profile indicates a maximum at ~20 µm after which the stress relaxes in two steps with different relaxation slopes. REFERENCES [1]
Kurz SJB, Meka SR, Schell N, Ecker W, Keckes J, Mittemeijer EJ. Residual stress and microstructure depth gradients in nitrided ironbased alloys revealed by dynamical cross-sectional transmission X-ray microdiffraction. Acta Materialia [Internet]. Acta Materialia Inc.; 2015;87:100–10. Available from: http://dx.doi.org/10.1016/j.actamat.2014.12.048
[2]
Schell N, King A, Beckmann F, Ruhnau HU, Kirchhof R, Kiehn R, et al. The High Energy Materials Science Beamline (HEMS) at PETRA III. In: AIP Conference Proceedings. 2010.
[3]
Kieffer J, Karkoulis D. PyFAI, a versatile library for azimuthal regrouping. In: Journal of Physics: Conference Series. 2013.
[4]
Scherrer P. Bestimmung der Größe und der inneren Struktur von Kolloidteilchen mittels Röntgenstrahlen. Nachrichten von der Gesellschaft der Wissenschaften zu Göttingen, Mathematisch-Physikalische Klasse. 1918;
[5]
Dölle H. The influence of multiaxial stress states, stress gradients and elastic anisotropy on the evaluation of (Residual) stresses by X-rays. journal of Applied Crystallography. 1979;
[6]
Stefenelli M, Todt J, Riedl A, Ecker W, Müller T, Daniel R, et al. X-ray analysis of residual stress gradients in TiN coatings by a Laplace space approach and cross-sectional nanodiffraction: A critical comparison. Journal of Applied Crystallography. 2013;46(5):1378–85.
[7]
Keckes J, Bartosik M, Daniel R, Mitterer C, Maier G, Ecker W, et al. X-ray nanodiffraction reveals strain and microstructure evolution in nanocrystalline thin films. Scripta Materialia. 2012;
[8]
Bartosik M, Daniel R, Mitterer C, Matko I, Burghammer M, Mayrhofer PH, et al. Cross-sectional X-ray nanobeam diffraction analysis of a compositionally graded CrNx thin film. 2013;542:1–4.
[9]
Staron P, Fischer T, Eims EH, Frömbgen S, Schell N, Daneshpour S, et al. Depth-Resolved Residual Stress Analysis with Conical Slits for High-Energy X-Rays. Materials Science Forum. 2013;
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Trattamenti termici
QUENCHING OF ALUMINIUM ALLOYS IN SODIUM SILICATE SOLUTIONS OF VARIOUS SILICATE CONCENTRATIONS P. Krug, T.Tenostendarp, W. Stips
Precipitation hardening is a common process in heat treatment of automotive components, e.g. chassis parts, cylinder heads and others. Due to distortions after water quenching subsequent levelling operations or machining is necessary, which causes additional cost. Avoiding or minimizing such distortions is desirable and can be achieved by applying alternative quenching media. The use of aqueous solutions of sodium silicate enables one to adjust the quench severity in between pure water and high concentrated polymer solutions. Since critical cooling velocity is low among the 6xxx alloys in comparison to 2xxx or 7xxx alloys, a smaller cooling rate can be acceptable. Sodium silicate solutions are non-toxic and will cause no environmental hazards. This is an overview about the use of silicate media, the mechanical properties and the hardness of aluminium alloys quenched in sodium silicate solutions of various concentrations and the possibility of compensating the reduced cooling velocity by modifying the ageing parameters.
KEYWORDS: QUENCHANT – ALUMINIUM – HEAT TREATMENT – SODIUM SILICATE – DISTORTION.
INTRODUCTION Distortion of aluminium components is often a result of quenching due to thermal gradients occurring during rapid cooling. Different quenching media with various quenching intensities are available, e.g. polymer solutions, oil, hot water, pressurized air or other gases, to adjust cooling velocity. Especially the 6xxx series aluminium alloys allow moderate critical quench rates in the range of 1 to 50 K/s [1]. A quenchant with adjustable quench intensity and - of course nontoxic - would be of competitive advantage in combination with easy handling, low cost and with reduced or even no efforts of cleaning necessary. Therefore it should be water based and available in appropriate amounts. Polymer quenchants will fulfil the boundary condition of thermal gradient reduction and will lead to a more uniform heat transfer but their cooling mechanisms are not fully understood yet and a sticky film will remain on the component´s surface which has to be removed prior to ageing. In water quenching usually four phases of heat transfer mechanisms can be found [2]: Immediately after immersion of the component a vapor film is formed on the component´s surface. Due to its poor heat conductivity and good insulating properties, heat transfer coefficients are very small in
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this phase. After reaching a distinct temperature (Leidenfrost point) turbulent boiling starts, which raises the heat transfer coefficient significantly up to several thousand W/m²K. This is followed by the phase of nucleate boiling and finally to convective cooling. This phases can be controlled by the surface condition of the component [3-7] the flow velocity, the composition as well as tap water hardness [8], the viscosity and the boiling point of the quenchant [9].
Peter Krug, Thomas Tenostendarp, Waldemar Stips
Cologne Technical University of Applied Sciences - Department of Automotive Engineering - Betzdorfer Strasse 2, 50679 Cologne - Germany
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Heat treatment An aqueous sodium silicate solution may fulfil the before mentioned requirements. It is water based with a pH value between 9 to 11 which is slightly alkaline. It exhibits a viscosity higher than that of water, is non-toxic and - since it is used in several applications like silicate paintings and inorganic binders in sand casting industry – is therefore available in huge amounts at low cost. Sodium silicate solutions are colloidal, mostly clear or slightly opaque solutions, consisting of small particles homogeneously distributed in water, causing light to be scattered (Tyndall effect). These particles exist as single nanometer sized spheres (colloidal solution) or as small polymerized networks (gel). By controlling pH value and/or the content of other ionic compounds one can influence the solution in on of both directions. Usually it is available along with different viscosity grades, which are still given in the non-SI unit Baumé (°Bé). Only scarce information can be gained from literature about the application of sodium silicate solutions as a quenchant. First trials has been conducted in 1931 by Thomas Hamill, from the US. Bureau of Standards [10]. Different concentrations of sodium silicate solutions has been used in steel hardening experiments, showing the ability to modify the cooling intensity somehow in between water and oil. As an advantage he found that there was no corrosive action in contact with steel which is rather rational since plain steels are in passive state at those pH-values. As disadvantage he mentioned the instability of some diluted concentrations in terms of floc-
culent precipitation, which can be fixed by adding a certain amount of sodium hydroxide. A transcript from the Russian author Yu. G. Krotov and co-workers is available in [9]. They used sodium silicate with a concentration of 50% quenchant as an admixture together with K- and Na- spruce wood tannins in comparison to sulphite-alcohol malt concentrate with different additives for quenching the aluminium alloy D16 (AlCu4Mg1) and V95 (AlZn5.5MgCu). Only limited information is given about the effect on cooling velocity, distortion and the resulting mechanical properties but film formation is mentioned although it remains unclear from which substance these films have originated. Effect of nanoparticle size and coating with ceramic layer of various thickness was reported in [11-13]. They found, that a certain content of nanoparticles may enhance cooling but care has to be taken, since altering of the boiling behaviour will occur. Taking all the information together, one could claim that the higher viscosity may lead to a retarded boiling, but nanosizes particles in a colloidal solution may improve cooling intensity. In addition, one may expect that a solid silicate film will form on the surface after immersion due to the calcifying effect of the hot component. Film formation can shift the Leidenfrost temperature to higher values, but this depends on the resulting film thickness and the growth velocity of the silicate film in comparison to heat transfer into the surrounding liquid.
EXPERIMENTAL For quenching trials rolled sheets (thickness 25 mm) and extruded rods (31 mm in diameter) made of aluminium alloy EN AW 6082high strength were used (see also Figure 1 and 2).
Delivery state was semi-hard for the plates and T6 peak aged for the rods. The chemical composition is given in Table 1.
Tab. 1 - Nominal chemical composition
Rods were machined down to a diameter of 25 mm and with two 3 mm holes drilled parallel to the axis of the rods, one in the centre and one near the surface, for placing two thermocouples (Type K, class 1, Sensortherm GmbH) 50 mm from the upper end of the rod. A drawing is given in figure 1 showing the relevant geometrical features. The surface was wet grinded with 4000 emerald paper to an average roughness of 5.5 +/- 0.7 µm. Temperatures were sampled by a TC-08 AD/ DA data logger (Omega Engineering GmbH) with sampling
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rates of 0.1 per second. For optimization the ageing time and temperature, small cylinders of 30 mm diameter and 30 mm height were machined and heat treated, too.
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Trattamenti termici
Fig. 1 - Rods with holes for thermocouples for recording cooling rate
Fig. 2 - Plates with fillet for characterization of distortion behavior.
For distortion characterization, plates with an asymmetrical fillet were machined out of the rolled plates (see figure 2). Dimensions were measured by using an ATOS II Scanner und ATOS Professional Software (GOM GmbH). The rods were solution annealed for 40 min. at 540 °C in a vertical resistance furnace (Gero Carbolite HTRV 40/500-16) and plates in a resistance furnace (Carbolite, Typ: ELF 11/14B) for the same time at same temperature. Ageing was done in a Memmert UF55 circulating air furnace. All furnace temperatures were recorded and controlled by the same type of thermocouples as above. For quenching commercially available sodium silicate solutions (WHC GmbH) was applied with a viscosity of
37/40 °Bé and various concentrations from 0 to 35,5 wt.-% sodium silicate content were adjusted by using simple tap water. The concentrations used and data of sodium silicate solutions are given in table 2. The quenching procedure was recorded by using a Sony RX100 Mark 4 camera with up to 960 fps and a high speed camera Photron FASTCAM Mini AX200 (Model 900K-M-16GB) with frame rates up to 30.000 fps. Tensile tests were performed with a 100 kN tensile test unit (Z100 TEW) from company Zwick Roell GmbH using short proportional tensile rods with a diameter of 10 mm and hardness testing was performed at a Wolpert hardness tester DIATESTOR 2Rc using Brinell hardness HB2.5/125.
Tab. 2 - Data of sodium silicate solutions
RESULTS Boiling behaviour The boiling behavior of pure tap water and 35.5 wt.-% sodium silicate solution is completely different. As shown in figure 3 a) to c) one can distinct clearly the phases of boiling during the first seconds. In contrast, there is no evidence for any boiling action with sodium silicate solution (figure 3 d) to f)). From high speed recordings it cannot be revealed if there is a vapor blanket around the rod or not. Nevertheless, no turbulent bubble formation can be observed. There are some areas where small clusters adhere at the surface and move very slowly upward. From visual assessment, this cooling behavior is similar to the quenching in ionic liquids. La Metallurgia Italiana - n. 10 2019
At the lower right picture beginning silicate film formation can be seen, but it will not grow immediately fast over the complete surface. During further cooling more and more silicate “islands” occur and grow slowly together while some small bubbles appear from time to time at and separating off the surface (see picture on the lower left).
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Heat treatment
Fig. 3 - a) Quenching in pure water 1 second after immersion.
Fig. 3 - b) Quenching in pure water 1.75 seconds after immersion.
Fig. 3 - c) Quenching in pure water 3 seconds after immersion.
Fig. 3 - d) Quenching in silicate solution (35.5 wt.-%) 1 second after immersion.
Fig. 3 - e) Quenching in silicate solution 1.75 seconds after immersion.
Fig. 3 - f) Quenching in silicate solution 3 seconds after immersion.
If the concentration is lowered to 20 wt.-% film formation will appear instantaneously during immersion with a growth speed of approximately 0.5 m/s (!). Even lowering the conCooling rate Comparing the cooling curves of both samples reveal for the high silicate concentration a period of about 20 seconds the existence of an insulating vapor film with subsequent slow
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centration down to 2.5 wt.-% will lead to the same result but with decreasing film thicknesses.
cooling but nearly no difference in the temperatures of the two thermocouples near the surface and at the center of the rod (see figure 4a) and b)).
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Fig. 4 - a) Cooling curves recorded during quenching of rods. Temperatures measured in the middle of the rod height. b) Cooling curve for rods quenched in silicate solution (35.5 wt.-%). Please note the different time scale! The sample quenched in pure water exhibits only a small time interval of film boiling (approx. 1 to 2 seconds) with subsequent nucleate boiling. Note that the time scale is different for both graphs by a factor of 3. This effect can be better
identified, when the calculated cooling velocity is shown as a function of temperature, as it is given in figure 5 a) and b). A difference of about 45 °C/s exists in the tap water quenched rod.
Fig. 5 - a) Cooling rates calculated from cooling curves recorded during quenching of rods in pure water. b) Same as in a) but quenched in silicate solution (35.5 wt.-%). The cooling rates measured in the center and near the surface show only very slight differences. Mechanical Properties Since cooling velocity is rather slow in the 35.5 wt.-% silicate solution it is of interest how mechanical properties are influenced by the cooling rate. Tensile test specimens were heat treated and quenched in silicate solutions of various concentrations and tested after an ageing procedure comprising 4 hours at 180 °C. The results of the tensile tests are shown in La Metallurgia Italiana - n. 10 2019
figure 6. By increasing the silicate concentration to 10 wt.-% a decrease in strength of approximately 20 MPa will occur. Nevertheless, elongation seem to be unaffected by the reduced quench rate. From 10 to 30 wt.-% silicate concentration the mechanical properties remain unchanged at the same level. Only when the concentration exceeds 30 wt.-% a further decrease of 20 MPa will occur, while elongation remains fur15
Heat treatment ther constant. Since strength (and elongation) can be altered by adjusting ageing time and temperature, small cylinders of the same aluminium alloy were solution heat treated, quenched in different media and aged at different times at 180 °C. Water quenched specimens will reach a Brinell hardness
of about 100 HBW2,5/125 after 4 hours. By prolongation of ageing time to 12 hours at 160 °C for the silicate solution quenched specimens, the same hardness value can be reached. Due to the lower ageing temperature one can expect that elongation values will be improved.
Fig. 6 - Mechanical properties after complete heat treatment (solution annealed, quenched and tempered) as a function of silicate content of the quench bath. Distortion In figure 7 the results of the distortion measurement is shown for different silicate concentrations and with/without stirring during quenching. With increasing silicate content in the quench bath the quenching velocity decreases which transfers into reduced geometrical deviations.
At about 18 wt.-% the distortions are diminishing. Stirring itself has a positive influence but is not sufficient enough to avoid distortions alone.
Fig. 7 - Distortion of fillet plate quenched in sodium silicate of varying concentrations. 16
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Trattamenti termici DISCUSSION AND OUTLOOK The results confirm that sodium silicate solutions are appropriate to quench aluminium alloys. By performing quench tests with high strength type EN AW6082, upper critical cooling velocity (UCCR) is too high to avoid the precipitation of zones during the quench process completely, as it is shown
in Figure 8 a) even when quenched in pure water at ambient temperature [1]. But it is worth to mention that with silicate admixtures up to 35.5 wt.-% high temperature reaction, i.e. formation of coarse Mg2Si, could be suppressed due to its lower critical cooling rate (LCCR).
Fig. 8 - a) Achieved cooling curves drawn into the underlying timetemperature-precipitation-diagram, taken from [1].
Fig. 8 - b) Cooling behavior can be adjusted by mixing silicate solutions continuously with water.
By choosing different silicate concentrations, it is possible to adjust quenching velocity form pure water to those which can be achieved by polymer, oil or ionic liquids (refer to figure 8 b)). No evidence could have been found so far, that the nano sized colloidal particles will enhance heat transfer. The loss of strength can be counteracted by modifying time/temperature regime in the subsequent annealing process. Although silicate solutions do not cause corrosion among the 6xxx aluminium alloys, the calcified coating which remains on the surface has to be removed prior to ageing since residual moisture will cause a significant volume increase of the silicate. This removal can easily be done by using a water jet and warm water. Nevertheless one has to take into account that some amount
of silicate will be taken out of the quenching bath. Therefore monitoring of the bath quality by controlling the pH-value and the viscosity is mandatory.
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Further investigations will be made to elaborate the film formation mechanism by adding inorganic compounds to the water silicate solution. Different surface structures will be tested, too, because controlling silicate film thickness at certain locations on the component may lead to a â&#x20AC;&#x153;designableâ&#x20AC;? cooling velocity within one component. These experiments will also include higher concentrations to gain more information about the boiling suppressing mechanism of silicate solutions with high viscosities.
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Heat treatment REFERENCES [1]
Milkereit B, Wanderka N, Schick C, Kessler O. Continuous cooling precipitation diagrams of Al–Mg–Si alloys. Materials Science and Engineering: A 2012; 550:87–96. Available from: doi: 10.1016/j.msea.2012.04.033.
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Nukiyama, S. The maximum and minimum values of the heat Q transmitted from metal to boiling water under atmospheric pressure. Int. J. heat and mass transfer, 27,7; 2012, 959-970 p.
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Kozlov N. Gezielte Beeinflussung der Flüssigkeitsabschreckung in der Wärmebehandlung metallischer Bauteile mittels vorheriger Drehbearbeitung [Dissertation]. Rostock: University Rostock; Germany, 2018. Available from: doi: 10.18453/rosdok_ id00002035
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Prabhu KN, Fernandes P. Effect of surface roughness on metal/quenchant interfacial heat transfer and evolution of microstructure. Materials & Design 2007; 28(2):544–50. Available from: doi: 10.1016/j.matdes.2005.08.005.
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Mitrovic J. How to create an efficient surface for nucleate boiling? International Journal of Thermal Sciences 2006; 45(1):1–15. Available from: doi: 10.1016/j.ijthermalsci.2005.05.003.
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Kozlov N, Keßler O. Influencing on liquid quenching by surface structuring. International Journal of Thermal Sciences 2016; 101:133–42. Available from: doi: 10.1016/j.ijthermalsci.2015.10.025.
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Kim H., Truong B., Buingiorno J., Hu L.-W. Effects of Micro/Nano-Scale Surface Characteristics on the Leidenfrost Point Temperature of Water. Journal of Thermal Science and Technology 2012; 7(3):453–62. Available from: doi: 10.1299/jtst.7.453.
