La Metallurgia Italiana - Gennaio 2020

La

Metallurgia Italiana

International Journal of the Italian Association for Metallurgy

n. 1 Gennaio 2020 Organo ufficiale dell’Associazione Italiana di Metallurgia.

Rivista fondata nel 1909

La Metallurgia Italiana International Journal of the Italian Association for Metallurgy Organo ufficiale dell’Associazione Italiana di Metallurgia.

House organ of AIM Italian Association for Metallurgy.

Rivista fondata nel 1909

Direttore responsabile/Chief editor: Mario Cusolito Direttore vicario/Deputy director: Gianangelo Camona Comitato scientifico/Editorial panel: Livio Battezzati, Christian Bernhard, Massimiliano Bestetti, Wolfgang Bleck, Franco Bonollo, Bruno Buchmayr, Enrique Mariano Castrodeza, Emanuela Cerri, Lorella Ceschini, Mario Conserva, Vladislav Deev, Augusto Di Gianfrancesco, Bernd Kleimt, Carlo Mapelli, Jean Denis Mithieux, Marco Ormellese, Massimo Pellizzari, Giorgio Poli, Pedro Dolabella Portella, Barbara Previtali, Evgeny S. Prusov, Emilio Ramous, Roberto Roberti, Dieter Senk, Du Sichen, Karl-Hermann Tacke, Stefano Trasatti Segreteria di redazione/Editorial secretary: Valeria Scarano Comitato di redazione/Editorial committee: Federica Bassani, Gianangelo Camona, Mario Cusolito, 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

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La

Metallurgia Italiana

International Journal of the Italian Association for Metallurgy

n. 1 Gennaio 2020 Organo ufficiale dell’Associazione Italiana

Editoriale / Editorial

di Metallurgia.

Rivista fondata nel 1909

Un 2019 positivo e un nuovo anno ricco di sfide, in primis quella ambientale

Federico Mazzolari, presidente AIM ...................................................................................... pag.4

Memorie scientifiche / Scientific papers Trattamenti Termici - Heat Treatments Bainitic transformation during the two-step Q&P process in a lean medium Mn steel containing silicon

S. Kaar, R. Schneider, D. Krizan, C. Béal, C.Sommitsch ......................................................... pag.6

n. 1 Gennaio 2020

Anno 112 - ISSN 0026-0843

High pressure gas quenching: assessment of velocity experimental measurements and steps for model validation E. Vyazmina, L. Bustamante-Valencia, V. Woimbee, G. Michel, A. Belhajria, Ch. David ................ pag.18

Influence of gas nitriding conditions on layer structure formation on greyand white-solidified cast irons (part I) A. Holst, A. Buchwalder, R. Zenker ....................................................................................... pag.28

Material-dependent aspects for the use of nitriding as a single or du-plex treatment in improving the tribological load-bearing behaviour of cast irons (part II)

indice

A. Buchwalder, N. Klose, R.Zenker ......................................................................................... pag.40

Simulation of heat treatment of Jominy specimen to improve quality of automotive gear components

E. Morgano, C. Viscardi, L. Valente ........................................................................................ pag.52

Neural networks-based prediction of hardenability of high performance carburizing steels for automotive applications

V. Colla, M. Vannucci, L. Bacchi, R. Valentini .......................................................................... pag.59

Heat treatment routes utilizing intercritical annealing of automotive medium-mn steel sheets

A. Grajcar, M. Morawiec ........................................................................................................ pag.66

Attualità industriale / Industry news New environmental-friendly developments and applications of the pulsed plasma heat treatment technology in the field of engine, transmission and power train engineering A. Gebeshuber, V. Strobl, T. Müller ..................................................................................... pag.73

Liquid Nitrocarburizing treatment on valves in internal combustion engines M. Legrand, V. Monteux .................................................................................................... pag.78

Non-destructive hardness / microstructure testing of heat-treated parts by mass production, with multiple frequency magnetic induction method A. Horsch ........................................................................................................................ pag.82

Atti e notizie / AIM news Eventi AIM / AIM events ................................................................................. pag.95 Corso Le prove non distruttive ..................................................................... pag.96 Normativa / Regulations ........................................................................... pag.97 Comitati tecnici / Study groups ...................................................................... pag.104 Eventi nazionali ed internazionali / National and international events .......... pag.105

editoriale - editorial

Un 2019 positivo e un nuovo anno ricco di sfide, in primis quella ambientale.

L’anno 2019 di AIM chiude con risultati più che soddisfacenti per l’attività svolta ed il numero di iscritti.

L’attività è stata intensa per quanto riguarda i Corsi e le Giornate di Studio, per un totale di 28 iniziative, a cui si aggiungono due manifestazioni di portata internazionale: la Conferenza Europea

Federico Mazzolari Presidente AIM

sui Trattamenti Termici - ECHT 2019 - tenutasi a Bardolino in giugno e la prima edizione dell’evento internazionale Metals for Road Mobility, svoltasi a Kilometro Rosso a fine novembre, che ha destato un buon interesse di settore con la partecipazione quale principale ospite del Presidente e fondatore di Brembo, Alberto Bombassei.

Best seller dell’anno è stato Bardolino con 200 partecipanti, seguito dalle Giornate Nazionali sulla Corrosione e Protezione e dalla Giornata di Studio sulle leghe di nickel e superleghe sempre con oltre 100 partecipanti.

“

La grande sfida alla quale siamo chiamati a dare una risposta è quella della competitività ambientale.

Federico Mazzolari

Anche per il 2020 l’attività di AIM sarà orientata alla diffusione della cultura metallurgica tramite attività ed iniziative a partire da fine gennaio.

ECC 2020 a Bari, a metà giugno, la Conferenza Internazionale sulla Colata Continua, e HTDC 2020 la conferenza internazionale sul DIE CASTING ad inizio luglio saranno gli eventi chiave del prossimo anno, insieme al 27° Convegno Nazionale Trattamenti Termici a Genova in maggio ed al 38° Convegno Nazionale AIM che si terrà a Napoli dal 16 al 18 settembre, e da ultimo il Primo Meeting Internazionale sulle Materie Prime ed il Recycling a Bergamo a inizio dicembre.

La Metallurgia Italiana - gennaio 2020

pagina 4

editoriale - editorial

La grande sfida alla quale siamo chiamati a dare una risposta è quella della competitività ambientale: tutto il settore dei metalli dovrà instradarsi verso un concetto di metallurgia pulita riducendo sempre più le emissioni e più in generale l’impatto ambientale.

Il problema di fondo è costituito dalla grande necessità di capitali. Le nuove direttive europee comporteranno la necessità di una massa di investimenti che dovrà essere supportata da misure di sostegno a livello europeo e nazionale.

Sotto l’aspetto dei risultati economici il 2019 è stato un anno caratterizzato da una prima metà con volumi e prezzi ancora buoni sulla scorta dello slancio ereditato dal 2018, e da una seconda progressivamente in frenata.

La siderurgia in particolare e l’industria dei metalli sono correlati ai piani governativi di investimento principalmente sulle infrastrutture. L’Italia, e anche l’Europa, dovrà tenerne conto poiché la competitività di un paese è molto condizionata dal sistema logistico che deve dare garanzie di efficienza con bassi costi nella totale sicurezza e col minimo impatti ambientale.

Anno nuovo e nuovo look della nostra rivista. Da questo numero, La Metallurgia Italiana si presenta con una nuova veste editoriale più consona al formato digitale, grazie alle idee che ci hanno portato Siderweb ed i nostri collaboratori di AIM. Oltre ad alcune variazioni di impatto prettamente grafico, quali il font utilizzato e l’eliminazione della asimmetria tra pagine pari e dispari - tipica della rivista stampata - avremo la possibilità di scegliere le pagine cui accedere direttamente dall’indice con un semplice click, così come potremo accedere direttamente ai siti degli eventi AIM e delle Associazioni collegate. Le rubriche saranno riorganizzate per maggiore fruibilità. Inoltre prevediamo l’introduzione di nuove regole per gli autori, che ci consentiranno progressivamente di visualizzare foto e grafici con migliore risoluzione.

Non resta che augurare a tutti un felice 2020 insieme ad AIM!

La Metallurgia Italiana - January 2020

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Memorie scientifiche - Trattamenti Termici

Bainitic transformation during the two-step Q&P process in a lean medium Mn steel containing silicon S. Kaar, R. Schneider, D. Krizan, C. Béal, C. Sommitsch Lean medium Mn Quenching and Partitioning (Q&P) steels with Mn contents up to 5 wt-% have recently gained a lot of interest due to their promising combination of strength and ductility. This steel group is characterized by a micro-

structure consisting of a carbon-depleted tempered martensitic (α’’) matrix and a considerable amount of retained austenite (RA) stabilized by both, C and Mn, in order to ensure a sufficient strain-induced austenite to martensite trans-

formation (TRIP-effect). Furthermore, the hard α’’ matrix contributes to the excellent performance of these steels in

bending, sheet cutting and hole expansion operations. In this contribution, the impact of the Q&P process on the transsformation behavior of a 0.2C-4.5Mn-1.5Si lean medium Mn steel is presented and discussed in detail. This includes

dilatometric experiments examining the influence of the Q&P process parameters on the phase transformations, light optical and scanning electron microstructural (LOM, SEM) investigations, as well as hardness measurements according

to Vickers. Moreover, the Q&P annealing approach was compared to the TRIP-assisted bainitic ferrite (TBF) process, in order to evaluate the influence of primary martensite (α’initial) being present in the Q&P samples at the onset of isother-

mal bainitic transformation (IBT) on the transformation kinetics. With increasing quench temperature (TQ) the amount

of α’’ significantly decreased, whereas the phase fraction of bainitic ferrite (αB) and RA continuously increased, leading to the formation of fresh martensite (α’ final) and the sharp decrease of the RA fraction at TQ exceeding 250°C. Furthermore, with increasing partitioning temperature (TP) a decreasing amount of α B, associated with an increasing amount of α’final, could be examined. The comparison of the Q&P and TBF process in terms of transformation behavior manifested a pronounced influence of the presence of α’initial, since the IBT was intensively accelerated in the Q&P samples.

KEYWORDS: MEDIUM MN STEEL - Q&P PROCESS - PHASE TRANSFORMATION BAINITE - RETAINED AUSTENITE

Introduction High demands on crash safety and environmental protection are driving factors for the application of new steel grades in the automotive industry. (1) Stringent requirements in terms of reducing CO2 emissions, improving crashworthiness and enhancing the formability of complex shaped parts are impelling motives for the steel industry to the development of Advanced High Strength Steels (AHSS)

Simone Kaar, Reinhold Schneider University of Applied Sciences Upper Austria, Austria

Daniel Krizan voestalpine Stahl GmbH, Austria

Coline Béal, Christof Sommitsch Graz University of Technology, Austria

(2,3). By the application of AHSS grades, the sheet thick-

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Scientific papers - Heat Treatments ness can be downgauged due to their increased tensile

Rm*A80 exceeding 30,000 MPa% (18).

strength, resulting in a significant weight reduction of body structures (4). Furthermore, AHSS are characterized by an

The combination of the Q&P process and the concept of

improved ductility, which has beneficial impact on the pro-

medium Mn steels resulted in the development of lean me-

duction of complex automotive parts (5). The 1st genera-

dium Mn Q&P steels. (19) In comparison to conventional

tion AHSS includes Dual Phase (DP), Transformation-Indu-

Q&P steels, this steel grade contains an increased Mn-con-

ced-Plasticity (TRIP) and Complex Phase (CP) steels, with

tent (up to 5 wt-%), leading to a larger volume fraction of

a multiphase microstructure (6,7). These steel grades are

RA, which is chemically stabilized by C partitioning from

already in industrial application, since they are characteri-

α’ to γremain during isothermal holding in the over-ageing

zed by a good combination of strength and ductility with

region. The increased amount of RA and the presence of

tensile strengths (UTS) up to 1200 MPa and total elonga-

the hard C-depleted martensitic matrix allows for superior

tions (TE) up to 40 % (8,9). The 2nd generation AHSS, which

performance of these steels in both, deep drawing applica-

has outstanding mechanical properties with UTS*TE rea-

tions and forming operations such as bending, hole expan-

ching 60,000 MPa%, is represented by Twinning Induced

sion and sheet cutting (20).

Plasticity (TWIP), Nano-TWIP, Duplex and Triplex steels (10). However, until now, they are hardly used in industrial

By the application of the constrained carbon equilibrium

applications due to their challenging processing and high

(CCE) model proposed by Speer et al. (21), the amount of

alloying costs. (11)

RA as a function of TQ can be predicted. This simplified mo-

Currently, significant effort is put into the development of

del allows the calculation of the C-content in γ and therefo-

the 3rd generation AHSS in order to fill the property gap

re the prediction of the volume fraction of RA under three

between the 1st and 2nd generation AHSS (12). Therefore,

main assumption: (1) full C partitioning from α’initial to γremain,

recent research activities focus on the development of

while the partitioning kinetics are ignored; (2) no move-

Q&P, TBF and medium Mn steels (13,14).

ment of phase boundaries during the partitioning process; and (3) no competing reactions like αB formation or carbide

Q&P steels have a microstructure consisting of a car-

precipitation take place during the Q&P heat-treatment.

bon-depleted martensitic matrix with a substantial amount

(21)

of RA (15). Therefore, first the steel is fully austenitized, followed by subsequent quenching to a specific temperature

According to (22-27), the transformation from γ to αB du-

below the MS temperature in order to adjust the optimal

ring isothermal holding at TP has been observed in several

amounts of initial martensite (α’initial) and RA. Afterwards, the

instances.

steel is reheated to the over-ageing region, which triggers

Clarke et al. (25) and Kaar et al. (26,27) investigated a de-

the α’ tempering. During this stage the carbide precipita-

clining volume fraction of RA due to the formation of car-

tion will be significantly suppressed by the addition of Si

bide-free bainite during the C-partitioning step. However,

and/or Al. This enables the carbon partitioning from α’ into

besides the impact of the bainitic transformation on the

the remaining austenite (γremain), resulting in its appropriate

amount of RA, an influence on the RA stability was also

stabilization to room temperature (RT) upon final cooling

found (22-24).

(16,17).

This is due to the fact that the formation of carbide-free bainite contributed to the chemical stabilization of RA.

Medium Mn steels, which refer to the 3rd generation AHSS, as well, typically contain 3 – 12 wt-% Mn (12). They are

Therefore, the present contribution focusses on the investi-

characterized by a microstructure consisting of an ultrafi-

gation of the transformation behavior of a 0.2C-4.5Mn-1.5Si

ne-grained ferritic matrix with typical grain sizes below 1

steel grade. The influence of TQ and TP on the transforma-

µm and approximately 30 vol.-% RA (13). Medium Mn ste-

tion kinetics and thus the final microstructure and resulting

els have excellent mechanical properties with products of

hardness was investigated in detail. Furthermore, the Q&P

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Memorie scientifiche - Trattamenti Termici heat-treatment was compared to the TBF process in order

ingot of approximately 80 kg. The material was hot rolled to

to evaluate the influence of the presence of α’initial formed

a thickness of 4 mm, followed by tempering in a batch-an-

during quenching to TQ on the IBT.

nealing-like furnace at 550°C for 16h in order to provide cold rollability. Subsequently, the strip was cold rolled to a

Experimental Procedure

thickness of 1 mm. In order to investigate the transforma-

In the present investigation, a steel grade containing 0.173

tion behavior depending on TQ and TP (TB), a Bähr 805 A/D

wt-% C, 4.46 wt-% Mn, 1.47 wt-% Si and 0.027 wt-% Al was

dilatometer was used. Therefore, specimens with dimen-

investigated (Table 1). The steel was laboratory fabricated

sions of 10 x 4 x 1 mm³ were produced by wire-electrical

by medium frequency induction melting and cast into an

discharge machining.

Tab.2 - Chemical composition of the investigated steel in wt-%

Fe-C-Mn-Si

C

Mn

Si

Al

0.173

4.46

1.47

0.027

Figure 1 displays the time-temperature regimes for the

50 K/s.

Q&P (left) and the TBF (right) processes. In both cases, first the samples were fully austenitized at 900°C (Tan) for

Microstructural investigations were performed by means

120 s (tan) using a heating rate HR1 of 10 K/s. Subsequently,

of LOM using LePera etching according to (28). Additio-

for the Q&P heat-treatment the samples were quenched

nally, after electrochemical polishing, SEM was applied

to various TQ in the range of 130°C - 290°C in steps of 20°C

for higher resolution examinations.

with a cooling rate CR1 of 50 K/s. After 10s isothermal hol-

Mechanical properties were investigated by hardness te-

ding the samples were reheated with a heating rate HR2

sting according to Vickers (HV1) using an Emco Test Du-

of 20 K/s to a specific TP (350, 400 and 450°C) for 600s (tp).

raScan 20 device. SMM (29) was used to determine the

On the contrary, for the TBF cycle after full austenitiza-

volume fraction of RA, which was compared to the RA

tion, the samples were directly cooled to a TB of 350, 400

content calculated according to the CCE-model propo-

or 450°C and isothermally held for 600s (tB). Finally, all

sed by Speer et al. (21).

samples were quenched to RT using a cooling rate CR2 of

a)

b)

Fig.1 - Time-temperature regimes used for a) the Q&P and b) the TBF heat-treatments (26)

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Scientific papers - Heat Treatments By the application of the Koistinen-Marburger (KM) equa-

MS-equations are applied twice. Initially, the MS tempera-

tion (30), the volume fraction of α’ formed during quen-

ture of the γ and the volume fraction of α’ formed during

ching to TQ was determined

quenching to a specific TQ (orange line) are calculated.

fM = 1 - e^(-0.011(MS - TQ)) [1]

The remaining amount of γ after quenching is represented by the green curve in Figure 2 a).

Here, fM describes the amount of α’ formed during quen-

Since the applied CCE-model assumes full C partitioning

ching from the γ-region to TQ. MS is the martensite start

from α’ to remaining γ during isothermal holding at TP, a

temperature, which was calculated according to Mahieu et

new MS temperature is calculated according to Mahieu.

al. (31):

Given that, the volume fraction of α’ formed during final cooling to RT is calculated (blue line), using the KM equa-

MS = 539 - 423C - 30.4Mn - 7.5Si + 30Al

[2]

tion once again.

The resulting RA content as a function of TQ is shown in

Figure 2 a) shows the volume fractions of α’ and γ which

dark green. It is evident that the largest fraction of γ can be

were calculated by applying the CCE-model.

retained at that TQ where no fresh α’ is formed upon final

Figure 2 b) depicts the related time-temperature regime

cooling.

for the Q&P process. In the CCE-model both, KM- and

Fig.2 - a) Calculated phase fractions of martensite (α’) and austenite (γ) as a function of TQ and b) the related Q&P heat-treatment with the schematic microstructural evolution

Results

phase fraction could be found, since with increasing TP a smaller length expansion was observed. As a result, TP in-

Transformation Behavior

directly influenced the formation of α’final during cooling to

Figure 3 a) exemplarily shows the influence of TP on the di-

RT: Only for the Q&P samples with a TP of 350°C the forma-

latometric curves for the Q&P samples quenched to a TQ

tion of α’final could be fully prevented, whereas at higher TP

of 270°C compared to the influence of TB for the TBF sam-

of 400 and 450°C the curves confirm the formation of α’final

ples (Figure 3 b)). In the case of the Q&P regime, α’initial was

at low temperatures.

formed during quenching to TQ, since the MS temperatu-

On the contrary, for the TBF samples hardly any formation

re was approximately 310°C. During isothermal holding at

of αB during isothermal holding at TB could be observed.

TP, γ partially transformed to αB, accompanied by a length

The main transformation was the martensitic one during

expansion apparent from the dilatometer curves. Although

final cooling, indicating that the holding duration of 600 s

the amount of αB was rather low, an influence of TP on its

was definitely too short for bainitizing.

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Memorie scientifiche - Trattamenti Termici

Fig.3 - Dilatometric curves at different TP (TB) of a) Q&P heat-treatment (TQ =270°C) and b) TBF heat-treatment

Figure 4 gives a detail view of the influence of TP and TB on

400°C the formation of αB was not completed during the

the αB formation as a function of isothermal holding time

investigated period of time due to the slower transforma-

for the Q&P (TQ = 270°C) and TBF heat-treatment.

tion kinetics at lower temperatures.

In both cases, with increasing TP and TB the amount of αB

When comparing the two processes, it is evident that in

formed during isothermal holding decreased.

general for the Q&P samples the larger volume fraction of

However, especially for a TP of 350°C and TB of 350 and

αB was formed.

Fig.4 - Dilatation due to αB formation as a function of isothermal holding time at different TP (TB) for a) Q&P heat-treatment (TQ = 270°C) and b) TBF heat-treatment

In addition, for the Q&P steels the transformation beha-

isothermal holding increased with increasing TQ, which is

vior depending on TQ was investigated (Figure 5).

confirmed by the larger dilatation in Figure 5 b).

With increasing TQ, lower amounts of α’initial were formed

Furthermore, as already displayed in Figure 3 a) at very

due to the lower driving force for the martensite transfor-

high TQ of 270°C, the formation of α’final during final cooling

mation.

to RT was observed.

On the contrary, the volume fraction of αB formed during

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Scientific papers - Heat Treatments

Fig.5 - a) Dilatometric curves at different TQ (TP = 400°C) and b) dilatation due to αB formation as a function of isothermal holding time at different TQ (TP = 400°C)

Microstructure

mixture of α’’ and small amounts of αB, appearing as bluish and brownish areas in the micrographs. Furthermore, RA

As already confirmed by means of dilatometry, TQ signi-

and/or α’final , represented by the white and brownish are-

ficantly influenced the phase transformation behavior,

as, could be observed in rising amounts finely distributed

and thus, the presence and the amount of the final micro-

in the matrix.

structural constituents of the investigated steel.

As confirmed by dilatometry, at a TQ of 190 and 230°C,

Figure 6 a) - c) depicts the LOM images of the Q&P sam-

the secondary phase solely consisted of RA, whereas

ples quenched to a TQ of 190, 230 and 270°C, respectively.

for the samples quenched to higher TQ the formation of

In comparison, the microstructure of the TBF sample is

α’final could not be prevented. On the contrary, the micro-

shown in Figure 6 d). All micrographs relate to the the mi-

structure of the TBF samples primarily consisted of α’final

crostructure obtained at a TP and TB of 400°C. The LOM

with small amounts of finely distributed RA islands and

images for the Q&P steels show a matrix consisting of a

only small fractions of αB.

Fig.6 - LOM images of the a) - c) Q&P samples (TQ = 190, 230 and 270°C) and d) TBF sample at a TP (TB) of 400°C (original magnification 1000x)

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Memorie scientifiche - Trattamenti Termici The SEM images of the Q&P samples quenched to a TQ of

independent from TQ. However, at a TQ of 270°C a sub-

190, 230 and 270°C are displayed in Figure 7 a) – c), where-

stantial amount of α’final was found in the microstructure.

as Figure 7 d) depicts the microstructure for the TBF sam-

In the present micrographs, RA appears as cavities, since

ple. By means of SEM, small amounts of cementite preci-

it was dissolved during electrochemical polishing due to

pitates with an obvious triaxial alignment, confirming the

its high C and Mn content.

presence of α’’, could be observed for the Q&P samples.

In comparison, the TBF treated sample largely consisted

Concerning the presence of αB, no significant influence of

of fine lath like martensite (α’final) with low fractions of RA

TQ was found, since the volume fraction was rather low,

and αB.

Fig.7 - SEM images of the a) - c) Q&P samples (TQ = 190, 230 and 270°C) and d) TBF sample at a TP (TB) of 400°C (original magnification 5000x)

In Figure 8, the phase fractions in the final microstructure

oling, in turn resulting in comparably low amounts of RA.

of the Q&P samples are summarized as a function of TQ for different TP. In addition, the bar charts, situated close to

In Figure 9, the amount of RA measured by means of SMM

the right rim of each diagram, depict the microstructural

is plotted in detail as a function of TQ at a TP of 350, 400 and

constitution for the steel after TBF treatment. Indepen-

450°C. Furthermore, the volume fraction of RA for the TBF

dent from T , for the Q&P samples with increasing TQ a

treated samples is shown in the form of dotted lines. In

considerable decrease of α’’ was observed, whereas the

addition, the diagram contains the calculated RA fraction

volume fraction of αB increased. However, particularly at

according to the CCE-model. This model proposed a RA

higher T , a vivid influence of TP is obvious: the increase

maximum (RAmax) of 19.8 vol.-% at a TQ of 173°C. In accor-

of TP led to a sharply decreasing fraction of αB, resulting in

dance with the model calculations, the experiments also

the formation of α’final upon cooling to RT. Therefore, the

confirmed the fundamental shape of the RA evolution as a

RA content steadily rose with increasing TQ until the on-

function of TQ. At very low TQ large amounts of α’initial were

set of the formation of α’final occured. For the TBF samples,

formed, leading to low RA fractions. With increasing TQ

it is obvious that only small amounts of αB were formed

the amount of α’initial decreased, and therefore the volu-

at all investigated temperatures as already shown by di-

me fraction of RA increased, until RAmax was achieved, fol-

latometry (Figure 4 b). Due to these low αB fractions, the

lowed by a decline in RA content. This decreasing amount

largest amount of γ transformed into α’final during final co-

of RA at higher TQ was due to the lower chemical stabiliza-

P

Q

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Scientific papers - Heat Treatments tion of austenite and the resulting formation of α’final upon

achieved compared to the model predictions. Moreover,

cooling to RT. The experiments further show an influence

RAmax occurred for all TP, at temperatures 60 to 70°C hi-

of TP on the volume fraction of RA: by increasing TP from

gher than predicted by the CCE-model.

350 to 400°C, RAmax rose from 19.7 to 22.4 vol.-%.

It is evident that the RA contents determined for the TBF

In this context, at the TP of 350°C, the amount of RA cor-

samples were much lower compared to the Q&P samples,

related well with the model calculations. On the contrary,

ranging from 4.6 vol.-% RA at a TB of 450°C to 7.1 vol.-%

at the TP of 400°C and 450°C, a higher amount of RA was

RA at TB = 350°C.

Fig.8 - Phase fraction as a function of TQ for a TP (TB) of a) 350°C, b) 400°C and c) 450°C

Fig.9 - RA content as a function of TQ for a TP (TB) of 350, 400 and 450°C

Fig.10 - Vickers hardness as a function of TQ for a TP of 350, 400 and 450°C

Hardness

ness could be observed, especially at a TP of 450°C. For this

Figure 10 depicts the hardness according to Vickers as a

reason, in the case of the Q&P samples, the hardness was

function of TQ. In the TQ range of 130 to 230°C, both increa-

between approximately 400 and 500 HV1. In contrast, for

sing TQ and TP led to a slight decrease in hardness. However,

the TBF samples hardness was quite constant with 510 HV1

when further increasing TQ, a considerable increase in hard-

at a TB of 400°C, and 511 HV1 at 350 and 450°C, respectively.

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Memorie scientifiche - Trattamenti Termici Discussion

Al leads to a larger amount of αB formed during isother-

In Figure 11, the IBT for the Q&P samples (exemplarily at

mal holding at TP. According to (33) this can be explained

a TQ of 270°C) and the TBF samples is compared. The TTT

by the fact that Al shifts the T0-line to higher C-contents,

diagrams show an evident acceleration of the αB formation

enabling the formation of a larger amount of αB.

in case of the Q&P process. This is in agreement with Kaar

Furthermore, in the present case, a general impact of the

et al. (26) and Wang et al. (32), since the presence of α’initial

TP (TB) on the transformation kinetics was found: espe-

pronouncedly accelerated the IBT due to the presence of

cially at a TP (TB) of 400°C, the transformation kinetics was

geometrically necessary dislocations, acting as nucleation

evidently accelerated due to the faster C diffusion into the

sites for the αB formation. When comparing the IBT of the

remaining γ, while at 450°C a reduced driving force for the

present steel grade with that of a Fe-C-Mn-Al steel inve-

bainitic transformation seems to limit the transformation

stigated in (26), it is evident that the substitution of Si by

kinetics.

Fig.11 - TTT diagrams for a) Q&P steels (TQ = 270°C) and b) TBF steels

Furthermore, in terms of RA investigations, a considerable

αB formed during isothermal holding. This indicated that in

difference between the calculated CCE-model and expe-

the present case, RA was mainly stabilized via C-partitioning

rimental data was found. This was particularly apparent for

during the partitioning step, whereas for the Al-alloyed ste-

the amount of RA and optimal TQ, which is in agreement

el investigated in (26, 27) the RA stabilization was achieved

with results reported in literature (21,26,27,34). One reason

by both, formation of αB and C-partitioning.

for this discrepancy is the fact that the application of the

In general, in terms of RA stability, both, chemical and me-

CCE-model assumes full C-partitioning from α’initial into the

chanical stabilization have to be considered (35). However,

remaining γ. However, in the present case, the presence of

the CCE-model only takes into account the chemical stabi-

cementite was observed in the microstructure (Figure 7),

lization, which explains the significant shift of the RAmax to

which contradicts this assumption. This led to both, a lower

higher TQ than predicted. Additionally, two empirical for-

amount and chemical stability of γremain, resulting in the for-

mulas, describing the α’ kinetics and the MS temperature,

mation of α’final during cooling to RT and thus lower RA con-

are included in the CCE-model. Therefore, slight differen-

tents. Furthermore, the CCE-model does not consider the

ces in the chemical composition due to segregations could

formation of αB, associated with a decline in RA content. In

lead to deviations from the model, as well (36). Neverthe-

comparison to the Fe-C-Mn-Al steel grade investigated in

less, the CCE-model proposed by Speer et al. (21) is a vital

(26,27), the RA content was higher for the Si-alloyed steel

tool for the first estimation of applicable annealing parame-

examined in this contribution due to the lower amount of

ters in case of Q&P steels.

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Scientific papers - Heat Treatments Conclusions

• The increase in T Q resulted in a lower volume fraction of α’ initial and slightly larger amounts of αB. Therefore, the vo-

The results of the present investigation indicate that the

lume fraction of RA increased, until that TQ where α’final

Q&P process has significant influence on the transforma-

was formed and thus RA content declined sharply. In case

tion behavior of lean medium Mn steels.

of the Q&P samples this led to a final microstructure consisting of α’’, RA and at higher TQ partially of αB and α’final.

It is of vital importance to carefully set both, T Q and T P, since they have significant impact on the microstructural

• Especially for the Q&P samples heat-treated at a low TP,

constituents and thus, certainly mechanical properties.

the comparison of the CCE-model and the experimental data showed a significant divergence in terms of RA con-

The main findings of the present contribution can be

tent. Additionally, a substantial shift of the RA-maximum

summarized as follows:

to higher TQ was observed compared to the CCE-model predictions.

• In case of the Q&P process, the formation of α’initial during quenching to TQ significantly accelerated the IBT,

• In the present case, for the Si-alloyed steel grade RA

acting in the form of vital nucleation sites.

was mainly stabilized by C-partitioning during isothermal holding at TP, whereas on the contrary for the Al-alloyed

• For both, Q&P and TBF processes, with increasing TP

steel investigated in (26, 27), the RA stabilization was per-

(TB) the IBT kinetics was accelerated, but the amount of αB

formed by both, formation of αB and simultaneous C-par-

formed during isothermal holding was decreased.

titioning.

Acknowledgements The authors sincerely acknowledge the support of the Austrian Research Promotion Agency (FFG) related to the frontrunner project No. 860188 “Upscaling of medium Mn-TRIP steels”.

La Metallurgia Italiana - January 2020

pagina 15

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Memorie scientifiche - Trattamenti Termici

High pressure gas quenching: assessment of velocity experimental measurements and steps for model validation E. Vyazmina, L. Bustamante-Valencia, V. Woimbee, G. Michel, A. Belhajria, Ch. David

Gas quenching is a key step in the process of heat treatment of metals. To reach optimizations of the process time and quality, it is essential to correctly distribute the load inside the quenching cell. Numerical simulations using CFD (Computational Fluid Dynamics) could help us to understand the phenomena occurring into the quenching cell and identifying criteria for improving the process. However, numerical simulation must be validated versus experimental data.

This paper presents the experimental and simulation works conducted in order to validate the commercial software Qobeo for the modeling of gas quenching. The computational domain has been refined, the mesh corresponds precisely to the geometries of the quenching cell and the rack. The gas flow for quenching is simulated in the atmospheric conditions for pressure and temperature. The numerical results are compared with experimental measurements of flow speed. The results obtained are very encouraging, coherent behaviours are observed in both cases. However, further improvements should be conducted for better approaching the actual quenching conditions. For modelling aspects, high pressure (up to 20 bar), high temperature (up to 900 °C) and various gases (N2, CO2, He) should be included. For experimental aspects, the uncertainty of gas flow measurement technique should be determined as well as improving the assessment of flow orientation.

KEYWORDS: GAS QUENCHING - CFD - SPEED MEASUREMENTS GAS FLOW - HEAT TREATMENT - HPGQ Context French Institute of Technology for Material, Metallurgy and Process (IRT M2P) In France, 8 French Institutes of Technology (IRT) have been created in the 2010 decade. The common target of such Institutes is allowing transfer of knowledge and development between academics and industrials in specific fields. The IRT’s sites are smartly distributed in French regions, close to the industrial activity sources as shown in Fig 1.

The works presented here have been conducted by the French Institute of Technology for Material, Metallurgy and Process (IRT M2P). The IRT M2P is a partnership built up in 2013 in order to accelerate innovation, allow integration of

La Metallurgia Italiana - gennaio 2020

E. Vyazmina Air Liquide R&D, France

elena.vyazmina@airliquide.com

L. Bustamante-Valencia Air Liquide, AL France Industrie, France

lucas.bustamante-valencia@airliquide.com corresponding author

V. Woimbee Faurecia, France

valerie.woimbee@faurecia.com

G. Michel IRT M2P, France

gregory.michel@irt-m2p.fr

A. Belhajria, Ch. David Sciences Computers Consultants, France, abelhajria@scconsultants.com

pagina 18

Scientific papers - Heat Treatments new technologies and to share a technologic platform of in-

M2P with shared resources between industrials. It gathers

dustrial scale for tests.

contributions of: Air Liquide (gas industry), Faurecia (au-

The Institute is located in the east of France, close to an im-

tomotive industry), IRT M2P (supplier of the technical pla-

portant network of heat-treatment and materials industries.

tform), SCC consultants (innovative company offering its

Also, universities and technical centres of excellence are in

knowledge in fluid simulation and computer systems), ECM

the surroundings.

furnace producer.