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Abdalrahman KHM, Sabariman, Specht E. Influence of salt mixture on the heat transfer during spray cooling of hot metals. International Journal of Heat and Mass Transfer 2014; 78:76–83. Available from: doi: 10.1016/j.ijheatmasstransfer.2014.06.070.
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Krotov YG, Khamidullin DK, Anan’in SN. Selecting the quenching medium for aluminum alloys. Metal Science and Heat Treatment 1989; 31(3):164–6. Available from: doi: 10.1007/BF00715817.
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Hamill, Thomas E., Aqueous solutions of ethylene glycol, glycerine and sodium silicate as quenching media for steels. Bureau of Standards Jounal of Research, RP357, Vol. 7, 1931, 555-571 p.
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Das SK, Putra N, Roetzel W. Pool boiling characteristics of nano-fluids. International Journal of Heat and Mass Transfer 2003; 46(5):851–62. Available from: doi: 10.1016/S0017-9310(02)00348-4.
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Soltani S, Etemad SG, Thibault J. Pool boiling heat transfer of non-Newtonian nanofluids. International Communications in Heat and Mass Transfer 2010; 37(1):29–33. Available from: doi: 10.1016/j.icheatmasstransfer.2009.08.005.
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Vassallo P, Kumar R, D’Amico S. Pool boiling heat transfer experiments in silica–water nano-fluids. International Journal of Heat and Mass Transfer 2004; 47(2):407–11. Available from: doi: 10.1016/S0017-9310(03)00361-2.
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Influence of RT soaking on the stability of retained austenite in 72NiCrMo4 tool steel E. Caldesi, D. Carlevaris, A. Dauriz, S. Lindholm, A. Ometto, M. Zampiccoli, M. Pellizzari
Tool steels are used in a wide range of manufacturing applications including the automotive field. Their heat treatment implies a rigorous scheduling to avoid scarce properties and reliability related problems. Stabilization of retained austenite (RA) due to prolonged soaking between martensite start (Ms) and finish (Mf) leads to its more difficult transformation during tempering or even after subzero treatment. Aim of this research is to analyse the influence of different room temperature soaking periods (0, 5, 120h) on the stabilization of RA and the effect of an additional deep cryogenic treatment (DCT, -196°C, 30min) before tempering in a 72NiCrMo4-2 cold work steel. A set of samples has been isochronally tempered just after quenching, so to avoid RA stabilization. Dilatometry highlighted the three classical tempering stages, namely the precipitation of transition carbides (100-200°), the decomposition of retained austenite (250-300°C) and the precipitation of cementite from transition carbides and segregated carbon (200-450°C). DCT carried out just after quenching causes the almost complete transformation of RA, so that the expansion accompanying the II tempering stage was suppressed. The RT soaking stabilizes RA such that even DCT is no more sufficient to achieve a fully martensitic structure. Furthermore, the stabilization process also affects the early the tempering stages, in particular the precipitation of transition carbides.
KEYWORDS: TOOL STEEL – CRYOGENIC TREATMENT – TEMPERING – DILATOMETRY – DIFFERENTIAL SCANNING CALORIMETRY - RETAINED AUSTENITE.
INTRODUCTION Low alloyed cold work steels are interesting materials for the production of tools (circular cutting blades, knives…). After proper machining tools are quenched leading to a steel microstructure consisting of primary martensite, retained austenite and, eventually, undissolved carbides. The fraction of austenite mostly depends on the C content, which in turn depends on the redissolution degree of carbides during austenitizing. The stability of austenite, i.e., its capability to be transformed into ferrite and carbides during tempering, is also largely determined by its chemical composition as well as by its stabilization during any room temperature aging between quenching and tempering. It’s widely known that a long time RT exposure before tempering (i.e. between martensite start Ms and finish Mf) makes the austenite decomposition more difficult [1]. Cold treatment (CT, T>-80 °C) and Deep cryogenic treatment (DCT, T << -80 °C, typically at LN temperature, -196 °C) were firstly introduced to help the martensitic transformation of austenite, in view of the larger and larger driving force obtained far below martensite start (Ms) temperature. In this sense, these treatments could be also considered as a powerful healing method to transform RT stabilized austenite. Furthermore, DCT has been claimed to be an effective way to induce a finer and more homogeneous precipitation of secondary car-
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bides during tempering resulting in higher hardness (due to the transformation of RA into martensite) and improved wear resistance [2,3]. The phase transformations occurring in low and medium alloyed steels during tempering have been extensively described in literature. The three classical stages are i) the precipitation of transition carbides from martensite, ii) the decomposition of retained austenite into ferrite and cementite and iii) the
E. Caldesi, D. Carlevaris, A. Dauriz, S. Lindholm, A. Ometto, M. Zampiccoli, M. Pellizzari University of Trento, Dpt. Industrial Engineering, Via Sommarive 9, 38123 - Italy
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Heat treatment precipitation of cementite from transition carbides and segregated carbon. Furthermore, pre-precipitation stages including C atoms redistribution by segregation and clustering at low temperature (<100°C) were reported. Clustering of C atoms is a local enrichment C atoms, which may act as preferential nucleation site for transition carbide precipitation. Segregation allows C atoms to leave the matrix to be associated to lattice defects and grain boundaries. All the above phase transformations can be conveniently studied by combining differential scanning calorimetry and dilatometry [4,5], considering that events associated to large enthalpy changes may correspond to small volume changes
Their cumulative effects are a net volume expansion and a large heat evolution. The conversion of transition carbides into cementite is accompanied by a large volume contraction and a large heat evolution, which are strongly superimposed to the signals related to the austenite decomposition. The influence of RT aging has been considered in [6,7] showing that it leads to C redistribution by clustering, thus anticipating what would happen during tempering immediately after quenching. In general, segregation at RT is much less important in steels showing high enough Ms, because big part of segregation (up to 0.2%wt C) can occur still during quenching. The stabilization of austenite was confirmed by the inability of a deep cryogenic treatment in liquid nitrogen (the immersion time is not reported in [4]) to (fully) transform this phase after 24h aging at RT [4]. The effect was much less evident after an ageing time of 5h only, confirming that the
and vice versa. Clustering does not cause any change of the unit cell volume or the lattice parameters of martensite, while it was associated to a low temperature shoulder in DSC curves. On the other hand, segregation leads to a decrease of unit cell volume and a limited enthalpy change [4]. Precipitation of transition carbides causes a large length decrease and an appreciable heat evolution. The transformation by which retained austenite (γR) decomposes into ferrite (α) and cementite (θ) (Eq.1) may be seen as the result of two separate stages including the intermediate formation of a high C austenite (Eq.2) and its final decomposition into ferrite and cementite (Eq.3) [6]:
γR→ α + θ γR→ α + γ↑↑C% γ↑↑C% → α + θ
[1] [2] [3]
aging time may play an important role with this respect. DCT (immersion in LN for 8 and 24h) has been shown to enhance the pre-precipitation step which leads to increased carbide precipitation from martensite during tempering [8]. According to this study it was proved that longer DCT soaking times lead to a higher degree of transformation of RA into martensite, as well as more clustering and segregation of carbon, which promote carbide precipitation. In this study a series of experiments has been carried out to study the influence of RT soaking on the austenite stability in a low alloyed cold work tool steel. The study also investigates if this phenomenon can be eliminated, or at least mitigated, by deep cryogenic treatment (DCT). Dilatometry and differential scanning calorimetry have been used to evaluate the effects of RT stabilization time and DCT on the tempering transformations.
MATERIALS AND EXPERIMENTAL PROCEDURES The material tested is a cold tool steel, 72NiCrMo4-2 (AISI 8670, 1.2703mod.), with the composition reported in Table 1. Tab. 1 - Chemical composition of steel (wt%)
Cylindrical samples (length l=10 mm, diameter ∅=4 mm), were obtained by mechanical machining from a rolled sheet along the longitudinal direction. All heat treatments have been carried out using a Baehr dilatometer, model 805A/D. Quenching (Q) was carried out heating up the samples to 20
900°C at 30°C/min, maintaining the samples at the maximum temperature for 15 minutes, then cooling to room temperature at 100 °C/s by means of a high pressure N2 flow. The influence of retained austenite stabilization on its decomposition during heat treatment was studied by performing La Metallurgia Italiana - n. 10 2019
Trattamenti termici an isochronal tempering at 10 °C/min up to 500 °C, after different aging times at room temperature (RT0h, RT5h and RT120h). The dilatometric strain data were mathematically averaged over time and subsequently differentiated by temperature, in order to define the different transformation peaks during tempering. For each soaking time two quenched specimens were prepared, one of which further underwent deep cryogenic treatment by direct immersion in liquid nitrogen at -196 °C for 30 min before being tempered. The complete list of treatments and samples is reported in Table 2. All samples were also isochronally tempered at 10 °C/min up to 500 °C using a differential scanning calorimeter (model
Perkin Elmer DSC7). DSC samples (discs ∅=4 mm, 50-80mg) were cut from quenched dilatometric samples by precision microcutting machine mounting a diamond blade. Lubrication and a very low feed rate were used to avoid any possible stress/strain induced transformation of austenite. Microstructural analysis was carried out by scanning electron microscopy (SEM) after standard metallographic preparation using emery papers and diamond paste. The amount of retained austenite was determined by XRD analysis using a Cu radiation. The HV1 hardness was measured according to ASTM-E92.
Tab. 2 - Heat treatments with the relative symbols
RESULTS AND DISCUSSION The dilatometric curve recorded during quenching shows the formation of martensite only, starting from 201 °C (Ms), suggesting that some auto-aging can occur still during quenching [7]. The positive slope of this curve at RT further indicates that the martensitic transformation is incomplete (Figure 1).
Fig. 1 - Dilatometric curve recorded during quenching
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X-ray diffraction analysis (Figure 2) allowed to determine about 7%RA. Accordingly, the Q sample shows a martensitic microstructure with polygonal areas of retained austenite situated between α' platelets (Figure 3) and some isolated undissolved carbides. The hardness is of 872±6 HV1.
Fig. 2 - XRD spectra of Q and Q+DCT samples
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Heat treatment DCT immediately after quenching promotes the partial transformation of RA: despite of the very low cryogenic temperature, about 4%RA is found in Q+DCT sample (Figure 2), confirming the inability of DCT to fully transform RA. According
to Villa this could be explained by the very fast cooling rate during direct immersion in liquid nitrogen [9]. The hardness of Q+DCT is of 930±5 HV1.
Fig. 3 - SEM micrographs of a) Q and b) Q+DCT samples
Tempering after quenching (Q) and DCT (Q+DCT) Dilatometric and DSC curves (Figure 4) highlight the stages occurring during isochronal tempering of steel in the as quenched state Q. According to literature [1-4-7], these can be ascribed to the following phase transformations: Peak I (150°C): Precipitation of transition carbides in martensite Peak II (280°C): Decomposition of retained austenite Peak III (330°C): Precipitation of cementite Precipitation stages (I and III) are accompanied by a volume contraction, while the decomposition of austenite (II) by a net volume expansion (Fig.1a). The superimposition between peaks II and III is well evident, so that their opposite volume change makes them distinguishable. The same transformations are accompanied by exothermic signals, the most intense being that related to the austenite decomposition (Fig.1b). Moreover, DSC evidences a low temperature shoulder (T<80°C) due to Carbon redistribution: in particular,
of dilatometric peak II and the lower intensity of peak II in DSC. The disappearance of dilatometric peak II highlights that the contraction due to the precipitation of cementite already starts at 220°C and corroborate the assumption that peak II in DSC is not uniquely associated to austenite decomposition. Clustering is also enhanced by DCT (see the shoulder in DSC), as found by one of authors for a similar steel [8]. This stronger clustering is in agreement with the stronger transition carbides precipitation shown by dilatometry (segregated C does not take part to transition carbides precipitation).
considering the C content of present steel (0.7%wt) and that a large part of segregation (0.2%C) may occur still during quenching, clustering is held to be responsible for this effect [4,6,7]. The effect of DCT can be summarized with a more intense precipitation of transition carbides and the (almost complete) transformation of austenite, as confirmed by the absence 22
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Trattamenti termici
Fig. 4 - a) Dilatometric curves and b) DSC curves for Q and Q+DCT samples
The above results are in agreement with the transformation of a large fraction of retained austenite into martensite during DCT. The higher amount of ι' will therefore lead to an enhanced precipitation of transition carbides in the first stage of tempering and enhanced cementite precipitation in the third stage. One more effect of DCT seems to be an anomalous small expansion at about 340°C, which does not correspond to any appreciable DSC signal: to the authors best knowledge this effect has not been described before and could be associated to an extra-precipitation (compared to Q) of segregated C to cementite [4]. This could be explained with the higher dislocation density introduced by DCT [10]. Alternatively, it could be also related to some further retained austenite sta-
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bilization that survived to DCT. According to Villa [11], the additional martensite transformed during the -196°C soaking produces some localized volume increase that puts the RA under a compressive stress, causing this phase to need a larger driving force to complete its transformation on tempering. As described in the introduction the RA decomposition during tempering (Eq.1) has been described as a two-step process (Eq.2-3). It is plausible that DCT may affect them, causing the dilatometric peak associated to the second transformation to be shifted towards higher temperatures. Other phenomena not completely explainable, like secondary carbides precipitation, or even non-detectable by dilatometry may be involved, and further investigation would therefore be needed.
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Heat treatment Tempering after RT aging (Q+RTx) and DCT (Q+RTx+DCT) Figures 5a-b display the influence of 5 and 120h RT soaking on the tempering behaviour of present tool steel, by
dilatometry and DSC, respectively. For purpose of comparison the curve of Q sample is also re-ported.
Fig. 5 - Influence of RT soaking on the tempering behavior of tool steel a) Dilatometric curves and b) DSC curves for Q, Q+RT5h and Q+RT120h samples A first interesting result is the markedly larger contraction due to transition carbides precipitation (peak I in Fig.5a). This effect becomes less evident after 120h than after 5h. Furthermore, the peak positions of Q+RT5h and Q+RT120h are shifted towards slightly lower temperature compared to Q, suggesting the lower activation required for precipitation. A similar, but less evident effect, can be observed also for the precipitation of cementite (peak III), which becomes more intense after RT ageing. The most plausible reason for the stronger precipitation of transition carbides is the Carbon redistribution during RT ageing, in particular the enhanced “natural” clustering oc-
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curring at RT compared to that occurring during “artificial” clustering during tempering. Clustering at RT is corroborated by previous work showing a decrease of unit cell volume of martensite during RT ageing of Fe-1.13%C for less than 50h [7]. In the same time frame no significant structural change in the austenite occurred. Aging for time longer than 50h further decreases the cell volume, due to a decrease of c lattice parameter, but less than during the early period. The lack of the first DSC shoulder in Q+RT120h is in good agreement with previous findings and confirms that clustering is almost completed after long time aging at RT. On the other hand, the limited enthalpy change associated to transition carbides pre-
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Trattamenti termici cipitation, makes the effects highlighted by dilatometry less obvious (Fig.5b). The stronger exothermic peak between 300 and 400°C for both Q+RTx samples is in good agreement with the stronger contraction due to cementite precipitation (peak III). Looking to the peak related to austenite decomposition (peak II), it does not show any appreciable influence of RT stay even if, it seems to be slightly higher in Q+RT5h than in Q. Obviously, this effect is not related to any change in RA content, since no decomposition of this phase is permitted at temperature lower than 200°C [1,12]. It could be directly related to Carbon redistribution at RT or, indirectly, to the effect of this redistribution on transition carbides precipitation. Ongoing investigations carried out by the same authors on a different tool steel (AISI D2), including a more systematic synchrotron X-ray diffraction analysis, seem to confirm the possible occurrence of C partitioning or a stronger stress relaxation of austenite after 5hRT ageing [13]. A preliminary annealing process at 150°C for 2h, leading to the precipitation of transition carbides (see green line in Fig. 5a), shows an expansion due to RA decomposition very similar to Q+RT5h. Considering the effect of DCT, it is possible to observe that it causes a stronger transition carbide precipitation and a shift toward lower temperature with respect to Q+RTx sam-
ples (Fig.6). Again, a similar conclusion can be drawn also for the precipitation of cementite, confirming that DCT plays the same influence described for sample cryotreated immediately after quenching (Figure 4). So, this effect is brought by the higher fraction of martensite after DCT, i.e., the higher amount of ε-carbides, which are available for the cementite precipitation. The latter process involves the participation of the carbon contained in transition carbides and the one contained within martensite, segregated or in solid solution [1,8]. A possible reason for the stronger cementite precipitation is also the stronger Carbon segregation produced by DCT: it has been shown that the soaking in LN can greatly improve the dislocation density responsible for segregation. This thesis is actually under investigation by other techniques (XRD diffraction), that can give direct information about the martensite structure. Independently from RT soaking time, the peak related to austenite decomposition is always present, suggesting that the effectiveness of DCT becomes clearly lower than just after quenching. The expected retained austenite stabilization produced by RT soaking is evident. It follows that the complete decomposition of RA can occur only during tempering. In other words, DCT cannot be used as a panacea to correct the effects of an unproper heat treatment practice.
Fig. 6 - Dilatometric curves recorded during isochronal tempering of quenched samples stabilized at RT for a) 5 hours and b) 120 hours (5days). La Metallurgia Italiana - n. 10 2019
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Heat treatment CONCLUSIONS The effect of RT aging on the microstructural transformations during tempering of a 72NiCrMo4-2 low alloyed tool steel has been investigated. A stronger contraction between 70 and 220°C due to transition carbides precipitation is well evident after 5 an 120h aging, suggesting some Carbon redistribution occurring at RT. In present steel, clustering completely occurred during aging for 120h. A stronger contraction due to cementite precipitation is also observed between 220 and 420°C. Retained austenite stabilizes during RT aging: a higher expansion during its decomposition occurs after 5h aging, suggesting two possible phenomena, to be further verified, i.e. a C partitioning or a stronger stress relaxation. DCT for 30min in LN produces the partial transformation of retained austenite. The higher amount of martensite leads
to stronger transition carbides precipitation, as well as to more intense cementite precipitation at higher temperature. The efficacy of DCT is higher when carried out immediately after quenching, as demonstrated by the disappearance of expansion in dilatometric curves. Indeed, after 5 and 120h RT ageing the expansion is still quite large, confirming that the complete transformation of retained austenite can occur only during final tempering. AKNOLEDGMENTS Present paper is the result of a laboratory activity carried out in the frame of the course of Advanced Metals, held by prof. Pellizzari at the University of Trento. All coauthors are students, which are acknowledged for their excellent work.