This paper illustrates the mutual works conducted by IRT

Fig.1 - Location of French Institute of Technology. The IRT M2P is located in the east of France

Introduction

not need any post washing, the homogeneity of heat tre-

Potentialities offered by gas quenching are huge, if the

atment of the load is enhanced and the intensity of quen-

process is correctly mastered. The present work is focu-

ching can be controlled precisely, but also, the distortion

sed on optimization of gas quenching, first experimentally

of parts is dramatically reduced compared to traditional

and secondly by developing sustainable models able to

technology. Nevertheless, the parts are concentrated in a

forecast cooling rate.

small volume into the quenching cell and the technology is not as mature as oil quenching. Large potentialities of

Gas quenching and oil quenching

development exists for gas quenching, this work aims to

A key project has been launched in IRT M2P to improve

provide knowledge in this field (2).

our knowledge in advanced heat-treatments, especially quenching (1). In industries, oil quenching is largely spre-

The gas quenching technology

ad but exhibits several disadvantages: absolute necessity

During gas quenching, the parts are cooled from the treat-

to wash parts after treatment, recyclability of oil after use,

ment temperature, often approximately 900 °C, until room

detergent rejection and compliancy with ecological regu-

temperature in a given time dependent on the quenching

lations more stringent. Moreover, flexibility is weak when

intensity required. In this cell, the gas is injected at high

oil containers overcome a volume of 1 m and temperature

pressure (20 bar max), then a turbine located on the top

should be regulated precisely at different set points. The

blows the gas towards the load. The gas is heated up when

key advantages of gas quenching are: the metallic parts do

enters into contact with the hot parts. The gas is cooled

3

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pagina 19

Memorie scientifiche - Trattamenti Termici down when passes throw a heat exchanger located in the

The main difficulties are, in one hand, the prediction of gas

wall of the cell, then directed toward the turbine again. The

circulation when the metallic parts are assembled in a rack,

gas flow in the cell is shown in Fig. 2. For certain confi-

and on the other hand, the calculation of the actual heat

gurations of load, in particular the massive ones, an unidi-

exchange coefficient and the change of metallographic

rectional gas flow showed to be insufficient to guarantee

properties of metal when cooling. Also, in gas quenching

the uniformity of treatment. In order to address this issue,

parts hidden by obstacles could receive less gas flow, le-

quenching cells equipped with two turbines were design,

ading to a higher dispersion level for a given cooling rate.

in general located on both lateral sides of the cell. Some

The composition of the atmosphere used for quenching

other manufacturers design alternative flow and rotating

has a critical influence in the final characteristics of the par-

deflector to orientate the flow in all the directions (3).

ts. It has been shown that an atmosphere composed of 90 %nitrogen and 10 % helium can provide a profile of tem-

The prediction of distortion of parts submitted to heat tre-

perature decrease similar to the one of oil (4). Changing

atment and quenching is tricky whatever the quenching

the gas or the proportions of a mix, can lead to improve-

media is, oil, water polymer mixture or gas. In the case of

ments of quenching performance, so conducting tests

gas, such prediction requires a correct prediction of the

with a mixing of helium and another gases could present

heat transfer by convection from the solid towards the gas.

interesting opportunities.

Fig.2 -Industrial gas quenching facility, gas flow follows the direction of blue arrows

Quenching experimental device Furnace and quenching cell In this research, experimental tests of gas quenching

The gas quenching cell is on the front, it allows also the

have been conducted in the industrial furnace of IRT M2P

entrance and exit of the load.

shown in Fig. 3.

The pressure of gas for quenching is set to 20 bar, which is

It is divided into two chambers: the first for low pressure

the maximum pressure admissible for the chamber.

carbo-nitriding and the second for gas quenching.

The higher pressure, the faster quenching; the heat

The parts for heat treatment are ranged in a rack of around

exchange capacity is optimum.

250 kg. After the high temperature stage in the cell at the

The gas is sprayed from top to the bottom of the parts rack

back of the furnace, they are transferred automatically to

(as presented in Fig. 2).

the quenching cell.

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Scientific papers - Heat Treatments

Fig.3 - Furnace on which test have been conducted and rack of parts

Gas flow measurements

Observations have been set by a transparent Plexiglas pla-

Experiments have been launched to get references to va-

te in place of the door, allowing to track smoke circulation.

lidate the gas flow simulation. The first trials (2) have been

A high speed camera ensures the shooting of the smoke

carried out to know qualitatively the gas flow circulation

transfer. Particular pictures are extracted from the film for

and assess visually the turbulence level. The furnace para-

expertize, see Fig. 4.

meters are set at ambient pressure and temperature.

Fig.4 - Left, quenching cell and installation for high speed camera. Right, smoke flow visualization

Fig. 5 shows honeycomb grid. On the left hand side, the

have been positioned. On the right had side, the simula-

real grid indicating the position of the gas flow meters that

tion’s honeycomb grid.

Fig.5 - Left, real honeycomb grid with the position of gas flow meters. Right, simulated honeycomb

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Memorie scientifiche - Trattamenti Termici Gas quenching simulation

into account by the estate law.

• Modelling of turbulence with k-epsilon or other models.

Quenching simulation tool qobeo®

Heat transfer:

The numerical simulation of the gas quenching has been

• Solving of heat transfer equation.

conducted with qobeo® which is a commercial software

• Modelling of radiative phenomena.

of 3D simulation for furnaces and installations for indu-

• For liquids, calculation of boiling effects.

strial quenching. This software has been developed by

Calculation of mechanics-metallographic effects:

the reseach consortium ThosT gathering heat treatment

• Possibility of coupling results from qobeo® with FOR-

industries, the Center of Forming of Materials (Cemef of

GE® NxT (or other software).

MinesParisTech, France) and the company SCC which is

• An internal calculation model of the mechanics-metal-

responsible for the integration and commercial activities

lographic effects is ongoing in the frame of the INFINITY

(5,7).

industrial chair (9).

qobeo® is a software design to be easily used by practitio-

Regarding the complexity of modelling, in this work, a sy-

ners of heat treatment. For being used, it does not require

stematic validation of the numerical results with regards to

deep knowledge in heat transfer, fluid mechanics or com-

corresponding experiments is conducted.

putational matematics. Pre-set scenarios help the user to build up his own simulation.

Simulation domain setup

The mesh is build up with tetraedric elements (8), for gas

The simulation domain and geometry used in CFD mode-

phase simulation the immersed boundary technique is

ling is presented in the Fig. 6, left. The flow is injected into

used and for solid phase calculation, the method of fini-

the quenching cell, then blown with a turbine. The asso-

te elements is used. It allows solving a single equations

ciated flow rate is 4 m3/s, as given by the furnace manu-

system with thermo physical properties adapted to the

facturer. This flow rate corresponds to a flow velocity of

sub-domain (solid or fluid). qobeo® solves equations in

18-20 m/s. Due to the geometrical shape of the turbine the

transitory state in three dimensions for Newtonian fluids,

injected flow has an axial and azimuthal components.

as described here after.

In the current simulations, a full 3D approach was chosen.

Fluid mechanics:

A tetrahedral mesh of 12 million cells was chosen for the

• Resolution of the coupling speed-pressure using the

simulations, as shown in Fig. 6, right. The mesh was refi-

non compressible Navier-Stokes equations.

ned in the zone of the turbine, the flow deflector and on

• For high pressure gases, the compressibility is taken

the honeycomb grids.

Fig.6 - Left, CFD simulation domain. Right, mesh description

La Metallurgia Italiana - gennaio 2020

pagina 22

Scientific papers - Heat Treatments In Fig 6 left, the outlet boundary conditions are shown in

CFD simulation results

blue. These are open boundaries allowing the flow to leave the computational domain, it avoids any recirculation

Fig. 7 shows the main flow trajectories. The gas injected

close to the outlet. This velocity is computed and in the

via the turbine circulates inside the cell, passes through

current simulation it is approximately 150 m/s at the outlet

the honeycomb grid and goes out of the computational

of the turbine. The extra space in the cell close to the door

domain through lateral outlets.

is also included into the simulation domain. Its presence

There are recirculation regions on the left and right of the

has a strong impact on the flow due to the symmetry brea-

honeycomb grid. In industrial situations, this region corre-

king in the Y direction (rear-front of the cell).

sponds to the position of the metallic load which is quenched.

The Fig. 5 shows the grid in which the metallic parts are

Here, the recirculation points creates enhanced turbulen-

located. The load submitted to quenching consists of 4

ce zone that could help to cool down even faster the sur-

plateaus of honeycomb shape. There is a small difference

rounding pieces.

between the real geometry (on the left) and the model-

The presence of this recirculation is also found experi-

led one (on the right): there are 4 holes for real geometry,

mentally, see Fig. 4 (right frame). In the experiments, the

whereas in simulations 5 holes were imposed. We assume

smoke accumulates in the recirculation region predicted

that this difference does not induce critical error in the si-

by simulations.

mulation results.

This qualitative comparison demonstrates that qobeoÂŽ correctly reproduces the physics of the flow inside the

An unsteady approach is used. Due to very small time step

cell.

the simulations took 3 weeks using 16 processors.

Fig.7 - Flow trajectory

Fig. 8 shows the average velocity field inside the cell. The

of X axis, mainly top to down flow. The blue arrows evi-

flow passes the turbine located at the top of the cell and

dence flow in the opposite direction. Fig. 8 shows clearly

passes through the grid before being evacuated throu-

that the numerical obstacles disturbs the gas flow field, as

ghout the lateral outlets. The presence of the honeycomb

evidenced experimentally.

grid creates multiple recirculation points. The red arrows indicate the resulting flow speed in the positive direction

La Metallurgia Italiana - January 2020

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Memorie scientifiche - Trattamenti Termici

Fig.8 - Flow field. Red arrow top towards bottom flow. Blue arrow, inverse flow

Fig. 9 presents the magnitude of the average velocity at

mogeneity between levels 4 and 1 could have a critical

the four levels of the honeycomb grid. At the top, on level

influence with regards to the quenching intensity of the

4, the velocity is higher than the ones observed in lower

metallic parts located in both grids. Assessing the effect

levels. The flow is inhomogeneous in the horizontal sli-

of such flow difference on the final mechanical properties

ce, this is due to the vorticity induced by the deflector.

of the metallic parts should be done in future works.

After passing through level 3 and 2 of the grid, the velocity distributions becomes more and more homogeneous

Analysing the global computational domain, the results

and the flow tend to be laminar. The velocity is homo-

of Fig. 9 and Fig. 8 show clearly that the flow is not sym-

geneously distributed at level 1, but the speed has been

metric in the volume of the quenching cell. Also, the ve-

divided by a factor of 10 with regards to the one of level 4.

locities become more homogeneous on the door side

The difference observed in terms of speed and flow ho-

and closer to the bottom of the cell (near to level 1).

Fig.9 - Magnitude of the average velocity at the four honeycomb grid levels

Experimental measurement (as presented in Fig. 5) and

1. Calculation of the average velocity fields in the time in-

numerical measurement points were located on level 4

terval from 15 s to 25 s.

with 17 measurement points and in levels 3, 2 and 1 with

2. Computation of the average speeds in a volume encom-

4 measurement points per level. The positioning of such

passing the measuring point. The volume corresponds to

detectors is detailed in Fig. 10. For the comparison of data

a cylinder of 2 cm diameter and 2 cm length (see Fig. 10,

from experiments and simulations, the following proce-

right).

dure was followed:

La Metallurgia Italiana - gennaio 2020

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Scientific papers - Heat Treatments

Fig.10 - Left and centre, position of instrumentation for the measurements of flow speed. Right, measurement point

Fig. 11 and 12 present the comparison of simulation re-

• The gas speed is higher in the rear and the front of the

sults (continuous lines) and experimental measurements

rack in which no physical barriers are stablished. Never-

(dots) at different levels of the honeycomb grid. Simula-

theless, these volumes does not play a role in the cooling

ted velocities are of the same order of magnitude as the

of the metallic parts.

experimental measurements, which are very encouraging results. Some coherent results should be highlighted:

• The turbulence induced by high speed gas avoid creating “preferential” path in which the flow is laminar and

• In the center of the honeycomb grid the speed is lower

established in steady state. This is the cause of the lack

that in the sides. This is due to the physical obstacles pre-

of linearity increasing or decreasing with regards to the

sent in the middle of the grid as shown in Fig 5, left. But

speeds of Fig. 11.

also, due to the fact that the momentum of the gas provided by the turbine and the deflector is lower in the center. This effect is clearly visible in the top view of Fig 7: the farther away from the center, the faster flow.

Fig.11 - Comparison of experimental and simulated velocity at level 4. Dot: experiment, line: simulation

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Memorie scientifiche - Trattamenti Termici

Fig.12 - Comparison of experimental and simulated velocity at level 1, 2 and 3. Dot: experiment, line: simulation

As shown, very good experimental and numerical results

experimental result can be an averaged value over a pe-

have been obtained.

riod of time (mean value), an instantaneous value whate-

Nevertheless, for getting a more accurate comparison

ver the value is, or the value obtained in steady state (if it

and improving our analysis, it is necessary to have the

is reached).

following information on the experimental data:

The key question is then, what is the standard deviation

• Calibration procedure and uncertainty of the gas flow

for the averaged gas velocity?

detectors.

• The error of the experiment produced by the repeatabi-

All this information is essential to find the error bar in the

lity of the experiment.

experiments. Without the error bar on the measuremen-

• The exact turbine flow rate and its variation in time.

ts, it is impossible to judge the quality of the experimental results. Hence, it is very complicated to perform an

As described, the turbulence in the chamber can indu-

accurate comparison with simulations.

ce a vibration in the flow speed measurement. Then, the

Conclusion and future works

experimental ones for speed distributions at the 4 levels of the honeycomb grids of the cell. Nevertheless, the

The gas quenching technology offers interesting oppor-

accuracy of experimental data used for the comparison

tunities for massive heat treatment of mechanical parts

should be improved, the uncertainty of the speed resul-

for automotive industry. This technique presents many

ts should be taken into account. Hence, improvement of

advantages such as reduction the distortion of the load

the experimental campaign is needed for more reliable

and close controlling of quenching intensity.

comparison with the simulations. It would be very beneficial to perform PIV measurements.

In order to optimise the gas quenching process, a comparison of experimental and simulation results has been

The numerical results showed important differences in

conducted. Trials have been carried out in the industrial

terms of speed and flow homogeneity between levels 4

plateform of IRT M2P. Modeling of a quenching cell have

and 1 of the rack. This could have a critical influence with

been performed using the 3D CFD software qobeo®. The

regards to the quenching intensity of the metallic parts

simulations evidenced a good agreement with measu-

located in both grids, the final properties of the parts tre-

rements qualitatively (the same flow behavior inside the

ated could present discrepancies. Assessing the effect of

cell). Encouraging results have been obtained, the nu-

such flow difference on the mechanical properties of the

merical results are of the same order of magnitude as the

metallic parts should be done in future works.

La Metallurgia Italiana - gennaio 2020

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Scientific papers - Heat Treatments References [1]

Woimbee V. Gas quenching. Internal IRT M2P, 2017.

[2]

Woimbee V, Belhajria A. Gas quenching process: simulation can provide better understanding. IRT-M2P, in A3TS conference Bordeaux, 2018.

[3]

https://www.bmi-fours.com/products/gas-cooling-furnaces/b8-t-high-temperature.html

[4]

Bustamante Valencia L. Performance du refroidissement sous flux gazeux : choisir son gaz et les paramètres de mise en œuvre.

Séminaire Bodycote - Air Liquide, France, October 2016.

[5]

E. Hachem, E. Massoni, T. Coupez, H. Digonnet. 8th. Heat Transfer Modeling Inside Industrial Furnaces. World Congress on Computational Mechanics (WCCM8) 5th European Congress on Computational Methods en Applied Sciences ans Engineerint (ECCOMAS 2008) - Venice, Italy, June 30-July 5, 2008.

[6]

E. Hachem, T. Kloczko, H. Digonnet and T. Coupez. Stabilized finite element solution to handle complex heat and fluid flows in industrial furnace using the immersed volume method. International Journal for Numerical Methods in Fluids, 2010.

[7]

https://www.scconsultants.com

[8]

G. Jannoun, J. Veysset, T. Coupez, E. Hachem. Edge-based anisotropic mesh adaptation for CFD applications. Proceedings of the 21st International Meshing Roundtable, pages 567–583, 2013.

[9]

Industrial chair industrielle INFINITY (on going) - http://chaireinfinity.fr

La Metallurgia Italiana - January 2020

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Memorie scientifiche - Trattamenti Termici

Influence of gas nitriding conditions on layer structure formation on grey- and white-solidified cast irons (part I) A. Holst, A. Buchwalder, R. Zenker

The wear and corrosion behaviours of nitrided cast iron depend strongly on the material-specific morphology of its graphite. Furthermore, and depending on its specific form, this morphology can have either a positive or a negative effect on the potential for improvement of such behaviour. For this reason, comparative investigations were carried out to determine the influence of gas nitriding (GN) conditions on the layer structure and the properties formed for both grey-solidified (with graphite) and white-solidified (without graphite, but with ledeburitic carbides) cast iron surface layers generated by electron beam remelting (EBR). By varying the nitriding potential within the limits KN = 0.1 (no compound layers), KN = 1 and KN = 4 (with compound layers), as determined on the basis of the Lehrer diagram, the influence of different Si contents in the cast irons on the layer formation and phase composition were investigated. In addition, the influence of the morphology of the two different C-containing modifications - the stable graphite (lamellar, spheroidal) and the metastable ledeburitic carbides (secondary dendrite arm spacing: 3 µm - 7 µm) – on the formation of both the compound and diffusion layers was analysed by means of SEM investigations, hardness measurements and EBSD analysis. This allowed conclusions to be drawn about the microstructurally dependent nitriding mechanism during gas nitriding. For the metastable carbide layers, the subsequent thermal regime of the nitriding process applied (EBR+GN) may have influenced the initial microstructure. For this reason, the tempering behaviour was investigated in separate annealing tests under temperature-time regimes relevant for nitriding (540 °C to 670 °C / 4 h to 16 h). The hardness decrease induced by the coarsening of pearlitic cementite in the ledeburite was strongly dependent on the chemical composition of the cast iron used. At temperatures above 630 °C, the hardness dropped remarkably due to graphitizing. In light of these results, the nitriding parameters could be optimized.

KEYWORDS: CAST IRON - ELECTRON BEAM REMELTING - GAS NITRIDING DUPLEX SURFACE TREATMENT

Introduction Grey-solidified cast iron is interesting for numerous applications due to some of its special properties, such as good castability and associated near-net-shape com-

Anja Holst, Anja Buchwalder, Rolf Zenker

ponent manufacture, high damping capacity, high com-

TU Bergakademie Freiberg, Institute of Materials Engineering

pressive strength and emergency running properties

Freiberg, Germany

[1,2]. Alloying with silicon (Si) promotes solidification according to the stable Fe-C system. Therefore, the microstructure consists of coarse graphite particles em-

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Scientific papers - Heat Treatments bedded in a ferritic or pearlitic matrix. As known from

particular, the role of silicon in nitriding, hypoeutectic

steel, nitriding is a common method for improving both

pearlitic and ferritic cast irons were investigated in the

the material’s wear and corrosion resistance [3]. Howe-

initial (grey-solidified) and EB remelted (white-solidi-

ver, because the graphite is not nitrided, i.e., the layer is

fied) states.

interrupted, the compound layer is inhomogeneous and

Gas nitriding was carried out with different nitriding para-

does not always fulfil its function satisfactorily.

meters (time t N, nitriding potential K N) at a constant tem-

Electron beam remelting (EBR) of cast iron surfaces car-

perature T N of 540 °C.

ried out before the nitriding treatment leads to the dissolution of near-surface graphite and, as a result of the

Experimental

process-related high cooling rate, to solidification according to the metastable Fe-Fe 3C system.

Base materials

The eutectic microstructure of white-solidified cast irons

The investigations were carried out on commercially

is referred to as ledeburite, which exhibits high hardness

available grey-solidified pearlitic (P) and ferritic (F) cast

and has been investigated extensively [4–6].

irons with lamellar (GJL), spheroidal (GJS) and vermicular

The morphology of ledeburite depends on the cooling

graphite (GJV), with the samples provided in the form of

rate. The degree of microstructural refinement can be

cast flat samples (100 × 60 × 20 mm³) by the foundry Kes-

influenced by varying the energy input during EB remel-

sler and Co. GmbH (Leipzig, Germany). Tab. 1 shows the

ting.

respective chemical composition and the carbon equi-

The higher the cooling rate, i.e., the thinner the remel-

valent (CE) calculated from it. Compared to conventio-

ted surface layers, the finer the remelted microstructure.

nal GJL-250 cast iron, small additions of chromium (Cr),

During nitriding of the ledeburitic surface layers, a clo-

copper (Cu) and titanium (Ti) served to modify the for-

sed, homogeneous compound layer is formed [7]. Such

mation of the graphite lamellae in order to improve the

combined treated surface layers (remelting + nitriding)

melting behaviour during EB remelting [10]. The nomi-

are significantly more resistant to abrasive wear and cor-

nal tensile strength was not affected. This was marked in

rosion [8,9].

the material designation by the imposed character mod

In order to understand the nitriding mechanism and, in

(GJL-250 mod).

Tab.1 - Nominal composition (Fe base) of the investigated alloys and the corresponding carbon equivalents CE (CE = wt.%C + ⅓ (wt.% Si + wt.% P))

The alloys were grouped with respect to their matrix

tic cast irons were 258 HV0.3 (P/L_2.0Si) and 288 HV0.3

microstructure (P - pearlitic, F - ferritic), graphite mor-

(P/S_2.3Si), respectively, while that of the ferritic cast

phology (L - lamellar, S - spheroidal, V - vermicular) and

iron was 167 HV0.3 (F/S_2.3Si, F/V_2.5Si).

silicon content, cf. Tab. 1.

The measured mean core hardness values of the pearli-

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Memorie scientifiche - Trattamenti Termici Treatment parameters

were electrolytically coated with a protective Ni layer

The electron beam remelting (EBR) was carried out in

prior to embedding. To facilitate light-optical examina-

a K26-15/80 universal electron beam chamber facility

tion, the specimens were etched with 3 % Nital solution.

(pro-beam systems GmbH), which operated at a beam

The specimens investigated by scanning electron micro-

power ≤ 15 kW and an acceleration voltage ≤ 80 V. Be-

scopy were not etched, but subjected to an additional

fore remelting, the specimens were preheated in-situ

polishing step using Buehler MasterMet ™. A LEO 1530

up to a temperature of approx. 450 °C using the EB field

FE-SEM (Carl Zeiss) equipped with an HKL Channel 5

technique. The preheating procedure was necessary to

EBSD system (Oxford Instruments) was employed, along

prevent cracking of the EBR layer. The remelting beha-

with a MIRA 3-SEM (Tescan) equipped with an Apollo XP

viour of cast iron materials is mainly determined by their

detector (EDAX/Ametek). Light-optical microscopy was

graphite morphology. For this reason, the beam guidan-

used for measuring the EBR layer thickness (t EBR) as well

ce technique (multi-spot technique, meander technique,

as the compound layer thickness (CLT) after nitriding as

etc.) and/or EB parameters (acceleration voltage, focal

an average value of a minimum of 20 individual values.

position, local and global frequency) were adapted in

Secondary dendrite arm spacing (SDAS) was determined

order to produce crack- and pore-free remelted layers

by means of light-optical micrographs according to the

with surface deformation minimized as much as possi-

BDG guideline P220 (2011) [11]. This value facilitated the

ble. Once the electron beam had passed, the remelted

estimation of both the degree of microstructural refine-

surface layer solidified again almost instantaneously due

ment achieved by the solidification process and the co-

to rapid heat dissipation into the cold core material – a

oling rate. The final SDAS values mentioned were mean

phenomenon known as self-quenching.

values from 10 individual measurements. The treated

Before nitriding, the surface deformations resulting from

surfaces were characterised by Vickers hardness-depth

the EBR process were machined and ground with 1200-

profiles perpendicular to the surface using nominal loads

grit SiC paper. Gas nitriding was carried out in a chamber

of 2.94 N (HV0.3) for EBR layers and 0.49 N (HV0.05) for

furnace at a nitriding temperature of T N = 540 °C for both

nidrided layers. For EBR layers, the minimum and maxi-

different time periods (4 h, 16 h) and nitriding potentials

mum hardness values were estimated from the profiles

K N = 0.1 / 1 / 4 [atm- 1/2].

measured.

The influence of the thermal regime during nitriding on

Results and discussion

the thermal stability of the remelted microstructure was investigated through examination of the pearlitic cast

Microstructure and properties of white-solidified EB

irons. They were annealed under an argon atmosphere

remelted surface layers

at different temperatures and for different time periods,

Independently of the EBR parameters used, all cast

which are associated with both low and high thermal lo-

irons solidified in a metastable manner according to the

ads typical of conventional nitriding treatments.

Fe-Fe 3C system. While the base materials investigated

The particular annealing conditions were 540 °C/4 h,

were hypoeutectic cast alloys, the main fractions of the

540 °C/16 h, 590 °C/12 h, 630 °C/4 h and 670 °C/12 h. The pa-

so-called white-solidified microstructure were primary

rameter combination of 670 °C/12 h is unusual for conven-

austenite dendrites (which decomposed to pearlite) em-

tional nitriding, and may be considered to be an upper limit.

bedded in a eutectic matrix (Fig. 1). The eutectic matrix exhibited typical features of ledeburite, i.e., cementite

Examination methods

plates and a rod-like eutectic. While with an increasing

Cross-sectional specimens were prepared for the mi-

cooling rate (equal to SDAS↓, Fig. 1a, c), the volume

crostructural characterization. All specimens were em-

fraction of rod eutectic (austenite rods in a cementite

bedded in Struers PolyFast ®, ground and polished (final

matrix) decreased, the volume fraction of plate-shaped

step: 1 µm diamond suspension). The nitrided specimens

eutectic microstructure increased. Moreover, analysis of

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Scientific papers - Heat Treatments the element distribution using energy dispersive X-ray

gations during solidification and the formation of another

spectroscopy (EDS) showed a strong heterogeneous

ternary Fe-C-Si phase [12,13] – which typically outlines

silicon (Si) distribution in the ledeburitic microstructu-

the edges of eutectic cementite plates – obviously led to

re. The eutectic cementite could be assumed to be Si-

the Si enrichments observed.

free. Local Si enrichments occurred along the boundary between eutectic cementite and pearlite. Si microsegre-

b) P/S_2.3Si, SDAS↑ = 5.7±0.8 µm

a) P/S_2.3Si, SDAS↓ = 3.7±0.8 µm

c) F/V_2.5Si, SDAS↓ = 3.0±0.4 µm

d) F/V_2.5Si, SDAS↑ = 5.7±0.6 µm

Fig.1 - White-solidified microstructure of remelted surface layers with small and large SDAS of pearlitic P/S_2.3Si (a+b) and ferritic F/V_2.5Si (c+d)

The cooling rate was influenced by the energy input

function of EBR layer thickness or SDAS (cf. Fig. 2a, b).

(beam current). As expected, the lower the energy input,

The hardness values measured in the remelted surface

the thinner the EBR layer thickness (t EBR) and the more

layers were in the range of 600 - 750 HV0.3. Compared to

compact the secondary dendrite arm spacing (SDAS) (Fig.

the initial hardness of the pearlitic (258 - 288 HV0.3) and

2). As a result of the preheating, the cooling rate calcula-

ferritic (167 HV0.3) base materials, the hardness values in

ted from the SDAS was ≤ 900 K/s [14]. Independent of the

the EBR layers were more than doubled (Fig. 2a) or qua-

beam guidance techniques or EBR parameters used, only

drupled (Fig. 2b) after EBR.

marginal differences in layer hardness were apparent as a a) Pearlitic cast irons

b) Ferritic cast irons

Fig.2 - Influence of the EBR layer thickness tEBR on minimum (hollow rhombus ◊) and maximum (solid rhombus ♦) layer hardness and mean secondary dendrite arm distance SDAS (solid circles ●)

During EB remelting, the spatially connected network

more widely separated manner in the microstructure –

of fine graphite lamellae of the GJL cast iron was almost

floated to the surface in the melt, and were thus not com-

completely dissolved, while the graphite morphologies

pletely melted due to their very short exposure time to

of the GJS or GJV cast irons – which were present in a

the electron beam. This graphite was removed during the

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Memorie scientifiche - Trattamenti Termici surface grinding that was necessary before the remelted

Due to its poor melting behaviour (pore formation, surfa-

surfaces could be nitrided. However, such a procedure

ce deformation), only EBR layers with SDAS↑ (4.9 - 6.7 µm)

can lead to a change in the local carbon content, espe-

could be provided for P/L_2.0Si.

cially in thin EBR layers. This lower carbon content led to a higher volume fraction of primary austenite (pearlite

Influence of thermal regime of nitriding on the micro-

after cooling) in the remelted microstructure and, thus,

structure of EBR layers

to somewhat lower layer hardness values (cf. Fig. 2).

For further studies regarding low-temperature annealing

Starting from the initial levels of hardness after EB remel-

and nitriding, EBR layers were selected for each cast iron

ting (hollow symbols in Fig. 3), the hardness decreased

with:

continuously with increasing thermal load.

- low SDAS↓: 3.0 µm - 3.7 µm (t EBR = 0.1 mm - 0.3 mm) and - high SDAS↑: 5.4 µm - 6.7 µm (t EBR = 1.0 mm - 1.3 mm).

Fig.3 - Influence of the thermal regime during annealing on the hardness of EBR layers of pearlitic cast iron.

The massive loss of hardness at 670 °C/12 h was due to graphitizing

SEM investigations of the microstructures annealed up to

3). The decrease in hardness was larger for P/L_2.0Si than

630 °C/4 h revealed that the changes in hardness (Fig. 3)

for P/S_2.3Si up to 630 °C/4 h.

were mostly related to changes in the microstructure of

The higher chromium (Cr) and manganese (Mn) contents

pearlite.

in P/L_2.0Si (cf. Tab. 1) - which are known to dissolve in

The lamellar morphology of the pearlite (Fig. 4a) was lost

cementite and, consequently, stabilize pearlitic cementi-

due to the coarsening of the pearlitic cementite (Fig. 4b).

te against coarsening during annealing [16,17] - could not

The morphology of eutectic cementite (EC) in the EBR

compensate for the lower Si content. At a thermal load of

layers was hardly affected by annealing up to a thermal

630 °C/4h, the pearlitic cementite only became coarser,

load of 630 °C/4 h.

but was largely retained (Fig. 4b).

As determined in [15], the presence of Si in the pearlitic α-iron (Fe) and along the boundary between eutectic cementite and pearlite significantly retarded the coarsening of the pearlitic cementite and stabilized its lamellar morphology.

Higher silicon content hindered the coarsening of pearlitic cementite more effectively, as indicated by the different changes in hardness of P/S_2.3Si and P/L_2.0Si (Fig.

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Scientific papers - Heat Treatments a) Annealed at 540 °C/16 h

b) Annealed at 630 °C/4 h

Fig.4 - SEM images (backscattered electron contrast) of the annealed microstructure of remelted P/L_2.0Si. The microstructural constituents, i.e., eutectic cementite (EC), pearlite (P) and ferrite (F), are labelled.

Cementite appears dark in contrast, while ferrite appears bright

At the maximum thermal load of 670 °C/12 h, the hard-

Refining the microstructure (SDAS↓) by changing the

ness dropped remarkably due to graphitizing. The EBR

EBR parameters increased the amount of possible

layer of P/S_2.3Si was almost completely graphitized at

graphite nucleation sites, while the redistribution of Fe

670 °C/12 h (Fig. 5b), whereas the remelted P/L_2.0Si was

and Si was accelerated by more rapid interface diffusion.

only partially graphitized under the same conditions. The

For this reason, graphitization and ferrite grain growth

microstructure of the remelted P/L_2.0Si contained co-

were more advanced in the smaller remelted micro-

arse temper graphite as well as coarse eutectic cementi-

structures, i.e., the temper graphite particles and the fer-

te plates (Fig. 5c). In contrast to the P/S_2.3Si, the higher

rite grains were larger in the smaller remelted layers than

thermal stability of the eutectic cementite in P/L_2.0Si

in the remelted layers with larger SDAS after 670 °C/12 h

should have resulted from the lower Si content and the

(Fig. 5a,b).

higher Cr and Mn content (cf. Tab. 1).

a) P/S_2.3Si, SDAS↓ = 3.7±0.8 µm

b) P/S_2.3Si, SDAS↑ = 5.7±0.8 µm

c) P/L_2.0Si, SDAS↑ = 6.7±1.1 µm

Fig.5 - Light-optical micrographs of remelted P/S_2.3Si (a+b) and P/L_2.0Si (c) annealed at 670 °C for 12 h, where temper graphite formation occurred

In summary, the low-temperature annealing of the cast

re T N ≤ 540 °C and a nitriding time t N ≤ 16 h without chan-

irons investigated showed that the EBR layers with small

ges to their microstructures or hardness.