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M. Villa, et al., “In-situ investigation of martensite formation in AISI 52100 bearing steel at sub-zero Celsius temperature”, Proc. of the 2nd Mediterranean Conf. on Heat Treat. and Surface Eng., 2013, Dubrovnik, Croatia
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Kelkar R, Nash P, Zhu Y (2003) The mechanism of property enhancement in M2 tool steel by cryogenic treatment. In: 45th MWSP Conference Proceedings pp 13–19
[11]
Villa M, Pantleon K, Somers MAJ. Evolution of compressivestrains in retained austenite during sub-zero Celsius martensite formation and tempering. Acta Mater 2014; 65:383–392
[12]
Krauss G. Principles of Heat Treatment of Steel, ASM Metals Park, Ohio 1980
[13]
Pellizzari M., Menegate V., Villa M., Somers M.A:J. On the influence of deep cryogenic tratment on tempering transformations in AISI D2 steels. Submitted to the 26th IFHTSE Congress, Moscow, 17-19/09/2019
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Trattamenti termici
Investigation on heat treatment of powder metallurgy carbon free Fe-Co-Mo alloy S. Roggero, D. Franchi, D. Magistroni, A. Rivolta
Carbon free Fe-Co-Mo alloys are relatively new cutting materials characterized by improved thermal stability. In general, tools materials used for cutting applications, have to be abrasive and wear resistant and additionally they exhibit high hardness at elevated temperature in order to withstand the thermal loading of the cutting edge which leads to softening of the tool. The ternary system Fe-Co-Mo exhibits as-heat treated hardness levels similar to those of high speed steels but with higher resistance to heat softening due to high thermal stability and thermal conductivity. For this reason, this alloy can be used in the fabrication of machining tools for application where high temperatures and stresses due to temperature changes play a major role, which is the case by machining hobs and mills or machining titanium alloys, Ni – based materials or stainless steels. Carbon free alloy type Fe-Co-Mo exhibits a precipitation hardening mechanism achieved by solution annealing followed by fast quenching and subsequent aging of the Fe-Co-Mo matrix at temperatures below austenite transition to form very hard nm-sized intermetallic precipitates. In this investigation, iron based Fe-25 wt.- % Co-15wt.-% Mo alloy BÖHLER MC 90 Intermet has been used to test different heat treatment parameters, such as temperature and quenching pressure, in order to study their effects on microstructure and mechanical properties in terms of hardness.
KEYWORDS: BÖHLER MC 90 INTERMET, CARBON FREE FE-CO-MO ALLOY, VACUUM HEAT TREATMENT, PRECIPITATION HARDENING, MICROSTRUCTURE, HARDNESS. INTRODUCTION High performance machining operations, especially metal cutting applications, require reliable and durable tools characterized by good combinations between strength and toughness. Different groups of tool steels have been designed for the manufacturing of cutting tools [1]. In the annealed state, tool steels are soft and machinable to the desired shape but they exhibit high strength, high hardness and high wear resistance as a consequence of the heat treatment of the tooling parts. Nowadays, high-speed tool steels are used for most of the common types of cutting tools including drills, reamers, taps, milling cutters, end mills, hobs, saws, and broaches. Their higher mechanical properties, such as abrasive resistance and hot hardness, are due to the presence of carbide forming elements as tungsten, vanadium, cobalt and molybdenum. The more severe the service condition (higher temperature, abrasiveness, corrosiveness and mechanical loading), the higher the alloy content and consequent amount of carbides required for the tool steel. The hardening mechanism of high speed steels is connected to the precipitation of secondary – hardening carbides during the heat treatment resulting in a high performance cutting tool for service temperatures up to about 600°C. Anyway, at higher temperature, a loss of cutting performance can be
La Metallurgia Italiana - n. 10 2019
observed as a consequence of the rapid over aging of carbides. For this reason, for higher operating temperatures low cutting speeds and cooling lubricants must be used. This undesired phenomena is particularly observed in case of difficult machining titanium alloys and nickel alloys. Those tough materials are very difficult to machine, since the thermal softening combined with high abrasive wear at the cut-
S. Roggero, D. Franchi
Trattamenti Termici Ferioli & Gianotti S.p.A. – Caselette, Italy
D. Magistroni, A. Rivolta
Böhler, divisione della voestalpine High Performance Metals Italia S.p.A – Milan, Italy
27
Heat treatment ting edge can lead to the premature failure of the cutting tool during application. A promising alternative to increase cutting parameters and productivity is represented by the use of the carbon free cutting material BÖHLER MC 90 Intermet produced through a powder metallurgical process. This iron based alloy Fe- 25 wt.-%Co-15wt.-%Mo is characterized by superior thermal properties, such as thermal stability and thermal conductivity, when compared to traditional high speed steels, resulting in an improved service life for the tool. For these features, the BÖHLER MC 90 Intermet is particularly suitable in case of difficult to cut materials, such as stainless steel, nickel alloys and titanium alloys. Nowadays, the carbon free alloy Fe-Co-Mo is mainly used in the automotive industry for dry machining of internal gears. It has been demonstrated that the use of the Fe-Co-Mo alloy instead of high speed steels made it possible to increase the cutting parameters and therefore increase the productivity [2]. For this reasons, Fe-Co-Mo alloy is used in the manufacturing process of several hobs and end mills geometries.
In the ternary system Fe-Co-Mo, the hardening mechanism is due to precipitation of very fine intermetallic phases from the matrix, the so called µ-phase (Fe, Co)7Mo6.Thanks to these fine precipitates, hardness values up to 68 HRc and high thermal stability are obtained. The hardening behaviour of FeCo-Mo alloys is similar to Al – alloys. After solution annealing and quenching the Fe-Co-Mo martensitic matrix is quite soft (~ 45 HRc) but its hardness increases with ageing process at elevated temperatures below austenite transition due to the formation of the intermetallic µ-phase precipitates [3,4]. The low hardness after solution annealing allows machining of tools after quenching and the final aging step takes place without dimensional changes or distortion of components as it doesn’t involve any phase transformation. Therefore it can be said that geometrical precision after hardening is higher for the ternary system Fe-Co-Mo when compared to the case of common high speed steels. The aim of this study is to investigate the effects of different heat treatment parameters on microstructure and hardness of the PM-alloy BÖHLER MC 90 Intermet.
EXPERIMENTAL All experimental heat treatments have been carried out on a
Fe-Co-Mo alloy with the nominal composition shown in table 1.
Tab. 1 - Nominal composition of Böhler MC 90 Intermet in supplied condition
From the bar of 82 mm diameter furnished by Böhler Edelsthal, voestalpine division of High Performance Metals Italia, different disc – shaped heat treatments samples of about 22 mm thickness have been cut. For control purpose, hardness test has been performed in the as – delivered state before the execution of the heat treatment.
The hardness test has been carried out using a Vickers testing machine with an applied load of 20 kg. The measured hardness in the as-delivered state was 380 HV20. A light microscope image of the microstructure in the as – delivered state is shown in Fig. 1
Fig. 1 - MC 90 Intermet microstructure in longitudinal direction prior to the heat treatment (500x) 28
La Metallurgia Italiana - n. 10 2019
Trattamenti termici As confirmed by the performed EDS analysis, the material in the as – delivery conditions contains a certain amount of coarse µ phase due to the manufacturing process of the material itself. In any case, the microstructural homogeneity is guaranteed. The BÖHLER MC 90 Intermet specimens have been heat treated in horizontal vacuum furnaces equipped with high pressure Nitrogen gas quenching system and convection technology. Every heat treatment cycle has been carried out in presence of a load sensor inserted in an appropriate heat sink in order to
monitor the temperature during the entire process. To study the effect of temperature selection during the solution annealing, the austenitization has been carried out at different temperatures in vacuum condition, varying from 1170°C to 1195°C for 30 minutes, followed by fast quenching with Nitrogen gas at 5.5 bar pressure, while the aging step has been carried out at the fixed temperature of 590°C for 3 hours in protective atmosphere condition with Nitrogen gas and convection mode.
Fig. 2 - Heat treatment process of BÖHLER MC 90 Intermet [5] For some test pieces, the heat treatment has been carried out at fixed optimal temperatures, i.e solution heat treatment at 1190°C for 30 minutes and aging at 590°C for 3 hours, but varying the Nitrogen gas pressure from 4 to 7 bar during the quenching phase. In order to study the influence of the stability of the quenched Fe-Co-martensite, aging treatments have been carried out immediately after solution heat treatment and a month later. HARDNESS RESULTS In table 2 the results of hardness measurements after solution heat treatment are shown. Hardness tests have been
Any differences have been found from hardness and microstructure point of view between the samples aged in continuous and aged after a month. Hardness tests have been performed on each specimen on the surface and in the longitudinal section but, as expected, no differences have been noted. Metallographic preparation and etching with Picric acid have been performed in all cases.
performed employing both a Vickers- Armstrongs hardness tester and a Future-Tech Rokwell hardness testing machine.
Tab. 2 - Measured hardness values of the different solution heat treated samples La Metallurgia Italiana - n. 10 2019
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Heat treatment Typical hardness values achievable after solubilisation are in the range within the range of 44 – 47 HRC. All hardness values are compliant with those expected saving the case of solution annealing performed at 1170°C.
In table 3 the results of hardness measurement after solution annealing in different condition and aging at 590°C are reported.
Tab. 3 - Measured hardness values of completed heat treated samples in different conditions
In the case of series n.1 of specimens, the Vickers hardness impressions have revealed the presence of a brittle microstructure, which is due to the high hardness of 68 HRc, as shown in
Fig. 3. This fact, can be related to the very high pressure used during quenching after solution heat treatment.
Fig. 3 - Vickers hardness indent in specimen n. 1a, 950 HV30 (100x), micro crack visible on the right end Hardness testes performed on specimens aged at 590°C a month later of solubilisation with different parameters, have 30
not revealed any substantial differences in hardness values between samples aged in continuous. La Metallurgia Italiana - n. 10 2019
Trattamenti termici MICROSTRUCTURE INVESTIGATION In table 2 the results of hardness measurements after solution heat treatment are shown. Hardness tests have been
performed employing both a Vickers- Armstrongs hardness tester and a Future-Tech Rokwell hardness testing machine.
Fig. 4 - MC90 Intermet microstructures after solution annealing with different parameters (500x).
Light microscope images of the microstructures of the specimens after austenitization in various conditions and aging at
590°C are shown in Fig. 5.
Fig. 5 - MC90 Intermet microstructures after solution annealing with different parameters and aging at 590°C.
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Heat treatment After solution heat treatment and aging, all samples are characterized by fine and regular microstructure but differences in the amount of precipitates can be observed. All samples contain primary precipitates of µ phase of 1 – 4 µm diameters and secondary precipitates of µ phase, which are mostly responsible for the final hardness, in the nm range, however these cannot be resolved by light optical microscopy. Differences from microstructural point of view, and, as a consequence, from the hardness one, are related to the austenitization temperature. At higher temperatures, the µm-sized µ-phase particles are nearly completely dissolved in the matrix and, as observed for high speed steel, the grains tend to coarse with subsequent loss of mechanical properties, e.g. ductility. At lower solution annealing temperature of 1170°C, less µ- particles are dissolved and therefore less molybdenum is SCANNING ELECTRON MICROSCOPY Scanning electron microscope combined with Energy Dispersive Spectrometry (SEM – EDS) has been used to evaluate the chemical composition both of the matrix and precipitates from a quality point of view. An example of SEM image acquired using BSD detector is shown in Fig. 6.
available for the subsequent precipitation of intermetallic nm-sized µ-phase particles to guarantee an adequate increase in hardness. The microstructure after solution annealing at 1170°C and gas quenching with 5.5 bar nitrogen, shows the presence of µ-phase precipitates which probably are the same observed in the as – delivery condition. After complete heat treatment, i.e. at the end of aging step at 590°C, no substantial differences in the microstructure have been observed expect for a microstructural refinement. By that it could be shown that the solution annealing must be carried out in that way to guarantee a small remaining amount of intermetallic µm-sized µ-phase particles in the matrix in order to avoid the coarsening of grains. However, too low solution annealing temperatures lead to lower hardness after aging.
The SEM-image in back-scatter image mode reveals Fe-Co matric and the µm-sized µ-phase particles. Using the scanning electron microscopy, also the measurement of precipitates’ size has been done.
Fig. 6 - Sample n. 3A, solubilized at 1190°C and aged at 590°C No substantial differences in precipitates dimension between the various heat treated specimens have been found. All aged samples are characterized by primary intermetallic phase with size in the µm range and secondary intermetallic phase with diameters in the nm range. The intermetallic µ-phase particles are mainly located at grain boundaries. 32
Using the SEM – EDS a martensitic matrix, which appears grey, consisting of Fe – Co with some Mo, has been observed in all fully heat treated specimens while both the primary and secondary precipitates consist of µ - phase with chemical composition (Fe, Co)7Mo6, as confirmed in previous investigations [6].
La Metallurgia Italiana - n. 10 2019
Trattamenti termici CONCLUSION The studies of the role of solution annealing temperature as well as quenching pressure selection have been performed. At lower temperature, i.e. 1170°C, the measured hardness is out of the expected value being the amount of dissolved molybdenum after solution annealing not sufficient to guarantee the precipitation of nm-sized µ-phase and, as a consequence, to achieve high hardness values. At higher temperature, the amount of dissolved molybdenum increases and also the amount of secondary precipitates formed during the aging step. However, the solution temperature of 1195°C is not high enough to appreciate substantial differences in the final hardness values when compared to those obtained at 1190°C, i.e. at optimal solution temperature. As expected, varying the quenching pressure, differences on final hardness values have been found. Using 4 bar nitrogen
gas pressure after austenitization, a maximum of about 66 HRc have been obtained on sample n. 2a while hardness level of 68 HRc has been measured on sample n.1a quenched whit 7 bar nitrogen gas pressure. However, a minimum of 5 bar nitrogen gas pressure must be used to obtain suitable hardness level. The solution treatment affects the final content of intermetallic µ-phase particles in the final microstructure as well as the grain growth and it should be done at a temperature at which a sufficient amount of molybdenum is dissolved in order to allow a proper subsequent secondary hardening. Moreover, also the quenching step must be carefully performed, selecting the proper gas pressure in order to avoid undesirable effects, , e.g. brittleness effects in case of too high gas pressure or lower hardness at lower gas quenching pressure.
REFERENCES [1]
Roberts G, Krauss G, Kennedy R. Tool steels. ASM International.
[2]
Klocke F, Döbbeler B, Seimann M. Dry broaching using carbon free steel as tool material. 7th HPC 2016-CIRP Conference on high performance cutting.
[3]
Danninger H, Harold Ch, Gierl Ch, Ponemayr H, Daxelmueller M, Simancik F, Izdinsky K. Powder metallurgy of carbon-free precipitation hardened high speed steels. Acta Physica Polinca A No. 5; Vol. 177 (2010).
[4]
Danninger H, Rouzbahani F, Harold Ch, Ponemayr H, Daxelmüller M, Simančík F, Iždinský K. Powder metallurgy carbon free tool steels Fe-Co-Mo with varying Co and Mo contents. Powder metallurgy progress, Vol. 13 (2013), No. 2.
[5]
MC90 DE - 05.2014 - 1.000CD.
[6]
Danninger H, Rouzbahani F, Harold Ch, Ponemayr H, Daxelmüller M, Simančík F, Iždinský K. Heat treatment and properties of precipitation hardened carbon-free PM tool steels. Powder metallurgy progress, Vol. 5 (2005), No. 2.