SDAS (≥ 3 µm) could be nitrided at a nitriding temperatu-

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Memorie scientifiche - Trattamenti Termici Microstructure and properties of grey- and

investigated.

white-solidified cast irons after nitriding

With the exception of P/L_2.0Si, the ledeburitic surface layers were decarburized considerably during nitriding

At the chosen nitriding temperature of 540 °C, the nitri-

without compound layer (Fig. 6a), such that after a pro-

ding potential K N was selected according to the Lehrer

cess time of 16 h, in particular, the first hardness values

diagram [18] (i.e., the phase diagram for the formation of

HV0.05 measured at a distance of 10 µm from the surface

iron nitrides, which is valid for pure iron) in such a man-

had a significantly lower value than the remaining remel-

ner that nitriding was performed in the α-Fe field (K N = 0.1,

ted and nitrided surface layer (Fig. 6c).

without compound layer), in the γ’ field (K N = 1) and in the

In the P/L_2.0Si additionally alloyed with Mn, Cr and Cu

ε field (K N = 4) for both short (t N = 4 h) and long nitriding

(Tab. 1), the eutectic cementite was not decarburized as

times (t N = 16 h).

much and no decrease in hardness was measured directly

As expected, no compound layer was formed after nitri-

below the surface (Fig. 6b,c).

ding using K N = 0.1 in the pearlitic and ferritic cast irons

a) F/V_2.5Si

b) P/L_2.0Si

c) Hardness-depth profiles, t N = 16 h

Fig.6 - SEM images (backscattered electron contrast) of remelted F/V_2.5Si (a) and P/L_2.0Si (b) with the corresponding hardness-depth profiles after gas nitriding at 540 °C/16 h, KN = 0.1.

While the near-surface pearlitic and eutectic cementite of the F/V_2.5Si was decarburized, the eutectic cementite of the P/L_2.0Si was still largely intact

The phase constituents of the compound layers of the

In simplified terms, Fe carbonitrides and Fe nitrides are

EBR layers, which were formed during nitriding with

not distinguished conceptually in the following, but are

K N = 1 and 4 [atm -1/2] for 16 h, are summarized in Tab. 2.

referred to as Fe nitrides (γ’, ε).

Tab.2 - Composition of the compound layer of the microstructural constituents (eutectic cementite; pearlite) in the remelted pearlitic and ferritic cast irons after nitriding at 540 °C for 16 h

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Scientific papers - Heat Treatments During nitriding in the γ’ field of the Lehrer diagram

For K N = 4, only the ferritic cast iron F/V_2.5Si with the

(K N = 1), all compound layers formed on the remelted and

high Si content and the low Mn content had macrosco-

nitrided cast irons were characterized by a strong lateral

pically homogeneous ε compound layers. For the cast

inhomogeneity in their phase constitution. As shown in

irons alloyed with Mn and Cu, the nitrided pearlite was

Fig. 7, ε/ γ’ double layers were locally present in nitrided

transformed into ε and γ’ nitrides at K N = 4. Furthermore,

eutectic cementite, which was identified by EBSD phase

above the nitrided pearlite, outwardly directed Fe-nitri-

maps. The nitrided pearlitic regions of the pearlitic cast

de growth occurred at the surface (see arrows in Fig. 7).

irons alloyed with Mn and Cu were transformed to γ’. For

These surface nitrides may have been associated with the

the ferritic cast irons with low Mn contents, both γ’ nitri-

formation of pores that appeared near the surface along

des and ε nitrides were detected in the nitrided pearlitic

the eutectic cementite plates and along former α/Fe 3C (θ)

regions.

phase boundaries in pearlite.

Fig.7 - Influence of the Si content and the nitriding potential KN on the phase constitution (blue: α; yellow: θ (Fe3C);

violet: ε; green: γ’) in the compound layer of the EBR layer, measured with EBSD (black: non-indexable areas especially in nitrided pearlite due to pore formation and very fine microstructures)

a)

b)

c)

Fig.8 - SEM images of remelted F/V_2.5Si after gas nitriding at 540 °C/16 h, KN = 1 with Si nitride precipitations in the compound layer (CL) and diffusion layer (DL): (a) cross section, (b) transition CL – DL and (c) end of DL

After all nitriding treatments and independently of K N, na-

the α/Fe 3C (θ) phase boundaries and in α-Fe matrix. The

nodimensional precipitations existed in the compound

exact chemical composition of the precipitations has not

layer (CL) and the diffusion layer (DL) (Fig. 8). These pre-

yet been defined. There is a high probability that the pre-

cipitations were amorphous and only occurred where Si

cipitates were Si nitrides (Si 3N 4) as described by [19-21]

enrichments had been detected, i.e., in particular along

for binary Fe-Si alloys.

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Memorie scientifiche - Trattamenti Termici a) Without prior EB remelting

b) With prior EB remelting (SDAS ⁓ 6 µm)

KN = 1

KN = 4

Fig.9 - Influence of Si content, nitriding potential KN and nitriding time tN on the mean compound layer thickness (CLT) of nitrided grey-solidified (a) and white-solidified (b) cast irons

Nitriding the grey-solidified cast irons with K N ≥ 1 led to

(2.0 - 2.5% Si) of the various cast iron materials had only

the formation of a compound layer, which was inhomo-

a marginal influence on the compound layer thickness

geneous with respect to the layer thickness.

(Fig. 9b). Compared to all other cast irons, P/L-2.0Si

No compound layer was formed on the graphite due to

contained the largest amount of other nitride-forming

the absence of iron. Since the nitrogen could penetra-

alloying elements, such as Mn, Cr and Ti (cf. Tab. 1).

te particularly rapidly to greater depths along the small

This resulted in the lowest compound layer thicknes-

gap between the graphite and the matrix, the local com-

ses of all of the nitriding variants. The mean compound

pound layer thickness (CLT) varied greatly and depen-

layer thicknesses of the EBR layers for K N = 1 were gre-

ded on the size of the graphite particles located near the

ater than those of the grey-solidified surfaces, while for

surface (Fig. 9a).

K N = 4, they were lower than the mean compound layer

The absence of graphite in the white-solidified remel-

thicknesses of the grey- solidified surfaces - independent

ted surface layers resulted in dense compound layers

of the nitriding time (Fig. 9).

with a homogeneous compound layer thickness. Both the dendrite arm spacing of the EBR microstructures (SDAS↓ ~ 3.5 µm; SDAS↑ ~ 6 µm) and the Si content

a) Pearlitic cast irons

b) Ferritic cast irons

Fig.10 - Influence of Si content, nitriding potential KN and nitriding time tN on the diffusion layer thickness (DLT) of nitrided and remelted pearlitic (a) and ferritic (b) cast irons

La Metallurgia Italiana - gennaio 2020

pagina 36

Scientific papers - Heat Treatments a) P/L_2.0Si

b) P/S_2.3Si

a) F/S_2.3Si

b) F/V_2.5Si

Fig.11 -Influence of Si content, nitriding potential KN and nitriding time tN (4h: hollow symbol; 16h: solid symbol)

on the increase in hardness ΔHV0.05 (ΔHV = HVDL – HVEBU or HVBM) of nitrided cast irons in the initial (green symbols) and remelted states

The diffusion layer thicknesses (DLT) were determined

ded ferritic cast irons, the increase in nitriding potential

by means of investigating the depths at which the nano-

to K N = 4 led to a further increase in hardness (Fig. 11c,d).

dimensioned precipitations were visible under the SEM. The DLT increased with increasing nitriding time and/or

Conclusion

nitriding potential up to K N = 1 (Fig. 10). With a further

The influence of gas nitriding conditions by means of

increase in nitriding potential up to K N = 4, the diffusion

the variation of the time and nitriding potential on the

layer thickness remained almost constant (Fig. 10). Wi-

formation of nitriding layers on white-solidified (EBR)

thin the diffusion layer, the inclusion of nitrogen and the

and pearlitic and ferritic cast irons was investigated.

precipitation of special nitrides (e.g. Si, Cr, Ti nitrides

The white-solidified surface layers were generated by

etc., cf. [22]) resulted in an increase in hardness, which

electron beam remelting. By means of the targeted va-

is shown in Fig. 11 as the difference between the initial

riation of the EB process parameters, and in particular the

hardness of the EBR layer (HV EBR) or the base material

energy input (beam current), crack- and pore-free remel-

(HV BM) and the maximum hardness of the diffusion layer

ted surface layers with a layer thickness of ≤ 1.3 mm could

(HV DL). The greatest increase in hardness due to nitri-

be produced. Since the pearlitic and ferritic cast irons

ding was measured by increasing the nitriding potential

investigated were hypoeutectic alloys, the white-so-

in the range of K N = 0.1...1. Nitriding at a nitriding poten-

lidified remelted layers consisted of primary austenite

tial of K N = 4 did not lead to any significant increase in

dendrites (decomposed in pearlite) and a eutectic ma-

hardness. An analogous relationship resulted for the ni-

trix. The thinner the remelted surface layers, the smaller

trided pearlitic cast irons (Fig. 11a,b). In the case of nitri-

the secondary dendrite arm spacing (SDAS) measured.

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Memorie scientifiche - Trattamenti Termici Regardless of the base material of the cast iron (pearlitic

be detected by X-Ray diffraction (XRD). No Si-nitride

or ferritic), the EBR layer hardness was 600 - 750 HV0.3.

precipitations occurred in Si-free eutectic cementite.

For successful implementation of the combination tre-

In compound layer-free nitriding (K N = 0.1), the EBR

atment, the thermal stability of the EBR layer during

layer was decarburized to a considerable extent, with

subsequent nitriding was of crucial importance.

the exception of the P/L_2.0Si alloyed with Mn, Cu and

Therefore, annealing experiments were carried out under

Cr. The compound layers of the investigated cast irons

typical nitriding conditions. It was shown that Si stabili-

formed during nitriding with a nitriding potential of

zed the pearlitic and eutectic cementite against coarse-

K N = 1 exhibited a strong lateral inhomogeneity in their

ning upon low-temperature annealing. The graphitizing

phase compositions. This laterally inhomogeneous

of white-solidified surface layers at temperatures above

constitution of the compound layers was also observed

630°C, however, was promoted by increased Si content

for K N = 4, with the exception of F/V_2.5Si. With its high

and an increased degree of microstructural refinement.

Si content and low Mn content, it was only this F/V_2.5Si

The annealing tests revealed that nitriding treatment at a

ferritic cast iron that formed macroscopically homoge-

temperature of T N ≤ 540 °C for a nitriding time of tN ≤ 16 h

neous ε compound layers at K N = 4.

did not lead to any changes in either the microstructure

The microstructural refinement of the EBR layers (SDAS↓

or the hardness of the EBR layer.

~ 3.5 µm, SDAS↑ ~ 6 µm) and the Si content (2.0 - 2.5 wt.%

During nitriding, nanodimensional precipitations were

Si) of the cast irons investigated had only a marginal in-

formed in both the compound layer and the diffu-

fluence on the compound layer thickness, the diffusion

sion layer, which led to an additional increase in har-

layer thickness and the increase in hardness within the

dness. These precipitations were detected in Si-enri-

diffusion layer.

ched mi-crostructural areas, i.e., in particular along the α/Fe 3C (θ) phase boundaries and in the α-Fe matrix. From this information, it was deduced that these were silicon nitrides. Due to their amorphous nature, they could not

Acknowledgements This study was supported financially by the German Research Foundation (DFG) within the framework of the BU 2568/7-1 project. The authors wish to thank Mr. E. Siegismund (Institute of Materials Engineering, TU Bergakademie Freiberg) for conducting the nitriding experiments, Ms. G. Bittner (Institute of Materials Engineering, TU Bergakademie Freiberg) for performing hardness measurements, and M.Sc. S. Kante (Institute of Materials Science, TU Bergakademie Freiberg) for his support in conducting EBSD measurements.

La Metallurgia Italiana - gennaio 2020

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Scientific papers - Heat Treatments References

[1]

Davis JR, editor. Cast irons. 1. print. Materials Park, OH; 1996. (ASM specialty handbook ®).

[2]

Berns H, Theisen W. Eisenwerkstoffe - Stahl und Gusseisen. 4., bearb. Aufl. Berlin: Springer; 2008.

[3]

Spies H-J. Nitrieren und Nitrocarburieren: Gefüge - Eigenschaften - kontrolliertes Nitrieren. HTM J. Heat Treatm. Mat. 2013; 68(2):86– 96.

[4]

Hillert M, Steinhäuser H. The Structure of White Cast Iron. Jernkontorets Annaler 1960; 144(7):520–59.

[5]

Rickard J, Hughes ICH. Eutectic Structure in White Cast Iron. BCIRA Journal 1961; 9:11–25.

[6]

Park JS, Verhoeven JD. Directional solidification of white cast iron. Metallurgical and Materials Transactions 1996; 27A:2328–37.

[7]

Buchwalder A, Spies H-J, Klose N, Jung A, Zenker R. Effects of Different Microstructural and Hardness Gradients Generated by Single and Combined Surface Treatments with a Nitriding Top Layer. Savannah, USA; 2016 (23rd IFHTSE Congress).

[8]

Buchwalder A, Zenker R, Fritzsch K, Rüthrich K, Nagel K. Potenziale von ausgewählten Randschichtbehandlungsverfahren sowie deren Kombinationen zur Verbesserung des Verschleiß- und Korrosionsverhaltens von Gusseisenwerkstoffen. Mat.-wiss. u. Werkstofftech. 2015; 46(6):550-62.

[9]

Buchwalder A, Spies H-J, Klose N, Hegelmann E, Zenker R. Contribution to the microstructural constitutions and properties of nitrided as-received and remelted cast irons. Nice; 2017 (24th IFHTSE Congress).

[10]

Buchwalder A, Zenker R, Rüthrich K, Nagel K, Griesbach W, Hartwig S et al. Eine neue kombinierte Randschichttechnologie für hochbeanspruchte Gusseisenwerkstoffe. HTM J. Heat Treatm. Mat. 2014; 69(3).

[11]

BDG-Informationszentrum Giesserei. Bestimmung des Dendritenarmabstandes für Gussstücke aus Aluminium-Gusslegierungen.

Juli 2011. Düsseldorf; 2011 Juli 2011. Available from: URL: https://www.bdguss.de/fileadmin/content_bdguss/Der_BDG/Richtlinien/P_220.pdf.

[12]

Spinat P, Brouty C, Whuler A, Herpin P. Etude Structurale de la Phase ‘MnsSi2C’. Acta Crystallogr. 1975; (B31):541–7.

[13]

Malinochka YN, Osada NG, KovaI’chuk GZ. A new structural constituent in Fe-C-Si alloys 1961:299–305.

[14]

Holst A, Buchwalder A, Hollmann P, Zenker R. Influence of Cooling Rate on the Microstructural Features of a Remelted White-Solidified Cast Iron Surface and its Effects on Nitriding Behaviour. Journal of Materials Processing Technology 2019; 271:377–83.

[15]

Kante S, Holst A, Buchwalder A, Leineweber A. Low-temperature Annealing and Graphitizing of Low-alloy White Cast Irons Produced

by Rapid Solidification Technique. Materialwiss. Werkstofftech. 2019; (50):682-95.

[16]

Zhang G-H, Chae J-Y, Kim K-H, Suh DW. Effects of Mn, Si and Cr addition on the dissolution and coarsening of pearlitic cementite during intercritical austenitization in Fe-1mass%C alloy. Materials Characterization 2013; 81:56–67.

[17]

Lv ZQ, Fu WT, Sun SH, Bai XH, Gao Y, Wang ZH et al. First-principles study on the electronic structure, magnetic properties and phase stability of alloyed cementite with Cr or Mn. Journal of Magnetism and Magnetic Materials 2011; 323(7):915–9.

[18]

Lehrer E. Über das Eisen‐Wasserstoff‐Ammoniak‐Gleichgewicht. Zeitschrift für Elektrochemie und angewandte physikalische Chemie 1930; 36:383–92.

[19]

Meka SR, Bischoff E, Rheingans B, Mittemeijer EJ. Octapod-shaped, nanosized, amorphous precipitates in a crystalline ferrite matrix.

Philosophical Magazine Letters 2013; 93(4):238–45.

[20]

Mittemeijer EJ, Biglari MH, Böttger AJ, van der Pers, N. M., Sloof WG, Tichelaar FD. Amorphous precipitates in a crystalline matrix; precipitation of amorphous Si3N4 in α-Fe. Scripta Materialia 1999; 41(6):625–30.

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van Landeghem HP, Gouné M, Bordère S, Danoix F, Redjaïmia A. Competitive precipitation of amorphous and crystalline silicon nitride in ferrite: Interaction between structure, morphology, and stress relaxation. Acta Materialia 2015; 93:218–34.

[22]

Mittemeijer EJ, Somers MAJ. Thermochemical Surface Engineering of Steels. Amsterdam: Woodhead Publishing is an imprint of Elsevier; 2014. (Woodhead publishing series in metals and surface engineeringnumber 62).

La Metallurgia Italiana - January 2020

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Memorie scientifiche - Trattamenti Termici

Material-dependent aspects for the use of nitriding as a single or duplex treatment in improving the tribological load-bearing behaviour of cast irons (part II) A. Buchwalder, N. Klose, R.Zenker The author’s previously published studies proved the potential of the combined surface treatment of electron beam remelting and nitriding for improving the load-bearing behaviour of cast irons. However, further investigations indicated that a general statement could not be made regarding the suitability and benefits of the duplex treatment. This paper deals with investigations regarding the influence of cast iron base materials differing in graphite morphologies (lamellar, spheroidal), Si content and/or matrix microstructure and the nitriding conditions on the nitriding layer achieved (thickness, phase composition, hardness). Furthermore, the paper reports on the impact of these factors on the depth-dependent wear behaviour under different loading conditions. The nitriding processes for the cast iron (single treatment) and for the remelted cast iron surfaces (duplex treatment) were carried out under identical conditions. Differences in nitrogen concentration profiles and phase compositions resulted. For this reason, detailed investigations were carried out on the depth-dependent microstructural constitutions (phase composition (ε : γ`), porous zone), the hardness and the wear behaviour of the nitriding layers. This was realized by incremental layer removal using a polishing procedure. Each depth state was characterized by XRD analysis (Cu-Kα radiation), the indentation hardness, the scratch energy density and the specific wear coefficient. It was shown that in the case of the nitrided cast iron after the single treatment, the graphite morphology and the load level had the greatest impacts on wear behaviour, followed by the material and the specific nitriding layer construction. This is crucial for defining the potentials and limitations of the single and combined surface treatments with respect to the improvement of the tribological load-bearing behaviour of cast iron.

KEYWORDS: CAST IRON - EB REMELTING - NITRIDING - WEAR - DUPLEX

Introduction With nitriding (N), a wear- and corrosion-resistant surface layer can be produced on cast irons. However, the very soft graphite - specific to this material - interferes with this nitriding layer. Thus, the functionality of the surface layer applied is limited. Electron beam remelting (EBR) of the cast iron surface leads to the dissolution of

Anja Buchwalder Normann Klose Rolf Zenker TU Bergakademie Freiberg

Institute of Materials Engineering, Germany

near-surface graphite and, as a result of self-quenching,

La Metallurgia Italiana - gennaio 2020

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Scientific papers - Heat Treatments to the formation of a white solidified microstructure

Analogously to untreated cast iron, the ‘free graphite’ re-

(graphite → ledeburitic cementite) with high hardness

sults in a low coefficient of friction µ after nitriding due

(600 HV0.3 - 750 HV0.3) [1]. During subsequent nitriding,

to its lubricating effect. Depending on the material, nitri-

a closed layer is formed [2,3]. The nitriding mechanism

ding treatment and test parameters, values of µ = 0.1...0.2

of the heterogeneous ledeburitic microstructure is very

[10]...0.3 [11] were measured in unlubricated tests.

complex due to the high Si content, which is distributed inhomogeneously (though primarily in the intercarbidic

In wear tests (pin-on-disc test, unlubricated, FN = 32.8 N, s = 3000 m),

regions) in combination with the high volume fraction of

Nicoletto et al. [10] show that the friction behaviour is

cementite.

determined more by the nitriding process while the wear

The formation of the Fe nitrides and the growth of the

behaviour is determined by the material (and by the graphi-

compound layer (CL) follow both the thermodynamics (cf.

te morphology in particular), since only the matrix micro-

Lehrer diagram [4]) and the kinetics known from the nitri-

structure is hardened and not the graphite. At low sliding

ding of unalloyed steels, though only to a limited extent.

speeds (< 0.5 m/s), the wear rate of spheroidal graphi-

On the one hand, Si inhibits the formation of Fe nitrides,

te cast iron is lower than that of lamellar graphite. When

while on the other hand, Si itself functions as a nitride

subjected to stress, the interconnected graphite lamellae

former with weak interaction [5,3], i.e., the proportion of

and their irregular shape favour the fractionation and de-

unbound silicon changes progressively during nitriding.

tachment (wear particles) of the hard nitriding layer areas

Consequently, the kinetics of Si nitride precipitation in-

immediately surrounding them. In contrast, the graphi-

fluences the kinetics of compound layer growth in the le-

te spheres remain undeformed, which indicates that the

deburitic microstructure. The exact influence of Si on Fe

harder nitriding layer matrix surrounding the soft graphite

nitride formation - and especially on the suppression of

spheres compensates for the applied contact stresses bet-

γ’ nitride formation at the expense of ε nitride - is poorly

ter (i.e., in a more elastic manner) [10].

understood [3]. One conceivable reason for this could be the low degree of solubility of Si in ε nitride. Kante et al. [6]

At higher sliding speeds ≥ 1 m/s, the graphite morphology

present an initial model of time-dependent microstructu-

has hardly any influence. In lubricated wear tests (ball-on-

ral development during the nitriding of Si-rich pearlite.

disc test; counter body (CB): Al2O3, FN = 44 N, s = 6000 m), the higher the roughness and the lower the hardness of

Along with microstructural composition, pore formation

the nitriding layer, the higher the wear rate [12].

also influences the load-bearing behaviour. The causes of pore formation can be manifold. The main cause is the

While Holst et al. [13] provide an overview of the influence

metastability of the ε nitride, which leads to the migration

of nitriding conditions on the layer formation and hardness

of nitrogen atoms into dislocations and voids which then

of cast iron, the depth-dependent layer microstructure

combine to form nitrogen molecules [7]. Furthermore, the

and the resulting tribological behaviour are presented and

hypothesis is presented that Fe diffusion along the grain

discussed in this paper. Therefore, one set of nitriding pa-

boundaries leads to an oppositely directed diffusion of

rameters (540°C, 16 h, KN = 6) was used for treating three

voids into the interior of the ε nitrides, which triggers pore

different cast irons in states either with or without electron

formation (the Kirkendall effect) [8,9].

beam (EB) remelting.

The detailed investigations were carried out by analysing

The harder remelted microstructure also serves as a sup-

samples, which were polished stepwise perpendicular to

porting layer for the hard, thin compound layer. In the case

the surface and down to a defined depth. Each surface le-

of surface layers treated in this way (EBR+N), a (partially)

vel was tested and analysed. This was crucial for defining

significant improvement in the tribological and corrosi-

the potentials and limitations of the single and combined

ve load-bearing behaviour could be demonstrated when

surface treatments with respect to the improvement of the

compared to the merely nitrided cast iron (N) [2,3].

tribological load-bearing behaviour of the cast iron.

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Memorie scientifiche - Trattamenti Termici Experimental

tion, and in particular in their Si contents (cf. column 4 in Tab. 1), graphite morphologies (speroidal (S); lamellar

Base materials

(L)) and/or matrix microstructures (pearlite (P), ferrite (F))

The investigations were carried out on three conventio-

(Tab. 1, cf. I. in Fig. 2a-c).

nal cast irons which differed in their chemical composi-

Tab.1 - Chemical composition (GDOES), phase composition (XRD) and hardness of the substrates used as well as roughness after nitriding

Substrate

Element concentration [wt.%]

Microstructure

Hardness HV1

EN-GJ…

Fe

C

Si

Mn

P

S

Cu

Cr

Graphite

Matrix

Initial

EBR

S400

Base

3.40

2.67

0.14

0.024

0.050

0,02

0.01

S

F

156 ± 26

865 ± 16

S600

Base

3.60

2.33

0.34

0.030

0.010

0.67

0.03

S

P/F

321 ± 17

794 ± 19

L250

Base

3.20

2.11

0.08

0.024

0.010

0.76

0.17

L

P/F

238 ± 22

644 ± 14

Surface treatments

Sampling was performed by two different methods (see

Electron beam remelting (EBR) was used to generate the

Fig. 1). The focus in this article was on longitudinal mi-

ledeburitic surface layer microstructure.

crosections, which were realized by the stepwise remo-

Preheated samples (450°C (725 K)) were remelted using

val of the surface (approx. Δ2 µm) by polishing (Fig. 1b).

the meander beam deflection technique with a continuous

The thickness of the material removed was calculated by

sample feed rate of 0.5 mm/s and a beam current of 15 mA.

measuring the diagonals d1, d2 produced by Vickers har-

As a result of the rapid solidification, a ledeburitic layer

dness indentation (HV1) (Fig. 1b). Each (subsequently re-

with pearlite in the intercarbidic areas was generated ac-

moved) surface Ai+n was characterized by means of:

cording to the metastable system Fe-Fe3C.

•

Roughness measurements

There were hardly any differences in the microstructure

•

Glow

of the EB remelted layers, irrespective of the base material used. Therefore, an example of the EBR microstructu-

optical

emission

spectroscopy

(GDOES): element concentration-depth profiles •

re of S600 is shown in I. of Fig. 2d. Further information about the microstructural details is given in part I of this

discharge

Surface hardness (SH) measurements with F N = 9.81 N (HV1): mixed hardness value

•

Hardness under increasing loading conditions up to

contribution [13].

Fmax = 0.05 N (50 mN) by determining indentation har-

The deformed EBR surfaces were machined flat and

dness HIT according to DIN ISO 14577 1-3: local layer

ground (to SiC 1200) before nitriding.

hardness

The gas nitriding process was carried out at a temperature

•

of 540°C (815 K) for 16 h and with a nitriding potential of KN = 6.

Characterization methods

La Metallurgia Italiana - gennaio 2020

X-ray diffraction (XRD) with Cu-Kα radiation: phase composition and quantity distribution (vol.%)

•

Light optical (LOM) and scanning electron microscopy (SEM): microstructure, geometry of the layer

pagina 42

Scientific papers - Heat Treatments

•

and wear tracks after scratch testing and pin-on-disk

abrasive/adhesive wear behaviour by calculated spe-

testing

cific wear coefficient kv= W v/(FN∙s) (W v… wear volume

Scratch testing using a Rockwell indenter with a

(measured by 3D profilometer), F N… normal load, s…

constant load of 20 N (preload: 2 N): abrasion wear

sliding distance (both were defined)) and recorded

resistance by calculating scratch energy density

friction coefficient µ

w R= F T /A R (F T… tangential force (recorded); A R…

•

scratched cross section (measured by 3D profilome-

In the diagrams presented herein, the respective sample

ter))

variant or measuring direction used is indicated by a co-

Unlubricated pin-on-disc tests (counter body: WC-

ordinate system.

Co ball (Ø 6 mm); FN = 20; s = 100 m, v = 0.15 mm/s):

Fig.1 - Variants of samples tested

Results and discussions

graphite was washed out during mechanical surface preparation, round craters remained on the nitrided surface

Characterization of the initial states after single and

of the cast iron with spheroidal graphite (II. in Fig. 2a, b).

combination treatment

For the cast iron with lamellar graphite, only small unco-

Despite identical surface preparation, the surface after ni-

vered areas were visible using SEM (II. in Fig. 2c).

triding was strongly influenced by the substrate’s micro-

For both graphite morphologies, bulging of the com-

structure (II. in Fig. 2) and, for cast irons, this is mainly the

pound layer was detected at the edges of the respective

graphite morphology (cf. Fig. 2a, b and c) or graphite size

graphite (see 3D profiles in Fig. II of 2a-c).

(cf. Fig. 2a and b).

The characterization of such inhomogeneous surfaces by the determination of classical roughness parameters such as Ra, Rz, etc. is not effective.

Near-surface graphite with a sufficiently large degree of expansion usually remains uncovered by a nitriding layer.

Depending on the proportion of outwardly directed compound layer growth and/or the degree to which the

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Memorie scientifiche - Trattamenti Termici

a) S400

b) S600

c) L250

d) EBR layer (e.g. S600)

Fig.2 - Influence of the (I.) initial microstructure of the different substrates on the (II.) surface formation after nitriding (540°C, 16 h, KN = 6)

Within the compound layer and depending on the che-

examined after gas nitriding. Exceptions to this, however,

mical composition, the N concentration-depth profiles

were large graphite spheres that lay on the surface.

(GDOES) of the respective treatment states (untreated or

These spheres could also be enclosed by a pore seam. In

remelted) exhibited hardly any differences (cf. II in Fig. 3a

the combined treated surface layers (EBR+N) under the

and b). Only the nitrogen diffusion depth was significantly

selected nitriding conditions, pores could be detected

lower in the EBR state.

over almost the entire thickness of the compound layer

There was also nitriding-induced decarburization in both

in the intercarbidic regions. The latter could also be seen

treatment states, which was much more pronounced in

in the significantly higher maximum contents and depth

the case of the EBR+N state (II. in Fig. 3b), i.e. the C con-

diffusion of oxygen for the combination treatment com-

tent dropped very steeply at the transition between the

pared to the single treatment (cf. II. in Fig. 3a and b).

compound and diffusion layers (Cmin = 0.3 ma.%). For the

The solely nitrided cast iron had the typical phase com-

nitrided cast iron base material (BM+N), however, the C

position of compound layers on steel - consisting of ε

content decreased gradually (II. in Figure 3a). Further-

nitride with a maximum fraction and γ` nitride with a mi-

more, the decarburization (measured at the minimum

nimum fraction at the surface, which then counter each

C content) was significantly lower than for EBR+N. The

other by increasing or decreasing with increasing distan-

decarburization (C min = 0.4→0.8 ma.%) and the depth of

ce from the surface (I. in Fig. 3a).

the porous zone PZ (t PZ = 1.3→2.9 µm) tended to incre-

The compound layers on the remelted substrates pre-

ase with increasing nominal C content in the cast iron

dominantly contained ε nitride (60...90 vol.%) over the

(L250→S400→S600, cf. Tab. 1).

entire depth (I. in Fig. 3b). This resulted from the low

The nitrided cast iron had a pore seam which was typi-

N solubility of carbides (Fe3C), as a result of which they

cally found in the near-surface area - analogous to steels

converted preferentially into nitrides. The formation of γ`

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Scientific papers - Heat Treatments nitride was suppressed due to the restricted carbon solu-

in the compound and diffusion layer, as was shown by

bility, and ε nitride was preferred (I. in Fig. 3b).

comparative investigations by the authors between a C15

It could be assumed that ε carbonitride was formed as a

carbon steel and a cast iron with a ferritic matrix (S400)

result of the increased carbon activity in the environment

[3].

[14].

These were mostly amorphous precipitates [15] and

In addition to the formation of Fe nitrides, the precipita-

could not be analyzed using XRD, so that they will be left

tion of Si nitrides also influenced the increase in hardness

out of subsequent discussions of the results.

a) Nitrided cast irons (BM + N)

b) Nitrided EBR layers (EBR+N)

Fig.3 - Influence of the substrate used on the depth-dependent (I.) fractions of ε and γ`nitrides (XRD Cu-Kα) and (II.) element concentration profiles (GDOES) (540°C, 16 h, KN = 6)

Investigations on the depth-dependent evolution of

of 700 HV1, the indentation depth was approx. 7 µm. With

hardness and wear behaviour

respect to the Bückle rule, it could be assumed that the

The depth-dependent hardness curves shown in Fig. 4

hardness values contained information from a depth of

resulted from the layer microstructure (phase composi-

up to approx. 50 µm. Essentially, the surface hardness of

tion and defects) described above. When interpreting the

the combined treated surfaces (EBR+N) in both the com-

results, the different measured loads (I. 9.81 N (HV1), II.

pound layer and the diffusion layer was approx. 200 HV1

0.050 N) and load directions (both vertical (I., III) and pa-

higher than on the solely nitrided cast iron (cf. I. in Fig. 4a

rallel (II.) to the surface) had to be taken into account.

and b).

The hardness measurement with a load of 9.81 N provi-

This was a reliable reflection of the significantly better

ded a composite hardness (I. in Fig. 4) as a function of

supporting effect for the compound layer in the case of

the distance beneath the surface, consisting of hardness

the much harder remelted carbidic microstructure.

parts of the compound layer and the underlying diffusion

While the hardness curves of spheroidal graphite cast

layer or substrate. For example, with an average hardness

irons were relatively comparable after both single and

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Memorie scientifiche - Trattamenti Termici combination treatments (cf. I. in Fig. 4a and b), those of

rous zone. In nitrided cast iron, this was formed relatively

lamellar graphite cast iron (L250) deviated significantly

uniformly up to a material- and nitriting-dependent dep-

from these and lay at hardness levels that were lower by

th.

approx. 150 HV1 (EBR+N) to 300 HV1 (L250+N).

The exception was the graphite environment. Due to the The local layer hardness resulting from the microstructu-

favourable diffusion and penetration conditions for the

re (II. and III. in Fig. 4) could be determined by means of

nitrogen at the graphite/matrix gap interface (cf. [14,16]),

hardness measurements under increasing loading condi-

nitrogen could penetrate to greater depths and led to

tions, with a maximum load of 0.05 N.

pore formation.

With a maximum penetration depth of 0.5 µm, this hard-

In contrast, pore formation during nitriding of the EBR

ness value contained information from a depth of approx.

layers occurred mainly in the intercarbidic regions or at

3 µm, i.e. with the exception of the influence of pores in

the interface with the ledeburitic carbides and under the

the area near the surface, the supporting effect (for exam-

selected nitriding conditions over almost the entire dep-

ple, that provided by the hard carbides) played only a su-

th of the compound layer.

bordinate role - at least in the compound layer.

This led to the influence of the layer hardness measured As already shown above, the compound layers genera-

perpendicular to the surface, which is shown in III. in Fig.

ted in the EBR layer consisted mainly of ε nitrides (cf. II.