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Heat
Le manifestazioni AIM treatment AIM meetings and events METALLURGIA PER NON METALLURGISTI Corso Milano, 15-16-22-23-29-30 ottobre METALLURGIA SICURA Corso itinerante 30 ottobre, 6-13 novembre SALDATURA AD ARCO DELLE LEGHE LEGGERE: STATO ATTUALE E SVILUPPI FUTURI (ML) Giornata di Studio Milano, 31 ottobre IGIENE DELLE LEGHE DI ALLUMINIO Corso Carmagnola c/o Teksid Aluminium, 5-6 novembre DALLE DUE ALLE QUATTRO QUOTE. GLI ASPETTI METALLURGICI Giornata di Studio Pontedera c/o Piaggio, 7 novembre PROVE NON DISTRUTTIVE Corso Milano, 20-21 novembre MeMo - METALS FOR ROAD MOBILITY International Meeting Bergamo, 21-22 novembre LE EMISSIONI IN ATMOSFERA NEL SETTORE METALLURGICO (AS) Giornata di Studio Verona, 5 dicembre
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metallurgia sicura
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Safety first: come trasformare una priorità in un valore aziendale 30 ottobre 2019, Aosta c/o Cogne Acciai Speciali 6 novembre 2019, Vicenza c/o AFV Acciaierie Beltrame 13 novembre 2019, Brescia / Roncadelle c/o Almag
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CENTRO DI STUDIO AMBIENTE E SICUREZZA
Corso valido come aggiornamento quinquennale perASPP/RSPP/Dirigenti (art. 37 comma 7 del Decreto Legislativo 9 aprile 2008 n. 81 e s.m.i.)
presentazione L’obiettivo del Corso è quello di presentare argomenti relativi alla sicurezza e salute del lavoro proponendo alcune soluzioni concrete per promuovere la sicurezza come parte integrante del business e vero e proprio valore per l’azienda. Il Corso si svilupperà in modo modulare partendo dagli aspetti legati alla Valutazione dei Rischi ed al ruolo dei Preposti, ai metodi più efficaci per sviluppare un sistema di gestione partecipato che metta l’uomo al centro della prevenzione, approfondendo poi i temi legati ad impianti e macchine sicure. Il Corso è rivolto in particolare ai Dirigenti e Delegati per la Sicurezza, Responsabili del Servizio di Prevenzione e Protezione (RSPP), Addetti al Servizio di Prevenzione e Protezione (ASPP), ed ai Rappresentanti dei Lavoratori per la Sicurezza (RLS), agli EHS Manager. Inoltre, il Corso è utile per i Responsabili di Reparto o di Squadra, di produzione e di manutenzione e per i tutti i Preposti che non hanno una competenza specifica nel campo della sicurezza ma che si trovano a dover gestire persone, impianti, tecnologie, nel rispetto delle normative vigenti e delle direttive aziendali. Questo Corso si articola secondo la consolidata formula che prevede l’integrazione tra la presentazione di alcune tematiche da un punto di vista tecnico e normativo, la visita agli impianti produttivi, il confronto con i tecnici che li gestiscono attraverso casi concreti di applicazione. Questa impostazione consente ai partecipanti di aggiornare le proprie conoscenze teoriche e nello stesso tempo di osservare soluzioni implementate da alcune aziende operanti nel settore metallurgico. Il corso è valido come aggiornamento quinquennale per ASPP/RSPP e Dirigenti per tutti i settori ATECO (art. 37 comma 7 del D.Lgs. 9 aprile 2008 n.81 e s.m.i.). La seconda giornata di Corso (6 novembre) è valida come aggiornamento per formatori (d.m. 6/6/2013). Il programma completo è disponibile sul sito: www.aimnet.it
Coordinatori del corso: Ettore Bordon, Valerio Forti, Matteo Gelmi
Via F. Turati, 8 · 20121 Milano Partita IVA: 00825780158 Tel. 02-76021132 / 02-76397770 e-mail: aim@aimnet.it www.aimnet.it
Industry news Ultra-fine-grained thermo-mechanically treated XTP® bars with high toughness for cold forming and machining edited by: M.D. Bambach, K. Helas, L. Oberli This contribution presents a new intensive forming method for producing ultra-fine-grained steels for cold forming and machining applications. As a thermomechanical process, the XTP®-Technology offers a possibility to enhance the mechanical properties of conventional grades beyond their limits. This is demonstrated on the hand of two commercially available steels: 7MnB8 and 13MnSiCr7. In the case of 7MnB8, the room temperature toughness is increased by 150 J and the transition temperature is reduced by 100 K compared to the hot rolled state. XTP®-processing of 13MnSiCr7 allows for the adjustment of UTS values of up to 1250 MPa and room temperature toughness values of 190 J. Further technology development aims at the processing of high strength materials such as Titanium alloys. Based on these examples, the advantages of the XTP®-Technology are discussed with respect to the resulting microstructures and following forming operations. The wide variety of adjustable microstructures enables the variation of both strength and toughness within a given range depending on each unique application. Thus, tailoring of the mechanical properties on customers’ demand becomes possible.
KEYWORDS: XTP®-TECHNOLOGY – INTENSIVE FORMING – THERMO-MECHANICAL TREATMENT – TOUGHNESS – BAINITE – COLD FORMING – TITANIUM
M. D. Bambach, K. Helas GMT mbH, Berlin, Germany
L. Oberli Steeltec AG, Emmenbrücke, Switzerland
INTRODUCTION Nowadays, the trend to constantly increase the performance of different technologies requires the development of new materials with enhanced material properties. Therefore, it is necessary to further develop existing alloying and processing concepts. In this contribution, the XTP® (Extreme Technology Performance)Technology is presented as a possibility to yield materials with enhanced mechanical behavior. This technology combines a one pass intensive rolling and thermomechanical treatment [1, 2]. All process temperatures can be varied to fit the specific requirements of the alloy and its application. Thus, the mechanical properties of common steels can be improved significantly. There is a significant difference in the mechanical properties of steel after conventional rolling and the XTP®-Technology. Among others, the latter leads to a much finer grain size resulting in an extraordinary combination of strength and toughness [3]. The key 36
challenge, however, is the determination of process parameters for the different steel grades that will lead to their unique material properties. Within extensive trials, the process parameters are varied one by one and their influence is recorded. This results in different possibilities for the XTP® processing of steel which can be applied in a flexible manner according to the requirements of the application. This paper is organized as follows: in the next section an overview of the XTP®-Technology is given. As next, the technology is applied to two commercially available steel grades: 7MnB8 and 13MnSiCr7. The adjustment of the process parameters, the resulting microstructure and material properties are presented and discussed in detail. Finally, the XTP® processing of high strength alloys such as Grade 6 Titanium alloy is introduced. After drawing conclusions based on the performed analysis, potential for future work is identified in the last section. La Metallurgia Italiana - n. 10 2019
AttualitàTrattamenti industriale termici THE XTP®-TECHNOLOGY: AN OVERVIEW The XTP®-Technology combines the two basic components to reach new categories of mechanical properties within one facility: (i) intensive forming and (ii) thermomechanical treatment (see Fig. 1). The temperature control is performed in a flexible manner thus enabling a wide range of heat treatment cycles. For example, as an initial step common steel grades can be heated, held at this temperature or subsequently cooled down at the Temperature Control Area (TCA) inlet to adjust the desired rolling temperature which can be as low as 700 °C. The intensive forming consists of one rolling pass with numerous single incremental forming stages. • • •
austenitizing temperature TA; degree of deformation φ; rolling temperature T R;
Degrees of deformation of 0.7 as well as diameter tolerances up to h11 can be achieved thus enhancing the possibilities of both conventional warm and hot forming processes. Moreover, the temperature profile after the rolling step can be varied thus providing an option for tempering. An enhanced flexibility of the facility is obtained by the possibility to adjust the individual process parameters at each process step. Figure 2 gives an impression on the different options enabled through the XTP®-Technology. Important process parameters that can be adjusted individually are the: • •
cooling intensity after rolling; tempering process.
Fig. 1 - A schematic overview of the XTP®-Technology
Fig. 2 - Temperature variation in combination with intensive forming during XTP® processing (schematic). Black solid line: austenitization followed by rolling in the austenite range. Blue solid line: austenitization followed by thermomechanical rolling in the two-phase region. Green solid line: heating and rolling in the two-phase region. Dotted lines: various possibilities for cooling and tempering. La Metallurgia Italiana - n. 10 2019
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Industry news XTP® PROCESSING OF STEEL Up to now, about 20 steel grades have been processed successfully with the XTP ®-Technology. In the next sections, the XTP ® processing of two steel grades, 7MnB8 and
13MnSiCr7 (chemical composition given in Tab. 1), will be presented in detail.
Tab. 1 - Chemical composition of the selected steels in mass-%
7MnB8 The cold heading steel grade 7MnB8 offers a wide range of mechanical properties. Due to its low Carbon content it is characterized by good weldability among others. Using the XTP ®-Technology, the optimized material properties of this steel grade make it possible for applications such as spring bows or ball studs. This steel grade contains a complex combination of microalloying elements: Titanium (Ti), Vanadium (V) and Boron (B). It is important to consider the temperature range at which they act in order to influence the microstructure. Especially for Ti and V, it is important to choose a propriate austenitizing temperature TA (see Fig. 3). With increasing TA, a slight strength increase of the material is observed along with a significant drop in the impact energy associated with the occurrence of grain coarsening. Coarse TiN particles do not
represent obstacles to the grain growth. Moreover, Vanadium precipitates are not effective for hindering grain growth at high temperatures. Figure 5 (a) shows a coarse austenitic microstructure with evolving granular bainite. In contrast, the high values of the low temperature toughness at low TA result from the fine-grained duplex microstructure of the recrystallized fine-grained polygonal ferrite and granular bainite (see Fig. 5 (b)). Continuing the parameter investigation, the heating temperature was kept low in the following trials. In order to improve the strength, the cooling rate was adjusted (see Fig. 4). It can be seen that the strength increases with raising the cooling rate as expected. The impact energy obtains its highest value after air cooling (CR 1).
Fig. 3 - Mechanical properties as a function of the austenitizing temperature (Tᴀ1 < Tᴀ2 < Tᴀ3; d0 = 30 mm; TR / φ / CR = constant)
Fig. 4 - Influence of the cooling rate on the mechanical properties (CR1 < CRn < CR6; d0 = 30 mm, TA / TR / φ = constant)
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AttualitàTrattamenti industriale termici (a)
(b)
Fig. 5 - The microstructure of 7MnB8 (TR = constant); magnification x 500; transverse section. (a) high TA; φ = 0,4, (b) moderate TA; φ = 0,6. The observed mechanical behavior can be explained by the microstructure (see Fig. 6). When air cooled, the microstructure consists of upper bainite and polygonal ferrite. The amount of ferrite decreases with increasing the cooling
intensity, resulting in a microstructure mostly consisting of upper bainite. Higher cooling rates yield acicular phases and transition towards lower bainite occurs.
(a)
(b)
(c)
(d)
Fig. 5 - The microstructure of 7MnB8 (TA / TR / φ = constant); magnification x 1000; transverse section; 1/2d. (a) CR1 < (b) CR2 < (c) CR3 < (d) CR4.
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Industry news The conventionally rolled 7MnB8 is usually used for cold heading applications. In addition, the possibility for further forming after XTP® processing was tested (see Fig. 7). XTP ® treated bars with a tensile strength of 1100 MPa – a strength level similar to the one of a 42CrMo4 quenched
and tempered steel grade – and an elongation at fracture of 13 % were used for cold bending with a bending radius of 10 mm. Even at bending of about 150 °, no cracks could be detected during microstructural investigation. [3]
Fig. 7 - Result of a bending test with 7MnB8 (d = 18 mm; Rm = 1100 MPa; Rp0.2 = 990 MPa; A5 = 13 %) [3] 13MnSiCr7 The low alloyed Cr-Mn-steel grade 13MnSiCr7 has a predominantly bainitic microstructure. Compared to 7MnB8, it is distinguished by higher strength levels. Therefore, it represents an attractive alternative to the presented low carbon steel grade for e. g. U-bolts. Again, the aim of the investigations with this material was to find process parameters leading to the best combination of high strength and high toughness. A slight variation in the austenitization temperature TA did not significantly influence the strength whereas the impact toughness was improved significantly (see Fig. 8). The rolling temperature TR was kept constant for the different trials and the cooling intensity was varied.
It was found out that the strength increases with increasing cooling intensity (Fig. 9). However, no linear relationship could be observed between the toughness and the degree of cooling. The highest toughness was obtained for medium cooling rates. It is assumed, that the amount of ferrite goes up with lower cooling rates and the amount of martensite increases with a more intensive cooling. Nevertheless, the low impact toughness values at low cooling rates cannot only be explained by the larger amount of ferrite in the microstructure. The tendency of this steel grade to temper embrittlement at around 350 °C could be another reason for the loss of toughness.
Fig. 8 - Mechanical properties as a function of the austenitizing temperature (Tᴀ1 < Tᴀ2 < Tᴀ3; d0 = 40 mm, TR / φ / CR = constant)
Fig. 9 - Influence of the cooling rate on the mechanical properties (CR1 < CRn < CR4; d0 = 35 mm, TA / TR / φ = constant)
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AttualitàTrattamenti industriale termici As Fig. 10 shows, the XTP® process offers the possibility to adjust a wide variety of mechanical properties for one and the same material. Compared to the standard conventionally rolled state, the room temperature impact toughness can be increased by around 140 J at the same tensile strength,
whereas the -40 °C toughness is improved by more than 50 J. Using the XTP ®-Technology, the material properties and microstructure of 13MnSiCr7 can be tailored in such a way to fit each unique application in a best way.
Fig. 10 - XTP® processing of 13MnSiCr7 with various process parameter combinations: a whole spectrum of achievable material properties compared with the standard condition (conventionally rolled and drawn) and the corresponding microstructures from the CCT-diagram (I: ferrite + pearlite (+ bainite), II: ferrite + bainite, III: bainite (+ martensite), IV: martensite) Summary on the XTP® processing of steel So far, the strategy to process new steel grades with the XTP®-Technology has been explained. Two commercially available steel grades were studied extensively in order to determine the optimum process parameters allowing for a superior combination of strength and ductility. With the XTP ®-Technology, a wide range of mechanical properties can be adjusted to one and the same steel grade thus
tailoring its material behavior to a specific application. The following figure summarizes the different levels of properties obtained by varying the process parameters for the two steel grades presented in this paper as well as for other steel grades during XTP ® processing compared to conventionally processed steels.
Fig. 11 - Expanding the mechanical properties of different steel grades via XTP® treatment. Modified from [4].
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Industry news XTP® PROCESSING OF HIGH STRENGTH MATERIALS Titanium and Nickel-base alloys are the next challenge for the XTP®-Technology and first trials have already been done. The processing of a Grade 6 alpha-Titanium alloy (Ti5Al-2.5Sn) was investigated. This alloy is a non-heat treatable Titanium alloy containing predominantly alpha phase [6, 7] and is usually used for cryogenic applications. The rod material was intensively rolled in two passes as follows: i. rolling from 24.8 to 19.5 mm followed by recrystallization annealing at 930 °C for 95 min and air cooling to room temperature; ii. rolling from 19.5 to 13.6 mm followed by recrystallization annealing at 930 °C for 60 min and air cooling to room temperature.
In both cases forming and annealing was performed below the beta-transus temperature (1030 °C). The resulting microstructure shows an extremely fine equiaxed and some plate-like alpha phase in the case (see Fig. 12 (a) and (b)) with an increasing fraction of acicular alpha grains in the core (see Fig. 12 (c) and (d)). From the decrease in the grain size both higher ductility and formability as well as higher strength and increased fatigue resistance are expected [5, 6].
(a)
(b)
(c)
(d)
Fig. 12 - The observed microstructure at different magnification (x 10 / x 100); longitudinal section; Kroll reagent (HF + HNO3 + H2O). (a) and (b) – case, (c) and (d) – core.
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AttualitàTrattamenti industriale termici Surface hardness of 32 HRC was measured at the case of the sample cut in transverse direction. This value is slightly above the average hardness according to literature values for annealed rods of the same grade [7]. The strength and the low temperature toughness of the material after XTP ® processing still need to be determined.
Because of the small reduction ratios used for Titanium and Nickel-base alloys, a new processing scheme is currently in development to process Titanium including a reverse intensive rolling (see Fig. 13). Thus, the material can be reheated after each rolling pass and higher reduction ratios can be made possible.
1 ... Rolling mill 2 ... Drive block 3 ... Drive shaft
Fig. 13 - The XTP®-Technology for Titanium and Nickel-base alloys
CONCLUSIONS AND OUTLOOK The XTP ®-Technology offers a possibility for a product-specific tailoring of both the microstructure and the mechanical properties of conventional steel grades in order to best fit a certain application. Through this technology, an optimum combination of increased strength and toughness is obtained. This was demonstrated on two commercially available steel grades, 7MnB8 and 13MnSiCr7. Moreover, the XTP®Technology seems promising for processing high strength alloys as well. First experiments on the processing of a Titanium alloy were made successfully and an extremely fine microstructure was detected.
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Future work will concentrate on using the XTP®-Technology to improve the mechanical behavior of both existing as well as new steel grades in order to enhance the field of their application. Furthermore, the XTP®-Technology will be further developed for the successful processing of high strength materials such as Titanium and Nickel-base alloys.
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Industry news REFERENCES [1]
Borowikow, A.: Integration of Cross Rolling Step into the Heat Treatment Processing of Steel Bars. In: European Conference on Heat Treatment 2011, Wels, Austria (2011), pp. 131-135.
[2]
Borowikow, A., Mašek, B., Jirková, H., Jeníček, Š., Bublíková, P.: New Technological Concepts of Long Products based on HDQT - High Deformation Quenching and Tempering System. In: ROLLING 2013. Venice: Associazione Italiana di Metallurgia (2013), pp. 1-7.
[3]
Lembke, M. I., Oberli, L., Olschewski, G., Dotti, R.: Surpassing steel performance by creating a very fine grained structure. In: METALLURGIA ITALIANA 6 (2018), pp. 31-36.
[4]
Keul, C. Mosecker, L. International SCT Conference; Steels for Cars and Trucks; June 5.-9. 2011, Salzburg Austria
[5]
Publisher: ASM International: ASM Handbook. Volume 2: Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, (1990), pp. 1782-1880.
[6]
Leyens, C., Peters, M. (eds.): Titanium and titanium alloys: fundamentals and applications. John Wiley & Sons, (2003), p. 22.
[7]
Welsch, G., Boyer, R., Collings, E. W. (eds.): Materials properties handbook: titanium alloys. ASM international, (1993), p. 300.
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AttualitĂ Trattamenti industriale termici Nanostructured CVD W/WC Coating Prevents Galling and Adhesive Wear of Mechanisms under Dry Sliding Conditions edited by: Y. Zhuk Heavy-duty engine and drivetrain parts operating under high loads in dry and boundary lubrication conditions can be badly damaged by galling and adhesive wear. Newly-developed nano-structured Tungsten/Tungsten Carbide Hardide coatings produced by Chemical Vapour Deposition (CVD) are proven to prevent galling and protect critical parts against wear and corrosion. The CVD coating is crystallised atom-by-atom from gas mixture and can be applied uniformly on both internal and external surfaces of complex shaped parts such as bearings, hydraulic cylinders and gears. 50 microns thick Hardide-T CVD coating combines high hardness (~1200Hv or 72HRC) with toughness, which give it enhanced abrasion resistance, outperforming Hard Chrome by a factor of 14x and HVOF WC-Co (12%) by 3x. The coating is free from porosity and provides an effective barrier against corrosion. The coating galling wear was tested using a TE77 tribometer cylinder-on-flat reciprocating dry sliding rig under gradually increasing loads. Baseline uncoated 316 stainless steel tested against the same 316SS galled quickly, the friction coefficient exceeded 1.0 and samples seized under moderate contact pressure of 200 MPa. Uncoated 316SS sliding against Hardide-coated 316SS demonstrated improved galling-resistance, but this combination still showed adhesive wear when dry friction coefficient reached 0.52 at 390 MPa contact pressure. The CVD coating tested against the same CVD coating showed the best anti-galling performance and did not gall even under the maximum contact pressure of 810 MPa, maintaining a low dry friction coefficient of ~ 0.2. This result contradicts the common practice of using dissimilar materials of mating surfaces to prevent galling. The fatigue-resistant coating is used on critical parts of oil drilling tools, heavy-duty pumps and valves to extend equipment service life and performance in abrasive and corrosive conditions. Airbus has qualified Hardide-A CVD coating as a better performing environmentally-acceptable replacement for REACH-restricted Hard Chrome plating.