4b.

in Fig. 3b), which led to somewhat higher average layer

While the hardness measurements carried out parallel

hardness values of 12.0 ± 0.9 GPa (II. in Fig. 4b) compared

to the surface (II. in Fig. 4) on the etched cross section

to the compound layers (cf. II. in Figure 3a) that contained

facilitated a selection of the ‘defect-free’ microstructural

more γ` nitrides on the cast iron (10.5 ± 0.8 GPa, II. in Fig.

areas - i.e., the matrix microstructure in nitrided cast iron

4a). The differences due to the chemical composition of

and the carbides in combined treated cast iron - it was

the different cast irons were rather marginal for the two

not possible to achieve this differentiation with certainty

states (BM+N; EBR+N).

in the hardness measurement perpendicular to the surface (III. in Fig. 4).

From this, it could be concluded that the large differences in the macro hardness measurements (I. in Fig. 4a)

The surfaces of the unetched ablated samples appeared

resulted from the influence of lamellar graphite and the

almost uniformly white, as the pores or graphite were

resulting strong fragmentation of the compound layer

partially lubricated during polishing. Thus, selection of

(loss of load support).

the pore-free carbides or graphite (especially with the

In the case of the combination treatment (EBR+N), the

narrow lamellae of the L250) was not possible and the

lower C content of the L250 (cf. Tab. 1) and the associated

hardness was at a lower level (III. in Fig. 4) than the hard-

lower hardness of the EBR layer influenced the suppor-

ness measurement at the cross sections (II. in Fig. 4).

ting effect, and led to a lower overall hardness of 150 HV1.

In the lower, pore-free area of the compound layer, the Interestingly, the profile of the layer hardness measured

true layer hardness was measured at 11.9 ± 0.6 GPa for

parallel to the surface (II. in Fig. 4) differed significantly

the nitrided cast iron (BM+N) and at 13.6 ± 2.0 GPa for the

from that measured perpendicular to the surface (III. in

combined treated cast iron (EBR+N).

Fig. 4).

This applied in particular to the significant decrease in hardness in the near-surface area resulting from the po-

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Scientific papers - Heat Treatments

a) Nitrided cast irons (BM + N)

b) Nitrided EBR layers (EBR+N)

c) SEM image of the CL

Fig.4 -Influence of the substrate used on the depth-dependent (I.) surface hardness SH (9.81 N), the indentation hardness HIT (0.050 N) measured both (II.) parallel and (III.) perpendicular to the surface

The microstructural formation described and the re-

gential forces F T were approx. twice as high after the sin-

sulting hardness profiles influenced the tribological lo-

gle treatment than after the combination treatment.

ad-bearing behaviour, which was investigated under both

At greater depths, the scratch energy density for the com-

purely abrasive (scratch test with constant normal load)

bination treatment (EBR+N) was 2-4 times higher

and under abrasive-adhesive loading conditions (pin-on-

than for the single treatment (BM+N)(cf. I. in Fig. 5a and b).

disk test) (Fig. 5).

Failure modes generated by scratch testing with constant

On the basis of the scratch track produced by a loaded

loading were quite different on both of the nitrided sub-

(20 N) diamond cone, the so-called scratch energy den-

strates. It should be noted that comparisons between dif-

sity w R was calculated from the scratched cross-sections

ferent samples are only valid if the mechanism of failure

measured. The higher the scratch energy density in this

is the same. As a result of bending of the hard compound

measurement, the better the abrasive wear resistance.

layer into the softer cast iron, conformal through-thick-

While the nitrided cast iron within the compound layer

ness cracking occurred, both perpendicular C1 and paral-

had an almost constant scratch energy density of approx.

lel C3 (at the edges) to the scratch direction v (see LOM

21. 9 ¹ 6 N/mm² (I. in Fig. 5a), it was highly depth-depen-

image in I. in Fig. 5a). Furthermore, cracking around the

dent for the combined treated surface layers (EBR+N).

graphite C2 was observed. It was assumed that after the

In the near-surface area, the scratch energy density was

indenter passed and the scratched region was unloaded,

comparable with the single treatment (I. in Fig. 5b). Com-

the elastic deformation in the coating-substrate system

parable scratch energy densities meant that the scratched

was relaxed to some extent. However, due to the pla-

cross-sectional areas were 3-4 times larger and the tan-

stic deformation in the substrate, which resulted in the

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Memorie scientifiche - Trattamenti Termici formation of the visible scratch track (see LOM image in

In contrast, the combination treated surfaces did not

I. in Fig. 5a), it was not possible to completely relax the

crack or form other modes of failure (see LOM image in

substrate’s elastic deformation, and a residual stress re-

I. in Fig. 5b).

mained.

With decreasing hardness in the diffusion layer, the scra-

It is possible that the scratch energy density was influen-

tch energy density also decreased.

ced by this process, i.e., that it increased.

a) Nitrided cast irons (BM + N)

b) Nitrided EBR layers (EBR+N)

Fig.5 -Influence of the substrate used on the depth-dependent (I.) scratch energy density wR, determined by scratch testing under constant load (FN = 20 N), (II.) specific wear coefficient kV, and (III.) average friction coefficient µ, both determined by unlubricated pin-on-disk testing (FN = 20 N)

The wear behaviour under abrasive-adhesive loading

Fig. 5a). The reason for this was that the coefficient of fri-

conditions was determined not only by the hardness of

ction of the EBR+N variants (µ EBR+N = 0.78 ± 0.04, III. in Fig.

the nitriding layer but also by the coefficient of friction

5b) was essentially three times higher than that of the nitrided

µ and pore formation. The latter led to an increase in the

cast iron with spheroidal graphite (µGJS+N = 0.26 ± 0.09, III. in

specific wear coefficient kV on both substrates (cf. II. in

Fig. 5a). The diagram in section III. of Fig. 5a shows very

Fig. 5).

impressively how the friction behaviour is influenced by

If the porous zone was neglected, for the EBR+N variant a

the graphite morphology and its size at comparable hard-

specific wear coefficient of kV, EBR+N = 0.69 ± 0.25∙10 N∙m/mm³

ness. Under the given load conditions, the large graphite

(II. in Fig. 5b) averaged over all materials results, which

spheres (dS = 36 ± 12 µm) of the S400 led to the formation

would be approximately twice as high as that of the only

of a closed graphite lubricant film (see black film in wear

nitrided cast iron (kV, GJS+N = 0.31 ± 0.11∙10 N∙m/mm³, II. in

tracks in the micrographs in II. of Fig. 5a) that exhibited

-6

-6

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Scientific papers - Heat Treatments excellent constant friction conditions (µ = 0.21 ± 0.02).

be determined:

The somewhat smaller graphite spheres of the S600 (dS

(1) The course of the N diffusion profiles (GDOES measu-

= 24.5 ± 0.4µm) led to somewhat unsteady conditions (in

rements) within the compound layer differed only margi-

terms of the formation of the graphite lubricant film) and a

nally between the two substrate variants (BM+N; EBR+N),

slightly increased coefficient of friction with a larger ran-

with the exception of the final depth reached. The diffu-

ge of fluctuation (µ= 0.32 ± 0.10). With the L250, the lubri-

sion layer was characterized by strong N enrichment due

cating effect of graphite no longer existed at all due to the

to the presence of Si.

small effective cross-sectional dimension, so that analo-

The nitriding-induced decarburization in the area of the

gous (non-lubricated) friction conditions (µ= 0.67 ± 0.04)

compound layer was much more pronounced for the

existed as in the EBR+N state, which was also reflected

combination treatment EBR+N.

in comparable specific wear coefficients (cf. L250 in II. in Fig. 5a and b).

(2) While the nitrided cast iron exhibited the typical phase

While the specific wear coefficient for the solely nitrided

structure of compound layers on steel - consisting of ε

L250 (II. in Fig. 5a) increased in the transition region of

nitride with a maximum fraction and γ` nitride with a mi-

the compound to the diffusion layer due to the drop in

nimum fraction at the surface, which then countered each

hardness (II. in Fig. 4a), it decreased successively for the

other by increasing or decreasing with increasing distan-

combination treated variant (II. in Fig. 5b).

ce from the surface - the compound layers on the remel-

As already mentioned in the scratch test, a crack network

ted substrates mainly contained ε nitride (60...90 vol.%)

formed around the soft graphite under loading conditions

over the entire depth.

(see magnified micrograph in III. in Fig. 5a). In addition to its good lubricating effect, graphite also acted as a reser-

(3) The nitrided cast iron had a pore seam typically found

voir for abrasive wear particles, i.e., it exhibited smearing

in the near-surface area, which was analogous to steels

(see magnified micrograph in III. in Fig. 5a) and, there-

after gas nitriding. Exceptions, however, included large

fore, the surface was partially hardened. When the wear

graphite spheres that lay on the surface, whereby these

coefficient was determined by the profilometric measu-

spheres could also be enclosed by a pore seam. In the

rement of the wear volume, the specific wear coefficient

combined treated surface layers (EBR+N) under the se-

was distorted (too low) for the cast iron with spheroidal

lected nitriding conditions, pores could be detected over

graphite. In contrast, the wear removal on the combined

almost the entire thickness of the compound layer in the

treated surfaces was uniform and flat (see wear tracks and

intercarbidic regions. Pores reduced the layer hardness

magnified micrograph in III. in Fig. 5b).

values measured at low levels of loading (50 mN), as did the scratch energy density as a characteristic of abrasive

Summary and conclusions

wear resistance. In contrast, the specific wear coefficient

Within the scope of this paper, comparative investiga-

- as a characteristic of the abrasive-adhesive wear resi-

tions were carried out on the tribological load-bearing

stance - increased. In all cases, this meant a deterioration

behaviour of various cast irons after nitriding BM+N (sin-

of the wear behaviour.

gle treatment) and remelting and subsequent nitriding EBR+N (combination treatment).

(4) In the layer area with low or no porosity, the scratch

Nitriding was carried out for all substrate variants at 540°C

energy density of the combination treated samples was

/ 16 h, K N = 6. The focus was on the effects of the chemi-

2-4 times higher than for the single treatment. The speci-

cal composition (particularly the C and Si contents) of the

fic wear coefficient of the solely nitrided cast iron depen-

cast irons as well as the graphite morphology (sphere,

ded strongly on the graphite morphology and the size.

lamella) on the depth-dependent phase composition in

On the one hand, the graphite served as a lubricant and,

the compound layer, the hardness characteristics and the

on the other hand, as a reservoir for abrasive particles,

wear behaviour. The following essential findings could

which resulted in completely different loading condi-

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Memorie scientifiche - Trattamenti Termici tions than in the case of the combination treatment. Only

The results presented show that a direct comparison of

cast iron with lamellar graphite was comparable, since

the tribological behaviour of the two nitrided substra-

the mechanisms described above did not work. Both the

tes (BM+N and EBR+N) on the basis of the determined

coefficient of friction and the depth-dependent course of

wear values is difficult. Both the stress mode (lubricated,

the specific wear coefficient were comparable with those

non-lubricated) and the resulting contact conditions as

of the combination treatment.

well as the wear mechanism were very different.

At the transition to the diffusion layer, the wear coeffi-

An essential aspect for further optimization of the nitri-

cient of the nitrided cast iron increased due to the absen-

ding process or the preferred nitriding result is the re-

ce of a supporting effect, while that of the combination

duction of pores and/or decarburization.

treatment decreased further.

Acknowledgements The authors wish to thank Dr. C. Schimpf (Institute of Materials Science, TU Bergakademie Freiberg) for his assistance in carrying out the XRD measurements and Ms. G. Bittner (Institute of Materials Engineering, TU Bergakademie Freiberg) for performing hardness measurements. Many thanks also go to Dr. H. Zimdars (Institute of Materials Engineering, TU Bergakademie Freiberg) for carrying out the nitriding experiments. Furthermore, the article contains partial results that were generated in the context of the diploma theses of Mr. B. Haas, Mr. R. Wagner and Ms. J. Fitz.

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Scientific papers - Heat Treatments

References [1]

Holst A, Buchwalder A, Hollmann P, Zenker R. Influence of cooling rate on the microstructural fea-tures of a remelted white-solidified cast iron surface and its effects on nitriding behaviour. Journal for Materials Processing Technology 2019;271:377-383

[2]

Buchwalder A, Zenker R, Fritzsch K, Rüthrich K, Nagel K, Griesbach W. Potenziale von ausgewählten Randschichtbehandlungsverfahren sowie deren Kombinationen zur Verbesserung des Verschleiß- und Korrosionsverhaltens von Gusseisenwerkstoffen.

Materials Science and Engineering Technology 2015;46:550-562

[3]

Buchwalder A, Spies HJ, Klose N, Hegelmann E, Zenker R. Contribution to the microstructural con-stitutions and properties of nitrided as-received and remelted cast irons. Proc. of 24th IFHTSE Congress, Nice, June 26-29 2017, CD

[4]

Lehrer E. Über das Eisen-Wasserstoff-Ammoniak-Gleichgewicht, Z. Elektrochem. Angew. Phys. Chem. 1930;36:383–392.

[5]

Mittemeijer EJ, Somers MAJ. Thermochemical surface engineering of steels: Improving materials performance, Elsevier, 2014.

[6]

Kante S, Leineweber A. Two-Phase and Three-Phase Crystallographic Relationships in White-Solidified and Nitrided Fe-C-Si cast iron. Acta Materialia, 2019;170:240-252.

[7]

Prenosil B. Einige neue Erkenntnisse über das Gefüge von um 600°C in der Gasatmosphäre carbonitrierter Schichten. HTM 1973;28:157-164

[8]

Matauschek J, Trenkler H. Neue Erkenntnisse über den Aufbau und die Porigkeit der beim Badnitrieren erzeugten Verbindungsschicht. HTM 1977;32:177-181.

[9]

Hoffmann R, Mittemeijer EJ, Somers MAJ. Verbindungsschichtbildung beim Nitrieren und Nitrocarburieren. HTM 1996;51:162– 169.

[10]

Nicoletto G, Tucci A, Esposito L. Sliding wear behavior of nitrided and nitrocarburized cast irons. Wear 1996;197:38-44.

[11]

Nie X, Wang L, Yao ZC, Zhang L, Cheng, F. Sliding wear behaviour of electrolytic plasma nitrided cast iron and steel. Surf Coatings Technol. 2005;200:1745-1750.

[12]

Yang Z. The Use of Nitriding to Enhance Wear Resistance of Cast Irons and 4140 Steel. Dissertation, Paper 4714, University of Windsor, 2013.

[13]

Holst A, Buchwalder A, Hollmann P, Zenker R. Influence of Gas Nitriding Conditions on Layer Structure Formation on Grey and White Solidified Cast Irons (Part I). Conf. Proc. ECHT 2019. 2019 Jun 5-7, Bardolino, Italy. Milan: AIM; 2019: paper 060

[14]

Keller K. Vergleichsuntersuchungen an gas-, bad- und ionitriertem Kugelgraphitguss und Nitrierstahl. Ph. D. Thesis. TH Aachen.

1969

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Mittemeijer EJ, Biglari MH, Böttger AJ, van der Pers NM, Sloof WG, Tichelaar FD. Amorphous precipitates in a crystalline matrix, precipitation of amorphous Si3N4 in α-Fe. Scripta Materialia 1999;41:625-630.

[16]

Wiegand H, Hentze H. Dauerschwingfestigkeit und Verschleißverhalten von salzbadnitrierten Zahnrädern aus Gusseisen mit Kugelgraphit. Metalloberfläche 1959;13:238-242

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Memorie scientifiche - Trattamenti Termici

Simulation of heat treatment of jominy specimen to improve quality of automotive gear components E. Morgano, C. Viscardi, L. Valente Microstructure and mechanical performance play a key role for gear box reliability. According standard quality assurance procedure steel grade analysis is a fundamental base to identify correct mechanical properties achievable after heat treatment cycle. With the aim to make more robust production cycle of gear box components, samples are cut from parts and Jominy test is performed on this specimen. In case of small components like gears, specimen for Jominy test cannot be obtained and the only approach to validate process is supplier certificate obtained on steel before heat treatment.

Quenchability of component can be different from data reported on nominal supplier datasheet with the risk to have failure during testing or exercise of gear box; tests on failed parts demonstrated that chemical composition variation, notwithstanding inside allowed range, drove to lower mechanical performance according microstructure not properly quenched.

In order to validate heat treatment cycle, microstructure and mechanical properties, Jominy test has been virtualized into heat treatment simulation software. Real Jominy tests have been transferred inside simulation software for different steel grades used for automotive gear box application. A model to forecast microstructure and mechanical properties for Jominy specimen is built and benchmarked with experimental data. Alignment between simulation results and reality allows to transfer model definition from Jominy samples to complete real components, improving quality assurance capability for real process thanks to usage of heat treatment cycle virtualization for gear box realized with different steel grades.

KEYWORDS: HEAT TREATMENT SIMULATION - QUENCHING - MICROSTRUCTURE HARDNESS - RESIDUAL STRESS - DISTORSION - DEFORM HT

Introduction Gear box performances are based on mechanical properties that each part can achieve after manufacturing cycle and in particular after final heat treatment. Reliability of gear box assembly requests to validate each component of the assembly and respect of specification. Gears for automotive transmission are in general realized using

Enrico Morgano FCA CRF Metal Bulk, Italy

Cristian Viscardi, Lorenzo Valente Ecotre Valente Srl, Italy

hardening steel and quality of heat treatment is defined by several standards that are mandatory for automotive applications [1,2,3,4].