KEYWORDS: COATING, TUNGSTEN CARBIDE, GALLING-RESISTANCE, WEAR-RESISTANCE, CVD, HARD CHROME REPLACEMENT.
Yuri Zhuk Hardide plc, United Kingdom
INTRODUCTION Galling is a severe form of adhesive wear between two sliding solid surfaces which is distinguished by macroscopic localized roughening and creation of protrusions above the original surface; it often includes plastic flow or material transfer, or both [1]. Galling can affect heavy-duty engine and drivetrain parts and mechanisms operating under high loads in dry and boundary lubrication conditions. The approach most commonly used to prevent galling in mechanical applications is use of oil and grease lubrication which act as a barrier between two sliding surfaces. Special lubricants with anti-seizure additives were developed for galling wear applications. Meanwhile in high contact presLa Metallurgia Italiana - n. 10 2019
sure situations the lubricant can be squeezed out of the sliding contact resulting in boundary lubrication and even in dry metal/ metal sliding. Reliance on oil, grease and polymeric coatings also limits the operating temperature of the sliding contact parts, as oil and polymers degrade above 200oC. Use of oil sets temperature and load restrictions on the mechanisms, requires frequent re-lubrication and maintenance. Another common approach to galling prevention is the use of two dissimilar materials in the sliding pair. The mutual solubility theory views self-mated materials with identical properties as prone to galling under lower loads, irrespective of the materialâ&#x20AC;&#x2122;s ability to resist galling when in contact with other materials [2]. Meanwhile this theory was mainly 45
Industry news based on the data for pure materials and might be unsuitable to describe performance of alloyed and composite materials. With more advanced coatings coming to the market the use of coatings and surface treatments increasingly becomes the preferred solution for prevention of galling and adhesive wear. Metallurgical and ceramic anti-galling coatings typically have high hardness and thus will not be squeezed out of the sliding contact area even under high contact pressure, making them a more durable and less restrictive solution. Ideal anti-galling coatings should have a combination of several characteristics: high
hardness is an important factor, but to perform under high load the coating should also have enhanced toughness, a hard material with low coefficient of friction may never gall but it can fail due to cracking [3]. This paper examines the galling-prevention properties of Hardide® Coatings produced by Chemical Vapour Deposition (CVD) by testing different materials combinations. The paper also describes other properties of this type of coatings useful for automotive and other mechanical applications.
THE MATERIALS USED AND THE SAMPLES PREPARATION
constituents bonded together at the atomic level. The material is free from inter-granular inclusions, impurities, porosity and other defects which can weaken the material’s mechanical properties and make it brittle. The CVD coating from gas media enables coating both internal and external surfaces and complex shapes, which are difficult or impossible to coat by most traditional hard coatings. Unlike traditional spray Tungsten Carbide coatings, the CVD coating does not use a Cobalt or Nickel metal matrix binder and is pore-free. This results in high resistance of the coating to corrosion and chemically aggressive media, making it an effective barrier against corrosion. The Nanostructured CVD Tungsten Carbide Hardide coatings are used to protect critical parts of oil drilling tools, pumps, valves, aircraft components against abrasive wear, erosion and chemically aggressive media. In several applications the Hardide-T coating also proven its ability of preventing galling in dry sliding situations under high contact pressure. Samples for this testing shown on Fig.1 were all made of austenitic 316 stainless steel, some of them were tested in uncoated condition as a baseline control, while others were coated with Hardide-T coating.
This project used Hardide-T type nano-structured Tungsten/Tungsten Carbide composite coating produced by CVD. The coating combines high hardness ranging from 1100 Hv to 1600 Hv (7077 HRC), with enhanced toughness, measured to be greater than 9 MPa*m1/2. This combination of hardness with toughness is benefitial for high load sliding applications as it prevents coating wear and brittle fracture under high loads. The coating is typically 50 microns thick, this thickness gives it enhanced load-bearing capacity. This coating was produced by a low-temperature thermal CVD technology [4] at a temperature of 500oC. This enables coating a wide range of engineering materials: various grades of stainless steel, Inconel, tool steels stable at 500oC, Ni-, Cu-, and Co-based alloys. The coating has a strong metallurgical adhesion to these substrates with the bond strength exceeding 70 MPa. The coating is crystallised from gas media at a low pressure, building a dense pore-free layer with the Tungsten and Tungsten Carbide TESTING METHOD Galling wear was tested using Phoenix TE77 high frequency reciprocating test rig with cylinder-on-flat line contact configuration in dry sliding friction conditions. Fig.1 shows the samples used in the tests and the “cylinder-on-flat” configuration schematic. The 13.5 mm long cylindrical pins with 6 mm diameter were moved with frequency of 0.5 Hz keeping line contact with a 3.90 mm thick flat plates in direction perpendicular to the pin axis. The normal load applied to the pin was gradually increasing during the test from 10 N to the maximum of 800 N, which are
equivalent to a Hertzian contact pressure increasing from 90 MPa to 810 MPa (or until galling was recorded and the test stopped). All the tests were performed at room temperature (22 ± 2 ºC) and relative humidity of 54 ± 5 %. The testing program focused on the effects of the use of different material combinations sliding against each other and the samples initial surface finish. Dry coefficient of friction (CoF) was continuously monitored during each test, and CoF increasing above 1.0 was used as the main indicator of galling.
Fig. 1 - Left: TE77 samples used in the recirocating dry friction tests; right: the “cylinder-on-flat” configuration schematic. 46
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AttualitàTrattamenti industriale termici Three materials combinations were used: - Self-mated uncoated stainless steel, as a baseline combination; - Hardide-coated against uncoated stainless steel; - Self-mated Hardide coatings when both pins and plates were coated. For each combination three different levels of surface roughness were tested, referred to as Ra1, Ra2 and Ra3: Ra1 = 0.1-0.15 µm, Ra2 = 0.3-0.4 µm, Ra3 ~ 0.8 µm. After testing it was found that the initial surface roughness of the samples had some effect on the friction coefficient only at the initial stage of each test, lasting just 60...100 seconds, then the roughness effect disappe-
ared as the mating surfaces either polished each other or galling and adhesive wear started. For this reason only the effects of the materials combinations will be presented in details below, all for the samples surface roughness Ra1 = 0.1-0.15 µm. Detailed results of this galling testing including results for all the tested surface roguhness samples were published in [5]. After the tests the samples were inspected using optical microscopy and using JEOL JSM-6500F scanning electron microscope. SEM’s energy dispersive X-Ray spectroscopy (EDS) was used to check for material transfer. Surface roughness Ra was measured in the wear scars using Alicona Infinite focus instrument.
TESTING RESULTS
galling wear applications. Meanwhile in high contact pressure situations the lubricant can be squeezed out of the sliding contact resulting in boundary lubrication and even in dry metal/metal sliding. Reliance on oil, grease and polymeric coatings also limits the operating temperature of the sliding contact parts, as oil and polymers degrade above 200oC. Use of oil sets temperature and load restrictions on the mechanisms, requires frequent re-lubrication and maintenance. Another common approach to galling prevention is the use of two dissimilar materials in the sliding pair. The mutual solubility theory views self-mated materials with identical properties as prone to galling under lower loads, irrespective of the material’s ability to resist galling when in contact with other materials [2]. Meanwhile this theory was mainly.
Self-mated stainless steel Galling is a severe form of adhesive wear between two sliding solid surfaces which is distinguished by macroscopic localized roughening and creation of protrusions above the original surface; it often includes plastic flow or material transfer, or both [1]. Galling can affect heavy-duty engine and drivetrain parts and mechanisms operating under high loads in dry and boundary lubrication conditions. The approach most commonly used to prevent galling in mechanical applications is use of oil and grease lubrication which act as a barrier between two sliding surfaces. Special lubricants with anti-seizure additives were developed for
Fig. 2 - Test results for self-mated uncoated 316 stainless steel samples: Left: Coefficient of Friction recorded during the test with load increasing gradually from 10 N to 200 N. Right: SEM image of the tested pin contact area after the test showing adhesive wear. Red arrows show direction of the test reciprocating movement. Uncoated 316 austenitic stainless steel samples tested against the same uncoated 316 developed galling and adhesive wear quickly with friction coefficient rising above 1.0 after just 200 second and the samples seizing at relatively low contact pressure of 200 MPa. This is a predictable result for the baseline uncoated samples as stainless steel is known for its poor galling resistance. Fig. 2 Left shows Coefficient of Friction (CoF) plot with many spikes typical for adhesive wear when both metal surfaces are “welded” together. Continuing movement of the samples
against each other then rupture the bond producing debris and material transfer from one sample to another. Fig.2 Right SEM image shows this extensive adhesive wear damage to the tested surface. Average dynamic CoF measured during the test for this pair was very high at 0.978 with standard deviation of 0.267. Visual inspection after the test shown significant damage and rough surface of both pins and plates in the contact areas. Ra surface roughness increased from 0.1-0.15 µm before the test to 4.62 µm on the plate and 4.95 µm on the pin.
Hardide-coated against uncoated stainless steel When one sample in the pair (the pin) was Hardide-coated and another (the plate) left uncoated the galling performance im-
proved significantly. As Fig.3 shows CoF increased initially and shown some spikes indicative of adhesive wear, it remained below 1.0 threshold and later stabilised at the level of ~0.5.
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Industry news The average dynamic CoF for this pair was measured at 0.522 with standard deviation of 0.057. Visual inspection after the test shown that the coating on the pin remained in place, but also had some debris attached to it, as shown on Fig.3 Right. The debris composition was close to that of 316 stainless steel consisting mainly of Iron, Nickel and Chromium.
Surface of both pins and plates in the contact areas became rougher after the test, but to a different degree: Ra increased from pre-test 0.1-0.15 µm to 4.26 µm on the uncoated plate and 2.43 µm on the coated pin.
Fig. 3 - Test results for Hardide-coated pin against uncoated 316 stainless steel plate: Left: Coefficient of Friction during the test with loads increasing gradually from 10 N to 500 N. Right: SEM image of the tested pin contact area showing some polished areas and also signs of material transfer. The material transfer from the uncoated plate to the harder coated pin resulted in the pin contact surface becoming partly covered with stainless steel, changing the tribological system to at
least partly self-mating Stainless steel. This can explain relatively high friction coefficient measured in this system.
Self-mated Hardide coating Self-mated Hardide coating showed the best tribological performance among the three materials combinations. All tested samples from this group shown relatively smooth movement without signs of galling wear or seizure. The dynamic coefficient of friction for this combination is shown on Fig. 4 Left – for loads increasing from 10 N to 500 N, and Right – for further load increase from 500 N to 800 N. The rig maximum load of 800 N produced 820 MPa contact pressure. During the initial running-in stage the CoF gradually increased to 0.25, but later reduced to ~0.2 and remained at this level until the end of the test. Visual-
ly Hardide-coated samples became polished in the contact area with shiny surfaces. The reduction of the CoF after initial stage can be explained by polishing of the contact surfaces, as shown on Fig. 5. Dry Friction Coefficient of ~0.2 is approximately 1/5 of CoF measured in the same test for self-mating Stainless steel. For comparison: the steady state dry friction coefficient of Hard Chrome Plating was reported at 0.55 [6] which is almost 3x higher than in the CVD Tungsten Carbide/ CVD Tungsten Carbide pair. According to [6] various spray Tungsten Carbide coatings also have dry CoF in the range from 0.56 to 0.61.
Fig. 4 - Dynamic coefficient of friction of Hardide-coated pin tested against Hardide-coated plate: Left: with loads increasing gradually from 10 N to 500 N. Right: with loads increasing gradually further from 500 N to 800 N. 48
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AttualitĂ Trattamenti industriale termici
Fig. 5 - Optical microscopy image of the Hardide-coated pin before (left) and after testing (right) against Hardide-coated plate. Area in the red rectangular is the contact area which became polished, with vertical machining marks still visible above and below it. The fact that self-mated Hardide coatings shown the best anti-galling performance contradicts concept of using dissimilar materials to prevent galling. One possible explanation for this contradiction is that many structural metals used in demanding applications depend on protective Oxide film formed on the metal surface: Stainless steel, Inconel and other high-Chromium alloys form thin but dense and passive layer of Cr2O3; Titanium alloys form TiO2; Aluminium alloys Al2O3. When two such surfaces are moving against each other in dry unlubricated conditions the thin passive oxide film is easily damaged by asperities, bringing active metals beneath the oxide in direct contact leading to adhesive wear and galling. The fact that Hardide/Hardide selfmating pair shown no galling even under very high loads involving surface polishing indicates that although Tungsten Oxide layer does exist it does not play critical role in the protection of
the coating material from adhesive wear, as even the oxide-free coating remains inert and resistant to adhesive wear. This galling-resistance of self-mating hard coating gives an additional advantage for applications where mating parts are additionally exposed to abrasive wear and debris. When pairs of dissimilar materials are used to prevent galling one material usually has lower hardness than another, making it less resistant to abrasive wear. Coating both mating parts with hard CVD coating will not only protect them against galling but also against abrasive wear, erosion and corrosion, as detailed in further sections of this paper. This combined protection against several types of wear can be useful for heavy duty diesel engine parts where exhaust gases recirculation introduces highly abrasive carbon particles in the mechanical parts surface.
OTHER PROPERTIES OF THE CVD COATING RELEVANT FOR AUTOMOTIVE APPLICATIONS Resistance to Abrasive Wear Fig. 6 presents the results of modified ASTM G65 standard wear resistance tests [7] for several hard coatings and materials. The abrasive is fed on rotating neoprene rubber-coated wheel rubbing against the coated metal samples, the bars on Fig.6 represent the material volume loss after 6000 rotations. In this test
both Hardide-A and Hardide-T coatings outperformed Hard Chrome Plating (HCP) and HVOF WC-Co (12%). As compared to uncoated Inconel 718 superalloy Hardide-A shown wear rate reduced by factor of x250, while harder Hardide-T coating outperformed Inconel by a factor of x460.
Fig. 6 - Results of ASTM G65 Wear Resistance Test show material volume loss after 6000 counter-body rotation cycles. Note that Inconel 718 lost 562 mm3 which is far beyond the chart scale. La Metallurgia Italiana - n. 10 2019
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Industry news Resistance to corrosion and chemically aggressive media Porosity in most traditional coatings leaves a path for corrosive fluids to attack the substrate. Sealing the porosity using low viscosity organic sealants can help, but it limits the operating temperatures to below 200˚C. In contrast, the CVD Hardide coating is effectively porosity-free as applied and does not require additional sealing. The CVD coating is crystallised from the gas phase atom-by-atom filling the coating micro-pores and defects while the coating is deposited. The Hardide coating’s corrosion performance was benchmarked against other coatings in 480-hour salt
spray tests to ASTM B117-07, using mild steel plates coated with Hardide, HCP and HVOF coatings. Fig.7 shows three coatings samples after testing. The HCP samples were badly corroded and removed from test after just 288 hours exposure. HVOF-coated samples had heavy rust stains and the coating cracked due to the intensive corrosion of the steel plate beneath. The Hardide samples in this test showed only light staining, the pore-free CVD coating effectively protected mild steel from corrosion.
Fig. 7 - Samples of three coatings after 480 hours salt spray corrosion tests: left - HVOF; centre – Hard Chrome (removed from testing after just 288 hours due to heavy corrosion); right – Hardide coating. Toughness, resistance to impact, high loads and substrate deformations Toughness, ability to survive impact and the substrate deformations are properties of significant practical importance for applications involving shock loads and impact such as aircraft landing gear or engine parts. Brittleness and poor impact resistance are among the main drawbacks of traditional WC/Co hardmetals. HVOF WC/Co coatings are known to crack and spall under high load and high cyclic fatigue conditions [8]. The CVD Tungsten Carbide coating’s structure and composition was optimized to maximize its toughness. Fig.8 and 9 illustrate the coating’s ability to survive impact, significant substrate deformations and shock loads without spalling or cracking. Two
methods were used in an attempt to measure the CVD Tungsten Carbide coating’s fracture toughness: repeat nano-indentations and making deep indentations using various indenters (diamond cube corner, Berkovich, Knoop and spherical indenters). In the first test, the sample did not fracture after 100 nano-impacts all the characteristics measured by the nano-indenter including the coating dynamic hardness, dynamic depth, and coefficient of restitution showed no signs of brittle behavior. The diamond cube corner indentation (shown on Fig. 8 Left) also failed to introduce cracks into the CVD Tungsten Carbide coating. In both these tests the coating did not show brittle behavior, thus its fracture toughness exceeded the level of 9 MPa*m1/2 which is considered the maximum that can be measured by these methods.
Fig. 8 - Illustrates enhanced toughness and ductility of CVD Hardide coatings: Left: Diamond cube corner indenter with sharp edges produced no cracks, which in brittle coatings extend from the indentation corners. Center: Steel test ring with 50 microns CVD coating crushed to test coating adhesion and toughness – no flaking or coating separation from substrate. Right: CVD-coated Inconel pin/bush assembly survived intense repeated hammer impacts without fracture or flaking despite significant deformations of the substrate 50
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AttualitàTrattamenti industriale termici Traditional views on CVD coatings describe them as being too thin and brittle to withstand high loads. In contrast to these stereotypes the Hardide-A coating showed good performance when tested under very high loads. Fig.9 shows a cross-section of an aircraft part with Hardide coating after a dry reciprocating sliding wear test under high loads >1100 MPa. The loads exceeded the
S99 steel substrate yield strength of 875 MPa. The substrate deformed plastically, the coating started to crack and form a “stepped” structure with steel extrusion flowing between the damaged coating blocks. Meanwhile in these extreme conditions the coating remained adhered even when separate coating “blocks” were pushed deep into the steel base.
Fig. 9 - Cross-section of a coated sample after high load dry sliding wear test. The loads in excess of 1100 MPa plastically deformed the substrate steel, but the coating remained adhered
Ability to Coat Complex Shape Parts
Fig. 10 - Examples of complex shaped parts with Hardide coating: ball valves, bushes, a flow diverter; good surface finish can be achieved by polishing, without grinding.
CVD gas-phase technology enables uniform conformal coating of complex shapes and internal surfaces, some examples are shown on Fig. 10 and 11. For many traditional coating technologies, internal surfaces or complex shape items are very difficult or even impossible to coat due to the line-of-sight nature of the coating processes. For example, spray coatings (HVOF, Plasma Spray, D-gun) can be applied to external surfaces which are easily accessible for the spray gun under optimal angle with the spray nozzle kept at a distance to prevent overheating of the part being coated. Spray coatings have a number of other limitations: due to residual stresses they are not suitable to coat both faces of either internal or external corners, or smaller cylindrical surfaces with a diameter <50 mm, thin-wall parts, etc. When spraycoating complex shaped items, it is difficult to avoid building a La Metallurgia Italiana - n. 10 2019
thicker coating layer on the more exposed edges while applying a thinner layer in the “shadowed” areas. This can distort the part shape. Similar limitations exist for PVD coatings, where planetary rotation of the parts being coated helps improve coating uniformity on the external surfaces, but the PVD coatings cannot be deposited uniformly inside deep holes. Electrolytic processes such as Hard Chrome plating can coat inside but often build a thicker layer on the edges where the current density is higher thus creating the “dog-bone” shape, which often requires post-coating grinding. The only exception is electroless metal plating such as Ni-P which can produce a uniform coating on both internal and external surfaces and complex shapes.
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Industry news
Fig. 11 - Scanning Electron Microscopy image of Hardide-T CVD coating cross-section: the 50-microns thick coating (marked with arrows) was applied on both internal and external corners with uniform thickness. CVD coatings are applied from gas media at a low pressure. The coating is crystallized atom-by-atom from the gas phase on every hot surface in contact with the reactive gas mixture at approximately the same rate. So if the gas mixture flows through a hydraulic cylinder bore, the coating will grow uniformly inside the
part. It does not depend on the current density thus is free from the “dog-bone” effect, as shown on Fig.11 presenting SEM image of a cross-section of a Hardide-coated part with both internal and external corners.
CVD TUNGSTEN CARBIDE COATING APPLICATIONS
As detailed in this paper, in testing, the CVD Tungsten Carbide coating has proven to be more than an equivalent replacement for Hard Chrome; it outperformed HCP in several important characteristics, in particular as a better barrier against corrosion and also showed enhanced fatigue properties. The CVD Tungsten Carbide coatings are particularly suitable for the following types of applications: - Complex shaped parts, or where internal surfaces need protection against wear and erosion. - Applications involving high loads, shock loads, impact, operating in high load sliding wear conditions where risk of galling is present. For these applications coating both mating parts is recommended. - Parts working against seals, such as hydraulic pistons and cylinders, actuators, gearbox rotating shafts: the coating retains a good finish in operation which reduces the seal’s wear [9]. - Fatigue-critical parts. Parts with complex shapes and parts operating under high loads/ risk of galling are illustrative of the applications most suitable for coating with CVD Tungsten Carbide as these parts are often difficult to coat by other technologies. CVD Tungsten Carbide coatings can complement other coatings adding to the range of surface engineering materials that automotive engineers can use when working on the challenging task of replacing Hard Chrome plating.
CVD Tungsten Carbide coatings are widely used in the oil and gas and flow control industries on critical parts of drilling and downhole tools, severe service metal-seated ball valves, pistons and cylinders of positive displacement pumps, industrial tooling and bearings. The coating can dramatically increase the life of critical parts operating in abrasive, erosive and corrosive environments. Longer lasting coated parts reduce expensive downtime and enhance the reliability of expensive equipment such as directional oil and gas drilling tools. Reduced wear of critical components helps maintain optimum performance for longer. In the aerospace sector, the CVD coating has been used on components on the Eurofighter Typhoon jet since 2005. Starting from 2017 the production of Hard Cr Plating in the EU was restricted by REACH regulations due to use of highly toxic Hexavalent Chromium salts. This opens wider opportunities for the use of CVD coatings in the automotive and aerospace sectors. In 2015 after several years of extensive in-depth testing the CVD Tungsten Carbide/Tungsten coating passed Airbus technical qualification requirements as a replacement for Hard Chrome. In 2017, the Hardide coating production facilities and quality control system passed stringent audit by Airbus and the company gained the status of an approved Airbus supplier, enabling the coating of flying aircraft parts. This opens a wide range of new coating applications on various aircraft components such as landing gear, hydraulic actuators and wing flap mechanisms. 52
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AttualitàTrattamenti industriale termici CONCLUSIONS Hardide is a family of advanced CVD Tungsten Carbide coatings used to increase the life of critical metal parts operating in abrasive, erosive and chemically aggressive environments. The paper presented results of TE77 reciprocating dry sliding wear tests investigating the coating galling-resistance. The self-mating CVD coated samples did not fall under the maximum contact pressure the test rig could produce 810 MPa. In dry friction conditions the coated samples maintained the friction coefficient at ~0.2 level, which is 5x times lower than friction between uncoated stainless steel parts. In applications where risk of galling is present, the coating of both mating surfaces is recommended. The ability of the gas-phase CVD technology to coat internal surfaces and complex shapes opens new potential applications for hard coatings on critical parts which could not be coated by traditional thermal spray and other line-of-sight coating methods. The nano-structure and composition of Hardide coatings allowed achieving an unusual combination of high hardness with enhanced toughness and ductility. This combination increases the coating’s wear and erosion resistance, and its ability to survive impact and part deformation. The CVD coating is pore-free and
provides a very effective barrier against corrosion and aggressive chemicals. The coating as deposited is under compressive stress. Together with the coating’s toughness and uniform structure this can achieve enhanced fatigue resistance of the coated parts. The coating has strong metallurgical adhesion to steel and high loadbearing capacity withstanding loads of 1.2 GPa without cracking or chipping. Due to its structure the coating is seal-friendly and reduces wear of the mating seals and bearings. The coating has already been used on critical parts of oil drilling tools and flow control equipment for 15 years, and on flying aircraft parts for more than 12 years. The recent qualification by Airbus opens a wide range of applications on civil aircraft components including landing gear, actuators, wing flaps and locking mechanisms. The CVD coating is especially beneficial for parts with complex shapes and also parts operating under high loads where it can outperform other available HCP alternatives. Hardide CVD coatings can be a useful addition to the selection of materials and surface treatments automotive engineers are using to enhance performance and reduce maintenance of critical drivetrain parts.
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Astm G 98-02, “Standard Test Method for Galling Resistance of Materials,” Standards, vol. 02, no. Reapproved, pp. 1–4, 2002.
[2]
I. Hutchings, Tribology Friction and Wear of Engineering Materials, First Edit. Oxford: Butterworth Heinemann, 1992.
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J. Vikstrijm, “Galling resistance of hardfacing alloys replacing Stellite,” vol. 179, pp. 143–146, 1994.
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www.hardide.com/coatings/coatings-range/ Accessed 25 April 2019
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C. Micallef, Y. Zhuk, and R. J. K. Wood, “Galling resistance of nanostructured CVD tungsten/tungsten carbide coatings,” Surf. Topogr. Metrol. Prop., 2019.
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David Lee “Performance Testing of HVOF Coatings and Comparison with Hard Chrome Plate (HCP)”, HCAT/CHCAT Meeting December 14, 2000, by Deloro Stellite Co. Inc. - Stellite Coatings Div.
[7]
Gant AJ, Gee M G, Roebuck B, “Rotating wheel abrasion of WC/Co hardmetals”, Wear 258 (2005) 178-188.
[8]
“An Updated Thintri MARKET STUDY: 2009: Chrome Plating Alternatives: Thermal Spray, Electroless Plating, and Others” from http://www.thintri.com/chrome-plating-report.htm Accessed 27 July 2018.
[9]
Konyashin, I., Ries, B., Hlawatschek, D., Zhuk, Y., Mazilkin, A., Straumal, B., Dorn, F., Park, D. “Wear-Resistance and Hardness: Are They Directly Related for Nanostructured Hard Materials?” Int. Journal of Refractory Metals and Hard Materials, Vol.49, (2015), pp. 203-211
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Kilometro Rosso Bergamo (Italy)
21-22 November 2019 International Meeting
METALS FOR ROAD MOBILITY
Scope
The Meeting will focus on the development and innovation of the metallic materials used for the car body structures, equipments and the new propulsion architectures that will condition the mobility technologies in the next years. The attention will be devoted, in particular, on cast irons, steels, aluminium and magnesium alloys. The discussion will involve the microstructural aspects and the technological ones for the production of high performance and reliable materials to fit the safety and the environmental requirements of both current regulation and its upcoming evolution. MeMo aims at providing a forum for information transfer among producers, components manufacturers and suppliers, research and academics, designers and engineers, active in the mobility (automotive and trucks) field. Trends and future of automotive sector will be widely discussed during the plenary session, organized by Siderweb in the second day of the Meeting. AIM is looking forward to welcoming you in Bergamo!
Organized Secretariat
Via F. Turati, 8 · 20121 Milano · Italy Tel. 02-76021132 / 02-76397770 e-mail: aim@aimnet.it · www.aimnet.it/memo.htm
Program • Opening lecture by Alberto Bombassei, President and Founder of Brembo SpA (*) • Vehicles and alloys: is it a marriage due to continue? Perspectives for the use of metals for automotive Auto e metalli: un matrimonio
destinato a continuare? Le prospettive per l’impiego di metalli nell’automotive (*) • Automotive: current and future scenarios Automotive: gli scenari di oggi e di domani (*) • Italy: perspectives for the metallurgical industry in the automotive sector Italia: mappa e prospettive del settore dei metalli per il
comparto automotive (*) • Future challanges and how to face them: players’ debate Come
affrontare il domani: operatori a confronto (*) • Round table Tavola rotonda (*) (*) PRESENTATIONS IN ITALIAN LANGUAGE. A SERVICE OF SIMULTANEOUS TRANSLATION (ITALIAN-ENGLISH) WILL BE PROVIDED.
• Perspective and opportunities for electric vehicles Prospettive e
opportunità per l’auto elettrica • The evolution of the vehicle L’evoluzione del veicolo • Development of new bainitic forging steels for automotive and truck applications • Metallurgical design and production of AHSS grades DP800 and CP800 by thin slab ISP and ESP technology at • Acciaieria Arvedi in Cremona, Italy • Effect of Niobium in the 3rd Gen steel • Development of a new methodology for online measurement of austenite fraction formed during continuous annealing process. Part I: Measure on high strength microalloyed steel grades • New light corrosion resistant Fe-Al-Mn-C alloy • Materials, design and manufacturing for lightweight vehicles: a solution to low carbon mobility • Safety and lightweight innovations for future mobility by using stainless steels • Soft Magnetic Composites for current and future Automotive applications • Generation of compensated geometry for Ti6Al4V Formula 1 aerodynamic parts produced by PBF • Corrosion behavior of AlSi10Mg alloys produced by additive manufacturing for automotive • ESR and PM tool steels for the production of high performance components for racing motorbike • Hot work tool steel solutions for light alloys diecasting in automotive industry • Dimensional stability on fatigue performance of wheel bearing rolling elements: case studies • Metal Powder Solutions designed to promote future growth of the PM Industry • Control of stiction phenomena in brake systems • A comparison of the overaging behaviour, tensile and fatigue strength at room and high temperature of cast AlSiMg and AlSiCuMg alloys for engine components • A systematic experimental approach to evaluate knocking induced damage on forged aluminum pistons after bench tests • Laserscanning&3D modelling technology as a valid tools for risk analysis: experience in siderurgic and metallurgic field
www.aimnet.it/memo.htm 54
La Metallurgia Italiana - n. 10 2019
Attualità industriale Benefits and properties of laserhardened tool steel surfaces edited by: S. Nĕmeček, I. Černý, J. Kec, N. Ganev, J. Čapek
Laser hardening brings a great increase in part life and it is an industry-proven process, typically used for hardening of molds, gears, shafts etc. It was mainly the recent advancement of diode lasers that turned it into a competitive process. The paper describes the nature of residual stresses, the influence on hardness and fatigue properties of such parts. Investigations have to be performed in order to gain a comprehensive knowledge about effects on microstructure, hardness, surface properties of treated materials, but also mechanical an particularly fatigue properties. Concerning fatigue resistance of material treated with this technology, results and knowledge recently published in the literature indicate that fatigue resistance can be either reduced or increased, even considerably, depending on numerous parameters of basic material, laser hardening parameters etc. This contribution contains results of a partial study of effect of laser hardening of relatively small specimens on fatigue resistance of 42CrMo4 steel. Two different parameters of the treatment were used, namely two speeds of laser beam on the material surface at constant beam energy. Unlike the lower speed, when fatigue resistance was slightly reduced, higher speed of laser beam resulted in a slight increase of fatigue resistance and fatigue limit. The results are discussed considering an occurrence of residuals stresses.
KEYWORDS: LASER – HARDENING – STRUCTURE – RESIDUAL STRESSES – CRACS – DEFORMATION – PROPERTIES
Němeček Stanislav
RAPTECH s.r.o., U Vodárny 473, 330 08 Zruč-Senec, Czech Republic
Černý Ivo, Kec Jan
SVÚM Praha, Podnikatelská 565, 190 11 Praha 9, Czech Republic
Ganev Nikolaj, Čapek Jiří ČVUT Praha, Trojanova 13, 120 00 Praha, Czech Republic
INTRODUCTION The advantage of laser hardening lies in its ability to alter the microstructure, properties and residual stresses in a workpiece within a localized area. The nature of the laser beam allows even hard-to-reach locations to be locally hardened in relatively short process times [1]. Thanks to a rather small size of the spot being hardened at any given moment, the surrounding material acts as a cooler, thus eliminating the need for external cooling media. Laser hardening is characterized by high cooling rates reaching up to 1000 K/s [2]. Some researchers [3] report cooling rates as high as 104 to 108 K/s and thermal gradients between 105 and 108 K/m. At such values, metastable microstructures with novel properties are obtained. Where cooling rates are this high, almost all austenite in steels transforms to martensite. In addition, the growth of crystallites within polycrystalline materials is severely limited during laser hardening [4,5]. The temperature of the laser-hardened surface and the depth of the heat-affected zone within the sample increase with the square of dwell time of the beam at the point in question [6]. La Metallurgia Italiana - n. 10 2019
The choice of the laser beam size, the speed of movement of the specimen and the laser power govern the surface temperature which in turn affects the resulting hardness and hardening depth. A question remains: what are the fatigue properties of a laser-hardened surface in comparison with, for instance, bulk or induction hardening? This question is of particular importance when it comes to parts operated under dynamic loads, such as gear wheels. Literature sources suggest that under favourable conditions and with optimised parameters, an increase in fatigue strength of up to tens of percent can be achieved [7, 8]. Residual stresses and their distribution within the surface or subsurface layers play an important role here [9]. An additional issue is the impact of laser surface treatment on propagation of an existing crack under fatigue load [10, 11]. In this case, too, substantial crack growth retarding effect can be achieved with laser-treated layers [12]. This paper deals with interrelationships between laser hardening parameters, microstructure, residual stress and fatigue life aspects. 55
Industry news EXPERIMENTS Surfaces of experimental specimens were laser treated using the facilities of the RAPTECH company, 3kW high power diode laser. Two specimens of 42CrMo4 steel were examined. The LOW specimen was processed at the laser beam velocity of 3 mm/min and the hardening temperature of 1200 °C. The HIGH specimen was treated at the laser beam velocity of 4 mm/min and the hardening temperature of 1200 °C). The cuboid specimens of the 42CrMo4 steel had a size of 8×8×120 mm. Prior to laser hardening, the specimens were affixed to a base in groups of three in order to provide symmetric dissipation of heat from the middle piece. To observe scale size effect, the LARGE specimen with the size of 190×50×20 mm made from the Czech Standard ČSN 12050 steel (eq. with C45 steel grade) was processed at
the laser beam velocity of 3 m/min. The hardening temperature of 1200 °C was kept constant and equal in all specimens with the aid of a pyrometer. Three-point bending fatigue tests of LOW and HIGH samples were conducted by the SVUM company in Prague, using at constant stress ranges and the load cycle asymmetry of R = 0.1 (Fmin / Fmax). The loading frequency was 40 Hz. The tests were performed in a SCHENCK PHG machine. The spacing between supports was 120 mm. X-ray diffraction measurement was conducted in the Laboratory of Structural Rentgenography at the Faculty of Nuclear Sciences and Physical Engineering of the Czech Technical University in Prague.
RESULTS
Residual Stress and Hardness Depth Profiles In hardness profile measurement, the maximum hardness is often found not on the surface but in the depth of several tenths of millimetre below the surface in many materials. This fact has various consequences, such as acceptance inspection disputes in cases where the prescribed hardness “has not been met”. Another example is the machining process. Laser-hardened surfaces need no final grinding. This is very beneficial with gear wheels, which are costly to grind. The allowance for grinding (e.g. 0.2 mm for correcting distortion in long parts) may thus contribute to higher final surface hardness. Another factor is the potential surface decarburisation during manufacture of a semi-finished steel product. Differences found in machined components are therefore less dramatic than in the above mentioned case. Considering the known relationship between the carbon content and
hardness of material, one can assume, that residual stresses will vary as well. This assumption was confirmed by the present experiment. The surface in the laser track on the LARGE specimen of the ČSN 12050 (C45) material was gradually etched away and residual stresses were measured by X-ray diffraction, Fig. 3. Deleterious tensile stress was detected in the surface, which can contribute to crack initiation. At the same time, no more than 0.1 millimetre below the surface, in a depth which was reached by removing the oxidised layer, the stress is zero. With increasing depth, favourable compressive stresses prevail, which tend to inhibit crack propagation. Illustrative evidence is shown in Fig. 1. The soft surface exhibits hardness of 320 HV. Below the surface, hardness increases and peaks at 550 HV in the depth of 0.5 mm, Fig. 2. The same depth is the point where maximum compressive stresses were found.
Fig. 1 - Residual stress depth profile from the surface through the hardened layer
Fig. 2 - HV1 hardness depth profile from the surface through the hardened layer
Fatigue tests Reference tests on the base material without laser hardening were performed using specimens of two variants: 42CrMo4 and 42CrMoS4 (the same material with an increased level of sulphur). Hardening depths in LOW and HIGH specimens evaluated from macrostructure observation on metallographic sections through the surface ranged between 1.5 and 1.8 mm. The harde-
ning depth determined as the depth where microhardness drops to one half was much less: between 0.5 and 0.8 mm.
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The large scatter in data (Fig. 3) is surprising and quite unusual, namely with the 42CrMoS4 material. (The reliability expressed as R2 is no more than 0.22.) To some extent, it was encountered in the case of the 42CrMo4 material as well: the R2 is no more
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AttualitàTrattamenti industriale termici than 0.33. Fatigue properties show that the two versions of the material exhibit differences which had not been detected in prior composition analysis, microstructure observation or hardness tests. Scatter was found predominantly in fatigue strength data. Unexpected premature failures occurred in several cases. In this respect, the 42CrMoS4 variant performed more poorly than 42CrMo4. The obvious cause is the large amount of sulphide inclusions. On the other hand, the results of fatigue tests of specimens hardened at higher beam velocity (HIGH specimens) are very promising. The scatter is very small (R2 = 0.99) and the treated mate-
rial exhibits clearly higher fatigue strength than the feedstock in initial condition. This surface treatment has virtually eliminated scatter and enhanced the fatigue performance throughout the entire loading range. This effect may be related to the relatively high compressive residual stresses which have been detected by structural X-ray methods at the Faculty of Nuclear Sciences and Physical Engineering of the Czech Technical University (Ganev N. et al). Namely, where fatigue cracks form from surface or subsurface microstructural defects, their growth at the short crack stage is impeded or even arrested by the compressive residual stress [11].
Fig. 3 - Fatigue S-N curve for different groups of specimens The continuous nature of {211} α-Fe diffraction lines recorded in laser-hardened areas of the surface is an evidence of the presence of fine-grained isotropic polycrystalline material in the surface layer with a thickness of approx. 3 to 5 μm. This is the effective depth of penetration of the X-rays into the examined material. Residual stresses detected in laser tracks exhibit qualitative and quantitative differences. In the large LARGE specimen, tensile stresses between 230 and 300 MPa were found in the longitudinal direction. By contrast, both small LOW and HIGH specimens contain compressive residual stresses whose absolute values increased in the direction of the laser beam movement.
In the HIGH specimen hardened at a higher laser beam velocity, the stresses are approx. 100 MPa lower (-5 → -58 → -130 MPa) when compared to the LOW sample (-139 → -165 → -235 MPa). In the transverse direction, tensile residual stresses were found in all cases, except the 2C region in the HIGH specimen. A higher hardening velocity at higher thermal load on the specimen therefore leads to higher residual stresses. Hardening at a velocity of 3 mm/s, which was the lower velocity variant (LOW specimens), did not improve fatigue properties and in part led to their deterioration. In this case, too, the longitudinal compressive residual stresses were considerably lower and even negligible.
DISCUSSION AND CONCLUSIONS X-ray analysis is a well-proven method for finding residual stresses but it is limited by its small penetration depth and sensitivity to preparation prior to measurement. Consequently, the results obtained with surfaces without grinding were distorted. Depth profiles (obtained by gradually etching the surface away), on the other hand, brought a good agreement with the hardness profile.
and enhanced the fatigue performance throughout the entire loading range. The hardening set-up used for the LOW specimen, i.e. the lower velocity variant, did not improve fatigue properties and in part led to their deterioration.
In 42CrMo4 and 42CrMoS4 steels, the results of HIGH specimens are promising. They were hardened using a higher beam velocity (1200 W and 4 mm/s), where the scatter is very small and the resulting fatigue strength is clearly higher than that of the initial feedstock. This surface treatment has virtually eliminated scatter
With both laser hardening procedures, interrelationship exists between the measured residual stress in the surface and subsurface residual stresses. Upon the procedure which led to enhanced fatigue properties, more notable compressive stresses were found in the direction of the main fatigue load. In the material which exhibited slight decrease in fatigue resistance, these compressive stresses were much smaller, even negligible.
Acknowledgment Contribution is supported by the Ministry of Education, Youth
and Sports of Czech Republic in Inter Eureka project E!11711 Astracomp.
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Industry news REFERENCES [1]
J.C. Ion, Surface Engineering 18 (2002) 14-31
[2]
N.S. Bailey, W. Tan, Y.C. Shin, Surface & Coatings Technology 203 (2009) 2003–2012
[3]
A. Roy, I. Manna, Mat. Sci. Eng. A297 (2001) 85-93
[4]
C. Soriano, J. Leunda, J. Lambarri, V. García Navas, C. Sanz, Applied Surface Science 257 (2011) 7101–7106
[5]
S. Němeček, LaRoux Gillespie ed.: Design for Advanced Manufacturing: Technologies and Processes, McGraw-Hill 2017, ISBN13: 978-1259587450, ISBN-10: 1259587452
[6]
J. Benedek, A. Shachrai, L. Levin, Optics and Laser Technology 12 (1980) 247-253
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Jeddi D., Sidhom H., Lieurade H.P., Nizery F., Fabbro R.: Surface Engineering, 2001, Vol. 17, No. 5, p. 371-378
[8]
De la Cruz P., Odén M., Ericsson T.: International Journal of Fatigue, 1998, Vol. 20, No. 5, p. 389-398
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Jablonski F., Varvarikes J.: Materials Science and Engineering: A, 2008, Vols. 483-484, p. 440-443
[10]
Kocanacuteda D., Kocanacuteda S., Tomaszek H.: Materials Science, 2001, Vol. 37, No. 3, p. 374-382
[11]
Tsay L.W., Lin Z.W.: Fatigue & Fracture of Engineering Materials & Structures, 1998, Vol. 21, No. 12, p. 1549-1558
[12]
Doong J-L., Chen T-J., Tan Y-H.: Engineering Fracture Mechanics, 1989, Vol. 33, No. 3, p. 483-491
[13]
ROWSHAN R.: Process control due laser transformation hardening. Doctoral work, University of Miskolc 2007.
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Scenari
Trattamenti termici
OVERVIEW ON STEELS, HEAT TREATMENT AND SHOOT PEENING IN AUTOMOTIVE FIELD edited by: E. Morgano, J. Tatti Overview on steel, heat treatment and post process operation as shoot peening used for automotive components, in powertrain application. The aim of this form is to understand which are the different raw materials used and the possible modification in terms of chemistry, heat treatment and shoot peening useful to achieve the mission of the component during the vehicle life, in terms of: mechanical properties, fatigue limit and wear resistance. Will analyse cost/benefit and warning concerning different alternatives.
KEYWORDS: STEEL – HEAT TREATMENT – SHOT PEENING – AUTOMOTIVE – MICROSTRUCTURE
Enrico Morgano, Jacopo Tatti CRF GML Metals An FCA company, Italy
INTRODUCTION In automotive field, steel is used to produce many components and architectures in different parts of the vehicle (ex. Powertrain, Chassis, Body etc..).This extensive use is due to the possibility of playing for example with chemistry of raw material, heat
treatment and post process techniques as shoot peening, in order to achieve the requirements requests by R&D department during product setting phase.
RAW MATERIAL Starting from eutectoidic phase diagram Fe-C at certain percentage of carbon with a proper cooling rate, we could have different types of steel with different microstructures. This structural aspect is very important because “playing” with chemistry and modeling the rate of solidification is possible to reach different mechanical properties useful to achieve the target. Fixed a certain percentage of Carbon (C), it’s possible to modify the microstructure and to increase the cooling rate using different quenching fluid (ex: water; oil..). Result of quenching depends not only on the fluid, but also on the shape of the raw material, generally OEM prefers to receive
La Metallurgia Italiana - n. 10 2019
in plant material in bar shape, for gear components, for example, instead of rod shape, in fact: ROD: shows a not homogeneous cooling, in fact during the cooling phase (stelmor) and after their conformation, the external turns cool very differently, compared to the internal ones; this induces disuniformity (shape factor) BAR: the cooling process of bar is homogeneous, it doesn’t induce disuniformity.
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Expertsâ&#x20AC;&#x2122; corner
Fig. 1 - Fe-C phase diagram
Fig. 2 - TTT Curve (Temperature Time Trasformation)
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Fig. 3 - CCT Curve (Coontinuos cooling rate) for steel, at different cooling rate itâ&#x20AC;&#x2122;s possible to achieve different microstructures at high speed we have only Martensite at lower speed Ferrite+ Pearlite
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Trattamenti termici
Fig. 4 - Mechanical properties (Vickers Hardness) depending on microstructure
Fig. 5 - bar and rod shape
CASE HARDENING STEELS FOR GEAR APPLICATIONS Case hardening steels are used in gear application for their good wear resistance and for their excellent fatigue limit. Carbon content is in 0.10% รท 0.25% range, considering Fe-C diagram, in order to avoid brittleness in the core of component. Wear resistance and fatigue are assured by adding alloying elements as Chrome (Cr), Molybdenum (Mo), Manganese (Mn);
Nickel (Ni), Boron (B), Vanadium (V) in different percentage: - Manganese 1.5% max percentage admitted - Nickel: 3,5% max percentage admitted - Chrome: 1.5% max percentage admitted - Molybdenum 0.5% max percentage admitted
Fig. 6 - Common case hardening steels used in automotive gear with the percentage of major elements alloyed TEMPERED STEEL FOR TRANSMISSION AND SUSPENSION SYSTEM Tempered steels are commonly used for transmission and suspension system for their good mechanical properties in terms of La Metallurgia Italiana - n. 10 2019
toughness and fatigue resistance. This steel category is generally surface hardened by local induc61
Experts’ corner tion hardening, and by nitriding in order to improve mechanical properties as fatigue limit and wear resistance. Carbon content is in 0.35% ÷ 0.60% range and the major alloying elements used are Chrome (Cr), Molybdenum (Mo), Nickel (Ni), Manganese (Mn), below are some tempered steel used in the automotive field: - 41Cr4 - 40NiCrMo4 - 40CrMo4 - C43 - C53
Instead of this type of steel, OEM uses micro-alloyed steel as 38MnSiV6 and 48MnV3, in fact this class of steel alloyed with Vanadium (V) and Manganese (Mn) could reach the same mechanical properties of tempered steel without thermal treatment. Carbon content is in 0.35% ÷ 0.50% range and content on Vanadium and Manganese are in 0,70% ÷ 1.50 range. In conclusion starting from a certain percentage of Carbon (C) using different percentage of element as V, Mn, Cr.. etc. and a proper cooling rate with the correct quenching fluids, it is possible to determinate a microstructure and then certain specific mechanical properties for raw materials.
BANDAGE LEVEL The bandage level begins from steelmaking and could affect distortion, after quenching phase. The next annealing isothermal can modify, reducing the level. The area to be investigated for measurement is from the surface up
to a depth of 2 mm. For acceptability of bandage level is define an atlas in order to define in product setting phase the maximum bandage admissible level in components.
Fig. 7 - Example of bandage level and indication of maximum admissible level for automotive components
PRELIMINARY HEAT TREATMENT The preliminary heat treatments have the purpose to guarantee a good machinability and to reduce distortions, after carburizing and quenching. There are three different types, and all of them guarantee equivalent hardness ranges, positioned between 150200 HB.
It is about: - Isothermal annealing - Globular annealing - Machinability tempering
Fig. 8 - Graph Temperature-Time for different massive heat treatment 62
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Scenari
Trattamenti termici HEAT TREATMENT: CARBURIZING The thermochemical carburizing process is carried out with two different technologies. The traditional one, with atmospheric carburizing and quenching in oil, and the innovative one with low pressure carburizing and
quenching in gas (nitrogen) with different levels of pressure. The latter guarantees many advantages, both on the product and on the process.
Fig. 9 - On the left: atmosphere carburizing: gear with carbon content % C 0.17-0.23 Carburizing temperature 920°C, Oil quenching from 870 Stress reliefe 180°. On the right: low pressure Carburizing: gear with carbon content % C 0.27-0.30 Carburizing temperature 960°C Gas quenching from 940 °C Stress reliefe 180°. OVERVIEW ON HEAT TREATMENT AND SHOT PEENING IN AUTOMOTIVE FIELD The mechanical characteristics that act on the performance of a component are different; let us consider the retained austenite
induced by the quenching after the carburizing process.
Fig. 10 - Overview on all item induced by process and heat treatment and parameters to control The mechanical characteristics that act on the performance of a component are different; let us consider the residual stress induced by the peening operation. Shot peening consists of hitting the surface with a jet of metallic
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hots of hardness similar to the one of the component to be treated, with controlled parameters. It allows reaching high residual stresses in compression, which guarantees the increase of the fatigue limit around 20-25%.
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La Metallurgia Italiana - n. 10 2019
Atti Trattamenti e notizie termici CT CORROSIONE (C) (riunione 4 luglio 2019) Consuntivo di attività svolte • Le “Giornate Nazionali sulla Corrosione e Protezione” si sono svolte a Palermo dal 3 al 5 luglio 2019 in contemporanea con la riunione del CT. Le prime considerazioni sulla manifestazione sono positive, in linea con i numeri relativi alle Giornate Nazionali sulla Corrosione e Protezione di Ferrara nel 2015. Si tratta di un buon risultato grazie in particolare all’importante contributo del comitato organizzatore locale, coordinato dalla Prof.ssa Santamaria. La rivista La Metallurgia Italiana dedicherà ampio spazio alla pubblicazione di una selezione delle memorie.
Iniziative future • E’ stata in via preliminare individuata la sede di Torino per la prossima edizione delle “Giornate Nazionali sulla Corrosione e Protezione” nel 2021.
Stato dell’arte e notizie • Il CT Corrosione ha provveduto al rinnovo delle cariche: vengono nominati all’unanimità come Presidente del CdS Corrosione il Prof. Fabio Bolzoni del Politecnico di Milano, Vicepresidente la Prof.ssa Marina Cabrini dell’Università di Bergamo e Segretario il Dott. Tiziano Bellezze dell’Università Politecnica delle Marche. • Ormellese illustra la situazione di NACE Italia, la cui sezione è stata riaperta e collaborerà con altre associazioni tra cui AIM. Questo dovrebbe portare ad una maggiore conoscenza scientifica dei fenomeni corrosionistici e, si spera, ad un aumento dei partecipanti alla manifestazione “Giornate Nazionali”.
CT AMBIENTE E SICUREZZA (AS) (riunione del 19 settembre 2019) Manifestazioni in corso di organizzazione • Bassani segnala che il corso itinerante Metallurgia Sicura (Aosta 30 ottobre, Vicenza 6 novembre, Brescia 13 novembre) ha già raccolto un buon numero di partecipanti. Si sta verificando il numero dei partecipanti che potrà chiedere i crediti formativi.
La Metallurgia Italiana - n. 10 2019
Iniziative future • Continua la preparazione della GdS su tema ambientale. Il coordinatore Filippini presenta gli interventi già selezionati (tre per il mattino e tre per il pomeriggio), sui quali i presenti esprimono i loro pareri. La locandina sarà pronta a breve e la data della manifestazione è fissata per il 5 dicembre 2019 a Verona, in sede da definire.
Stato dell’arte e notizie • Per proseguire con l’esperienza nata durante la precedente riunione del CT, i presenti condividono esperienze sulla sicurezza e su eventuali incidenti avvenuti nelle rispettive aziende. Ne nasce uno scambio di opinioni molto interessante che fornisce spunti di riflessione e di miglioramento.
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Aim news CT TRATTAMENTI TERMICI E METALLOGRAFIA (TTM) (riunione del 26 settembre 2019) Consuntivo di attività svolte • Corso di Metallografia: si è concluso il quarto modulo sulle leghe leggere. In assenza del coordinatore Bavaro, il presidente Petta illustra l’esito complessivo del corso, che ha visto un buon numero di partecipanti nelle sue varie fasi. Il giudizio espresso nei questionari di soddisfazione è tra il buono e l’ottimo; in particolare, è stata molto apprezzata l’attività pratica di laboratorio.
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Manifestazioni in corso di organizzazione • “Convegno Nazionale Trattamenti Termici 2020” (Genova, 6 e 7 maggio 2020): tutti le informazioni sono sul sito dell’evento: www.aimnet.it/tt.htm; Petta sollecita memorie tecniche sulla filiera dei trattamenti termici per l’area convegno. Ci sarà però anche una zona dello spazio espositivo, che si auspica ricco di espositori di tutta la filiera, per presentare memorie tecnico-commerciali e dare visibilità a espositori e sponsor. Ora si devono definire i dettagli per garantire un ottimale funzionamento del nuovo format del Convegno. • GdS “Dalle due alle quattro ruote: gli aspetti metallurgici” (coordinatore Morgano) si terrà al Museo Piaggio di Pontedera il 7 novembre 2019. Il programma prevede memorie tecniche al mattino, una visita guidata al museo e una tavola rotonda in chiusura di giornata.
Iniziative future • GdS sugli stampi (coordinatore Alberto Rivolta): nella seconda settimana di giugno si svolgerà questa giornata sulla scia di altre precedenti. Il focus sarà l’ottimizzazione del raffreddamento degli stampi e le informazioni di base sui loro trattamenti termici, in particolare la nitrurazione. I dettagli saranno messi a punto nel prossimo futuro. • GdS presso Iveco Torino: il coordinatore Morgano organizzerà questa giornata per il 19 marzo 2020, scegliendo tra presentazioni su prodotti lunghi e relativi trattamenti termici o problematiche degli ingranaggi. Il tema e la locandina saranno definiti entro metà gennaio 2020. • Corso “Metallurgia di base”: il corso sarà riproposto come d’abitudine, attorno alla metà di maggio 2020. I coordinatori verificheranno i dettagli dell’organizzazione e la durata, che potrebbe essere estesa a 4 giornate. • GdS “Trattamenti termici e modellazione”: i coordinatori Pellizzari, Valente e Vicardi organizzeranno questa giornata che potrebbe anche chiamarsi, alternativamente, “La potenzialità della simulazione numerica nei trattamenti termici”. Gli interventi potrebbero coinvolgere anche le leghe di alluminio per i trasporti. • Diverse altre manifestazioni vengono suggerite dai presenti; i temi saranno ridiscussi nelle prossime riunioni.
La Metallurgia Italiana - 66 n. 10 2019
Atti Trattamenti e notizie termici CT CONTROLLO E CARATTERIZZAZIONE PRODOTTI (CCP) (riunione del 30 settembre 2019) Consuntivo di attività svolte • Il corso “Prove Meccaniche” è stato tenuto il 29 e 30 maggio e 5 giugno 2019 a Milano (sede Fast e lab OMECO) ed il 6 giugno a Crema, con visita ai laboratori Element. Il coordinatore Trentini) segnala un numero di partecipanti non soddisfacente, a fronte di un elevato livello di docenti e di contenuto delle lezioni. I questionari di soddisfazione compilati riportano commenti fortemente positivi. Per il futuro si ribadisce l’idea di suddividere il costo tra prove di base e prove specialistiche, con una possibile sinergia organizzativa con il CT Metallurgia Fisica.
La Metallurgia Italiana - n. 10 2019
Manifestazioni in corso di organizzazione • Il corso “PnD” (coordinatori Trentini e Cusolito) si svolgerà il 20 e 21 novembre 2019 a Milano presso la sede Fast. La locandina è stata diffusa per tempo e tutta l’organizzazione è stata portata a termine. • La GdS “Microscopia elettronica applicata alla Failure Analysis” (coordinatore Toldo) è programmata per l’11 ottobre presso la SMT di Pozzo d’Adda (BG). Gli iscritti hanno già raggiunto il numero massimo di posti disponibili.
Iniziative future • Il corso “Analisi chimiche”, visto il successo della precedente edizione del 2016, si potrà ripetere nel secondo semestre del 2020 con il titolo “Analisi chimiche applicate alla metallurgia”. Si cercherà di approfondire, tra l’altro, l’argomento delle tecniche di analisi sul campo, settore in grande sviluppo. • GdS su “Accreditamento e prove interlaboratori”: si discute sulla pertinenza di questo corso con le tematiche del CT CCP: alcuni argomenti sono già trattati nel corso di Analisi Chimiche. Nelle prossime riunioni si deciderà se organizzare il corso ed eventualmente con quali collaborazioni. • GdS sulla corrosione: lo scopo di questa giornata sarebbe quello di inquadrare le problematiche della corrosione per i non corrosionisti. Trentini prenderà contatto con il CT Corrosione per valutare la fattibilità della giornata. • Il corso “Failure Analysis”, che ha avuto una grande partecipazione di pubblico nella precedente edizione (quasi 100 iscritti) sarà ripetuto nel primo semestre del 2020. Nella prossima riunione si definirà lo schema del nuovo corso.
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Aim news CT PRESSOCOLATA (P) riunione dell’1 ottobre 2019 Consuntivo di attività svolte • GdS “Le deformazioni dei getti pressocolati: cause e rimedi”: la giornata, effettuata in una sede prestigiosa, ha avuto una partecipazione di pubblico limitata. Il comitato discute a lungo sulle possibili cause e soprattutto sulle attività da svolgere in futuro per evitare il ripetersi di questa situazione.
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Manifestazioni in corso di organizzazione • Il “Master Progettazione Stampi – II edizione” (coordinatori Timelli, Garlet, Citterio, Martina) è programmato tra settembre e dicembre 2019 a Brescia, suddiviso il 6 moduli. La manifestazione è cominciata e il numero di iscritti e al momento soddisfacente. I commenti saranno comunque espressi dai coordinatori alla fine dell’evento. • La sesta edizione del Corso “Igiene delle leghe” (coordinatore Muneratti), coorganizzato con il CT Metalli Leggeri, si terrà a Carmagnola (TO) presso Teksid Aluminium nei giorni 5 e 6 novembre 2019. I membri del comitato sono invitati a diffondere la notizia. • Il convegno internazionale “High Tech Die Casting 2020” avrà luogo a Vicenza dall’1 al 3 luglio 2020. Il Call for Paper sarà pronto entro ottobre e le memorie più Interessanti saranno pubblicate sul numero di luglio/agosto della rivista AIM. Questo convegno è il più importante organizzato dal CT P e tutti i membri del comitato sono invitati a contribuire alla buona riuscita dell’evento.
Iniziative future • La GdS “Zama HPDC 2021” (coordinatori Pola, Valente) sarà posticipata alla primavera 2021 a causa della mancata disponibilità e del principale esponente internazionale dello zinco. Si attende una conferma della possibile data. • Per la GdS sulla sostenibilità in fonderia si cerca un coordinatore e si discute dei contatti con Assofond, che potrebbe dare aiuto nell’organizzazione. • GdS sulla fatica termica: valente presenterà una bozza di programma alla prossima riunione, dopo aver esaminato i questionari e i commenti relativi alla scorsa edizione. Si cercano anche nuovi temi da affrontare. Stato dell'arte e notizie • Un nuovo membro proveniente dal mondo dell’industria si è proposto per partecipare al CT come membro permanente.
La Metallurgia Italiana - n. 10 2019
Atti Trattamenti e notizie termici AIM - UNSIDER Norme pubblicate e progetti in inchiesta (aggiornamento 30 Settembre 2019)
NORME PUBBLICATE E PROGETTI ALLO STUDIO (ELENCO)
PROGETTI UNSIDER IN INCHIESTA prEN E ISO/DIS - OTTOBRE 2019
NORME UNSIDER PUBBLICATE DA UNI NEL MESE DI SETTEMBRE 2019
prEN - PROGETTI DI NORMA EUROPEI
UNI EN ISO 945-1:2019 Microstruttura della ghisa - Parte 1: Classificazione della grafite mediante analisi visiva. NORME UNSIDER RITIRATE DA UNI NEL MESE DI SETTEMBRE 2019 UNI EN ISO 945-1:2018 Microstruttura della ghisa - Parte 1: Classificazione della grafite mediante analisi visiva. NORME UNSIDER PUBBLICATE DA CEN E ISO NEL MESE DI SETTEMBRE 2019 EN ISO 20074:2019 Petroleum and natural gas industry - Pipeline transportation systems - Geological hazards risk management for onshore pipeline (ISO 20074:2019).
prEN ISO 10426-2 Petroleum and natural gas industries - Cements and materials for well cementing Part 2: Testing of well cements (ISO/DIS 10426-2:2019). EN ISO 11961:2018/prA1 Petroleum and natural gas industries - Steel drill pipe - Amendment 1 (ISO 11961:2018/ DAM 1:2019). prEN 17415-1 District cooling pipes - Bonded single pipe systems for directly buried cold water networks - Part 1: Factory made pipe assembly of steel or plastic service pipe, polyurethane thermal insulation and a casing of polyethylene.
EN 1753:2019 Magnesium and magnesium alloys - Magnesium alloy ingots and castings.
prEN 17414-1 District cooling pipes - Factory made flexible pipe systems - Part 1: Classification, general requirements and test methods.
ISO 15590-4:2019 Petroleum and natural gas industries - Factory bends, fittings and flanges for pipeline transportation systems - Part 4: Factory cold bends.
prEN 17414-2 District cooling pipes - Factory made flexible pipe systems - Part 2: Bonded system with plastic service pipes Requirements and test methods.
ISO 10893-3:2011/Amd 1:2019 Non-destructive testing of steel tubes - Part 3: Automated full peripheral flux leakage testing of seamless and welded (except submerged arc-welded) ferromagnetic steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 1: Change of dimensions of the reference notch.
prEN 17414-3 District cooling pipes - Factory made flexible pipe systems - Part 3: Non bonded system with plastic service pipes â&#x20AC;&#x201C; Requirements and test methods.
ISO 10679:2019 Steels - Cast tool steels. ISO/TS 2597-4:2019 Iron ores - Determination of total iron content - Part 4: Potentiometric titration method. PROGETTI UNSIDER MESSI ALLO STUDIO DAL CEN (STAGE 10.99) - OTTOBRE 2019 ==
La Metallurgia Italiana - n. 10 2019
prEN 993-10 rev Methods of test for dense shaped refractory products - Part 10: Determination of permanent change in dimensions on heating. prEN ISO 6892-1 Metallic materials - Tensile testing - Part 1: Method of test at room temperature (ISO/ FDIS 6892-1:2019). prEN ISO 4947 Steel and cast iron - Determination of vanadium content - Potentiometric titration method (ISO/DIS 4947:2019). prEN 10219-3 Cold formed welded steel structural hollow sections - Part 3: Technical delivery condi-
tions for mechanical engineering purposes. prEN 10210-3 Hot finished steel structural hollow sections - Part 3: Technical delivery conditions for mechanical engineering purposes. EN ISO 10893-1:2011/prA1 Non-destructive testing of steel tubes - Part 1: Automated electromagnetic testing of seamless and welded (except submerged arc-welded) steel tubes for the verification of hydraulic leaktightness - Amendment 1: Change of dimensions of the reference notch; change acceptance criteria (ISO 10893-1:2011/DAM 1:2019). EN ISO 10893-2:2011/prA1 Non-destructive testing of steel tubes - Part 2: Automated eddy current testing of seamless and welded (except submerged arc-welded) steel tubes for the detection of imperfections - Amendment 1: Change of dimensions of the reference notch; change acceptance criteria (ISO 10893-2:2011/DAM 1:2019). EN ISO 10893-3:2011/prA2 Non-destructive testing of steel tubes - Part 3: Automated full peripheral flux leakage testing of seamless and welded (except submerged arc-welded) ferromagnetic steel tubes for the detection of longitudinal and/ or transverse imperfections - Amendment 2: Change acceptance criteria (ISO 108933:2011/DAM 2:2019). EN ISO 10893-8:2011/prA1 Non-destructive testing of steel tubes - Part 8: Automated ultrasonic testing of seamless and welded steel tubes for the detection of laminar imperfections - Amendment 1: Change acceptance criteria (ISO 10893-8:2011/ DAM 1:2019). EN ISO 10893-9:2011/prA1 Non-destructive testing of steel tubes - Part 9: Automated ultrasonic testing for the detection of laminar imperfections in strip/plate used for the manufacture of welded steel tubes - Amendment 1: Change acceptance criteria (ISO 10893-9:2011/DAM 1:2019). EN ISO 10893-10:2011/prA1 Non-destructive testing of steel tubes - Part 10: Automated full peripheral ultrasonic testing of seamless and welded (except submerged arc-welded) steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 1: Change the ultrasonic test frequency of transducers; 69
Aim news change of acceptance criteria (ISO 10893-10:2011/DAM 1:2019). EN ISO 10893-11:2011/prA1 Non-destructive testing of steel tubes - Part 11: Automated ultrasonic testing of the weld seam of welded steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 1: Change the ultrasonic test frequency of transducers; change of acceptance criteria (ISO 10893-11:2011/DAM 1:2019). EN ISO 10893-12:2011/prA1 Non-destructive testing of steel tubes - Part 12: Automated full peripheral ultrasonic thickness testing of seamless and welded (except submerged arc-welded) steel tubes - Amendment 1: Change of acceptance criteria (ISO 10893 12:2011/DAM 1:2019). EN 10222-4:2017/prA1 Steel forgings for pressure purposes - Part 4: Weldable fine grain steels with high proof strength. EN 10222-2:2017/prA1 Steel forgings for pressure purposes - Part 2: Ferritic and martensitic steels with specified elevated temperatures properties. ISO/DIS – PROGETTI DI NORMA INTERNAZIONALI ISO/DIS 22605 Refractories - Determination of dynamic Young’s modulus(MOE) at elevated temperatures by impulse excitation of vibration. ISO/DIS 21736 Refractories - Test methods for thermal shock resistance. ISO 11961:2018/DAmd 1 Petroleum and natural gas industries - Steel drill pipe - Amendment 1.
ISO 10893-8:2011/DAmd 1 Non-destructive testing of steel tubes - Part 8: Automated ultrasonic testing of seamless and welded steel tubes for the detection of laminar imperfections - Amendment 1: Change acceptance criteria. ISO 10893-9:2011/DAmd 1 Non-destructive testing of steel tubes - Part 9: Automated ultrasonic testing for the detection of laminar imperfections in strip/plate used for the manufacture of welded steel tubes - Amendment 1: Change acceptance criteria. ISO 10893-10:2011/DAmd 1 Non-destructive testing of steel tubes - Part 10: Automated full peripheral ultrasonic testing of seamless and welded (except submerged arc-welded) steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 1: Change the ultrasonic test frequency of transducers; change of acceptance criteria. ISO 10893-11:2011/DAmd 1 Non-destructive testing of steel tubes - Part 11: Automated ultrasonic testing of the weld seam of welded steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 1: Change the ultrasonic test frequency of transducers; change of acceptance criteria. ISO 10893-12:2011/DAmd 1 Non-destructive testing of steel tubes - Part 12: Automated full peripheral ultrasonic thickness testing of seamless and welded (except submerged arc-welded) steel tubes - Amendment 1: Change of acceptance criteria. ISO/DIS 10802 Ductile iron pipelines - Hydrostatic testing after installation. ISO/DIS 10426-2 Petroleum and natural gas industries - Cements and materials for well cementing - Part 2: Testing of well cements.
ISO 10893-1:2011/DAmd 1 Non-destructive testing of steel tubes - Part 1: Automated electromagnetic testing of seamless and welded (except submerged arc-welded) steel tubes for the verification of hydraulic leaktightness - Amendment 1: Change of dimensions of the reference notch; change acceptance criteria.
ISO/DIS 4947 Steel and cast iron - Determination of vanadium content - Potentiometric titration method.
ISO 10893-2:2011/DAmd 1 Non-destructive testing of steel tubes - Part 2: Automated eddy current testing of seamless and welded (except submerged arc-welded) steel tubes for the detection of imperfections - Amendment 1: Change of dimensions of the reference notch; change acceptance criteria.
FprEN – progetti di norma europei
ISO 10893-3:2011/DAmd 2 Non-destructive testing of steel tubes - Part 3: Automated full peripheral flux leakage testing of seamless and welded (except submerged arc-welded) ferromagnetic steel tubes for the detection of longitudinal and/or transverse imperfections - Amendment 2: Change acceptance criteria.
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PROGETTI UNSIDER AL VOTO FPREN E ISO/FDIS – OTTOBRE 2019
== ISO/FDIS – PROGETTI DI NORMA INTERNAZIONALI ISO/PRF 643 Steels - Micrographic determination of the apparent grain size.
La Metallurgia Italiana - n. 10 2019
Genova 2020 6-7 maggio
2
convegno nazionale
trattamenti termici
27th AIM National Conference & Exhibition on Heat Treatment
www.aimnet.it/tt.htm AIM è lieta di annunciare la 27° edizione del Convegno Nazionale Trattamenti Termici, il più autorevole ed affermato evento sui trattamenti termici a livello nazionale. Il Convegno – mostra si svolgerà nei giorni 6-8 maggio 2020 a Genova, dove si tenne la prima edizione nel 1960. Sede dell’evento saranno i Magazzini del Cotone, nel bellissimo contesto del porto antico di Genova. Maggiori informazioni saranno presto disponibili sul sito dell’evento: www.aimnet.it/tt.htm
Segreteria organizzativa Organising Secretariat
Via F. Turati 8 - Milano (Italy) Tel. +39 02 76021132 info@aimnet.it www.aimnet.it
AIM is proud to announce the 27th National Conference on Heat Treatment! The Conference &Exhibition will be held on May 6-8, 2020 at the Magazzini del Cotone, a former cotton storehouse, located in a picturesque setting right on the old harbour of Genoa. All information will be soon available at: www.aimnet.it/tt.htm
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10
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The 10th European Continuous Casting Conference - ECCC 2020 - will be organised by AIM, the Italian Association for Metallurgy, in Bari (Italy) on 17-19 June 2020 with focus on the status and future developments in the casting of steel. The ECCC is a unique forum for the European continuous casting community to exchange views on the status and the future development of the continuous casting process. The Conference program is abreast of the latest developments in control and automation, advanced continuous casting technologies, application of electromagnetic technologies and mechanical devices to improve the core microstructure, the lubrication issues for improving the surface qualities. Steel metallurgical issues will be addressed as well as their physical and numerical simulation. The exchange of experience in operational practice, maintenance and first results from the recently commissioned plants will integrate the program. The Conference aims at promoting the dialogue among the delegates with industrial and academic background and among the participants in former Conferences and new members of the continuous casting community.
Call for Papers - Abstract Submission
1st announcement call for papers
10th european conference on continuous casting 2020
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Bari . Italy 17-19 June 2020
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www.aimnet.it/eccc2020
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Prospective authors wishing to present papers are invited to submit, by 31 October 2019, a tentative title and an abstract of about 300 words (in English), specifying a maximum of two topics for each proposal, to the Organising Secretariat (aim@aimnet.it). The abstract should provide sufficient information for a fair assessment and include the title of the paper, the author’s names and contact details (address, telephone number and e-mail address). The name of the presenting author should be underlined. A poster session might be organized as well. There are two ways to submit papers: • fill in the form on the Conference website at: www.aimnet.it/eccc2020 • send the requested information by e-mail to: aim@aimnet.it.
Exhibition and Sponsoring As an integral element of the event, a technical exhibition will be held during the event. Companies have the opportunity to reinforce their participation and enhance their corporate identification by taking advantage of the benefits offered to them as sponsors of the event. The detailed sponsorship packages will be available on the Conference website: www.aimnet.it/eccc2020 Companies interested in taking part in the Exhibition or sponsoring the event may contact: e-mail: commerciale@siderweb.com tel. +39 030 2540006 THE ITALIAN STEEL COMMUNITY
siderweb
Contacts ECCC 2020 Organising Secretariat AIM - Associazione Italiana di Metallurgia Via Filippo Turati 8, 20121 Milan - Italy Tel. +39 02 76021132 / +39 02 76397770 aim@aimnet.it - www.aimnet.it/eccc2020