La Metallurgia Italiana - gennaio 2020

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Scientific papers - Heat Treatments Steel composition of component is defined by standards

Standard size of specimen represents a problem in case

and drawings; small variations in chemical composition

of small parts because it’s not possible to get exact sam-

are allowed. These fluctuations of composition impact on

ple accordingly technical specification.

quenchability of the part and so on its final mechanical

Moreover, Jominy test is a water quench exam [5] while

performance.

gear components are in general quenched with different

To keep under control these variations, a simple but effi-

media and in this way it’s difficult to compare Jominy cur-

cient test to validate mechanical performance of a quen-

ve provided by steel mill certificate with hardness profile

ched and tempered gear in hardening steel is Jominy Test.

obtained on real component.

This methodology requests a specimen with an exact geometry (Fig. 1) and water quenching condition. Quen-

A new approach is necessary to allow proper usage of

ched specimen is tested to get a hardness profile across

Jominy Test results for components hardened with diffe-

thickness which is a curve with hardness value versus di-

rent methods and with geometry not allowing extraction

stance from quenched end.

of standardized sample.

Fig.1 - Jominy Test geometry and slice of a real specimen showing hardness test indentation

Jominy test Jominy test virtualization In order to get a reliable model to forecast performance

Jominy specimen is in steel 17 NiCrMo7 and chemical

of hardened steel gear, FEM software DEFORM-HT has

composition in reported in table 1.

been used to simulate Jominy test [6].

Tab.1 - 17 NiCrMo7 analysis 17 NiCrMo7 composition C

Ni

Cr

Mn

Si

Mo

V

Cu

Al

P

S

Nb

Ti

B

Fe

0.17

1.91

0.53

0.58

0.34

0.26

0.0027

0.16

0.015

0.01

0.03

0.002

0.001

0.0005

Rem.

Initial temperature of sample is 865 °C with full austenite microstructure; applied heat transfer coefficients for test is shown in figure 2 [7,8,9].

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Memorie scientifiche - Trattamenti Termici

Fig.2 - Initial condition and boundary condition applied to DEFORM-HT simulation for Jominy Test

Due to standardization, effects of specimen geometry

because it provides hardness profile result as a function

and quenching can be neglected, which is an advantage

of transformed phase fraction and its hardness. This ap-

of this test. In DEFORM-HT determination of hardness

proach is independent by geometry or quenching system

profile can be approached with 2 different methods: Jo-

because resulting hardness is coming from transforma-

miny Curve from Jominy Distance or Jominy Curve from

tion from austenite “mother phase” into different “chil-

Phase Fraction. Jominy Curve from Jominy Distance is an

dren phases” like martensite or bainite; HRC hardness

approach valid only for Jominy Test; given specimen geo-

is the result of the percentage of each phase at the end

metry and water quenching condition from standard; har-

of heat treatment. DEFORM-HT uses the TTT curves in a

dness profile result is affected only by sample chemical

dynamic way for microstructural calculation; this means

composition and grain size, inclusion and oxides. Coo-

that, despite the use of TTT is always taken into account,

ling rate achieved across thickness is driven only by con-

the kinetics of phase transformations and hence the co-

ductivity and specific heat of steel. This kind of test is very

oling rate have a fundamental role for predicting micro-

easy to carry out but it’s not flexible in case of application

structure. Microstructural transformation, like Austenite

to different quenchant media or geometry. Jominy Curve

→ Pearlite or Austenite → Bainite, is described based on a

from Phase Fraction propose to a more refined approach

TTT Diagram, like shown in Figure 3.

Fig.3 - Kinetics of microstructural transformation defined for DEFORM-HT: Austenite to Pearlite and Austenite to Bainite

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Scientific papers - Heat Treatments Jominy test experimental results Jominy test results on sample in steel 17NiCroMo7 are reported in table 2

Tab.2 - 17NiCrMo7 Hardness profile 17 NiCrMo7 Hardness Profile [HRC] as a function of distance from surface [mm] Distance [mm]

1.5

3.0

5.0

7.0

9.0

11.0

13.0

15.0

20.0

25.0

Hardness [HRC]

44.5

44.0

41.5

37.0

33.0

30.0

29.0

27.0

24.0

23.0

Jominy test simulation results In order to validate reliability of simulation compared to reality, in particular for Fraction of Phase approach, HRC hardness profiles have been plotted in the same graph, shown in figure 4.

Fig.4 - Hardness profile of experimental results, DEFORM-HT simulation results and scatter band for H+ from Standard

Experimental results are inside H+ scatter band impo-

results; a representation of phases distribution at 90 se-

sed by EN Standard. Both Jominy Distance and Volume

conds after quench beginning is shown in Figure 5.

Fraction of phases are well aligned to experimental resul-

In every step “Mother” phase coexists with the other

ts. Jominy Distance result aligned to experimental result

“Children” Phases.

is a confirmation of correct input and correct calculation of DEFORM-HT simulation.

Other information that can be used to monitor heat treatment stage are Time between 800 °C and 500 °C that is

Regarding Volume Fraction of Phases, this result confir-

called also  and it’s an index of cooling rate of material

ms that approaching hardness estimation based on tran-

[7].

sformation of single phases can allow to reach correct

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Memorie scientifiche - Trattamenti Termici

Fig.5 - DEFORM-HT distribution of temperature, Time 8/5, Bainite and Martensite after 90 s from water quench beginning

This method, as itâ&#x20AC;&#x2122;s not based on fixed geometry and

particular for small components that are hardened and

fixed quenching conditions as defined in standards, al-

quenched with different technology.

lows to use simulation with different geometries and in

Application to pinion gear Heat treatment cycle for gear Results obtained from validation of Jominy test has been transferred on a pinion gear in hardening steel 17NiCroMo7, shown in Figure 6.

Fig.6 - Pinion gear simulated with DEFORM-HT.Hardening Heat treatment cycle in DEFORM-HT simulation software is described in table 3

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Scientific papers - Heat Treatments Tab.3 - Hardening heat treatment cycle 17 NiCrMo7 Pinion – Hardening Heat Treatment Cycle Duration [s]

Operation

Temperature [°C]

Carbon concentration [%]

1800

Heating Furnace

550

0

3600

Carburization Furnace

850

0.8

1200

Oil Quench

100

0

1800

Heating Furnace

225

0

3600

Air Cooling

30

0

Results obtained on pinion gear

different slice locations as well as Hardness distribution.

In this simulation, carburization stage introduces an ad-

Both pictures are captured at the end of Oil Quench ope-

ditional variable.

ration.

Carbon content is modified across gear thickness during

This simulation has been run coupling together Thermal

a heating in carburization furnace and this means that it’s

calculation, Microstructural calculation, Stress calcula-

necessary to introduce in TTT curves in DEFORM-HT

tion; as in Jominy Test, for all “Children” phases Thermal,

also dependence from Carbon content. In this way it

Microstructural and Elastic-Plastic properties are defined

will be possible to consider different TTT curve location

as a function of temperature and carbon content to obtain

in Temperature-Time diagram and different Martensite

an accurate solution.

temperatures.

Deformed gear can be exported from simulation in order to be compared with 3D scanning of experimental com-

According this modification in material input data, simu-

ponent.

lation has been run on shown pinion gear.

Simulation results are in alignment with experimental

In Figure 7 some results of simulation are shown; it is visi-

data measured on pinion gear in a range of ± 5%.

ble how Carbon content is various across thickness and at

Fig.7 - Pinion gear simulation with DEFORM-HT: Carbon content and HRC Hardness on two different slices

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Memorie scientifiche - Trattamenti Termici Conclusion DEFORM-HT simulation results are aligned with experi-

Fraction of Phases approach has been transferred on a

mental Jominy Test results.

pinion gear and comparison between experimental and

Simulation results of approach based on Volume

simulation data is in a range of ± 5%.

Fraction of Phases, valid for every type of heat treatment, are aligned with simulation results of Jominy Di-

Finally, DEFORM-HT predicts microstructure, residual

stance approach, valid only for Jominy Test.

austenite, mechanical properties, hardness, residual

According

this

alignment,

DEFORM-HT

Volume

stress and deformations of gear.

Acknowledgements Authors would like to thank M.Sc.Eng. Paolo Martelli for great contribution in development of this work.

References [1]

EN 10083-1, Steels for quenching and tempering - Part 1: General technical delivery conditions

[2]

EN 10083-2, Steels for quenching and tempering - Part 2: Technical delivery conditions for non alloy Steels

[3]

EN 10083-3, Steels for quenching and tempering - Part 3: Technical delivery conditions for alloy steels

[4]

EN 10083-4, Case hardening steels – Part 4: Technical delivery conditions

[5]

ISO 642-79: Steel Hardenability Test by End Quenching (Jominy test)

[6]

DEFORM User Manual, v11.3

[7]

Quenching Theory and Technology - B. Liščić, H.M. Tensi, L.C.F. Canale, G.E. Totten - (second edition) CRC Press

[8]

ASM Metals Handbook – Volume 4 (10th edition – 1990) – Heat Treating

[9]

Measurement of quenching intensity, calculation of heat transfer coefficient and global database of liquid quenchants – B. Liščić, T. Filetin – Materials Engineering 19 (2012) p.52-63

[10]

Prediction of Quench-Hardness within the Whole Volume of Axially Symmetric Workpieces of Any Shape - B. Liščić, S. Singer, B. Smoljan – Journal of ASTM International, Vol.7 No.2 – Paper ID JAI 102647

[11]

Simulation of Heat Treatment of a mold for HPDC: Cooling in gas or salt to get best performance?

A. Magistrelli, D. Mattavelli, P. Martelli, L. Valente, C. Viscardi EUROPEAN CONFERENCE ON HEAT TREATMENT 2015 & 22nd IFHTSE CONGRESS

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Scientific papers - Heat Treatments

Neural networks-based prediction of hardenability of high performance carburizing steels for automotive applications V. Colla, M. Vannucci, L. Bacchi, R. Valentini

The new quenching processes for automotive applications, which follow the cementation stage, include the application of pressurized gas for cooling during quenching. Therefore, it is of utmost importance to have an accurate estimate of the hardenability behavior of carburizing steels, which show a higher Carbon content with respect to traditional materials. These new cooling processes also require properly designed new steels in terms of alloying contents, which ensure a proper response to heat treatment.

In the present paper a neural network-based approach to the prediction of the hardenability profile is proposed, which can be applied both for the design of the steel chemistry and for assessing the suitability of the steel at the steel shop level, in order to suitable adjusting the cooling process after quenching.

KEYWORDS: LOW PRESSURE CARBURIZING - NEURAL NETWORK BASED PREDICTION - JOMINY END QUENCH TEST - TRANSMISSION GEARS

Introduction

face to the core. This is followed by a quenching process

Case hardening is a surface treatment able to confer a

and a low temperature stress relieving tempering.

component elevated resistance to wear, fatigue and in

The cementation can be carried out with different appro-

general surface damage. Moreover, this treatment gives

aches, solid, salts bath and gaseous. Between these, the

at the same time a high toughness at the core, thanks to

gaseous method is the most used because of process

the lower cooling speed, that is why it is commonly used

controllability and speed of the process itself. In this

for parts as transmission gear teeth and diesel fuel injec-

case the surface carburizing is due to carbon dissociation

tors. The process is typically applied to low carbon steels or low alloyed steels as Cr-Ni, Cr-Mo or Mn-Cr. These elements can be added to improve steel hardenability reaching a fully martensitic transformation till the core and thus sufficiently uniform mechanical properties.

Case hardening is a thermochemical diffusion treatment, it consists of a first heating of the component up to austenitizing temperature where the carbon enrichment of the surface is possible and diffusion within the steel is

Valentina Colla, Marco Vannucci TeCIP Institute, Scuola Superiore Santâ&#x20AC;&#x2122;Anna, Italy

Linda Bacchi Letomec s.r.l., Italy

Renzo Valentini DICI; UniversitĂ di Pisa, Italy

enhanced by the elevated carbon gradient from the sur-

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Memorie scientifiche - Trattamenti Termici from a hydrocarbon (propane, acetylene or methane).

provides not only the hardness value at each point but

The innovative variant of the treatment is the low-pres-

also an estimate of the reliability of such prediction.

sure carburizing (LPC), which is carried out at very low

This approach is also exploited within the present work

pressure using hydrocarbon gases and followed by high

for the simulation.

pressure gas quench.

This model was used to optimize gaseous hardening steThis process allows constancy of metallurgical characte-

els chemical composition predicting and comparing Jo-

ristics and lower deformations, surfaces without oxidi-

miny curves. Chemical composition was varied in terms

zing products and no retained austenite. Clean surfaces

of alloying elements and concentration, focusing on Mn-

lead to no need of cleaning products, resulting in lower

Cr and Mn-Cr-B steels.

cost for the treatment (1).

These steels can return a good compromise between mechanical properties and sufficient toughness, but they

Case hardening steels standards are usually very restri-

are a rational choice considering economical reason.

ctive in terms of Jominy curves, as the case of FCA internal material standard number (2).

Materials and methods Low carbon and low alloyed steels are usually used for

The Jominy End Quench Test is very expensive and time

case hardening process, to guarantee the major carbon

consuming thus a lot of models to predict hardenability

gradient between surface and core.

from the steel chemical composition have been proposed.

Considering the low carbon content, some alloy elements are introduced in the chemical composition to

The first attempts were based on traditional statistical

increase the hardenability of the steel reaching a com-

techniques, subsequently, more complex models based

plete martensitic transformation also at the very core of

on Machine Learning (ML) techniques have been em-

the component.

ployed outperforming classical approaches. In particular, Artificial Neural Networks (ANN) have been succes-

For this study, some Jominy tests were industrially per-

sfully applied by providing very satisfactory predictive

formed on various castings, considering the chemical

capabilities. In literature several approaches concerning

ranges of FCA internal standard (Table 1). Jominy End

ANN-based Jominy profile predictors have been propo-

Quench tests were carried out with reference to ASTM

sed (3,6), which predict the steel hardness for each sin-

A255 (11).

gle point of the Jominy profile, while in (7,9) a parametric approach has been proposed and refined, which predicts

Resulting hardness profiles did not match every time

each single point as part of a parametric curve, whose pa-

with desired ranges, in Figure 1 some examples were re-

rameters are correlated to the steel chemistry.

ported and the high deviation of results is evident.

A further relevant advancement is proposed in (10), where a sequantial ANN-based predictor is proposed, which

Tab.1 - Chemical composition ranges (%), balance: Fe C

Mn

Si

Cr

B

Ti

Cu

S

P

Al

0.23÷0.31

1.10÷1.40

0.10÷0.40

1.00÷1.30

≤0.0005

≤0.010

≤0.30

0.025÷ 0.040

≤0.030

0.015÷ 0.040

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Scientific papers - Heat Treatments

Fig.1 - Jominy curves returned on castings with FCA standard chemical composition

A comparison between test results and hardness ranges

ANN-based (refer to paragraph 2). Chemical composi-

reported in FCA internal standards was made. Resulting

tion was varied within new chemical limits in accordance

hardness profile were often quite far from required ones,

with Table 2, in particular Boron was added just to impro-

thus chemical composition was optimized by means of

ve hardenability.

Tab.2 - Modified chemical composition ranges (%): balance: Fe C

Mn

Si

P

S

Cr

Ni

Mo

Cu

Al

V

Nb

Ti

B

0.23÷ 0.31

1.10÷ 1.40

0.10÷ 0.40

≤0.030

0.025÷ 0.040

≤0.40

≤0.40

≤0.10

≤0.30

0.015÷ 0.040

0.003÷ 0.006

0.001÷ 0.003

0.045÷ 0.055

0.0025÷ 0.0035

Neural networks-based prediction of the Jominy profile

prediction of the different points of the profile has been

The Jominy profile prediction is performed by exploi-

performed both on the basis of some literature results

ting a set of 2-layers standard perceptron-based Fe-

related to the influence of chemical elements on the Jo-

ed-Forward ANNs according to the hierarchical structu-

miny profile and of a preliminary statistical analysis ai-

re, which is depicted in Figure 2.

ming at highlighting the correlation between the poten-

The inputs and the number Nh of neurons in the hidden

tial input variables and the target hardness values(12,13).

layer are different for each ANN and include the most correlated chemical elements and eventually the previously predicted hardness of other points of the Jominy profile.

The selection of the input variables to be used for the

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Memorie scientifiche - Trattamenti Termici

Fig.2 -Sequential structure of the Jominy profile predictor Table 3 shows in detail the list of the inputs which are

points of the Jominy profile are predicted in sequence

fed to each neural network for the prediction of specific

in order to suitably exploit the relations between them.

Jominy profile points: each row of the table corresponds

Due to the dependencies between single points expres-

to one point and reports both the included chemical ele-

sed in Table1, the order of prediction is not the natural

ments and the predicted hardness values fed as input to

one, but the following order, which is more fruitful for

the ANN (where Ĵ i represents the hardness value estima-

the purpose of this work, is adopted: 1-2-3-4-5-14-15-

ted at i-th distance value). Table 1 also reports the num-

13-12-11-10-9-8-7-6, where each number represents

ber of neurons in the hidden layer N h of each ANN. The

the corresponding profile point.

Tab.3 - Inputs of the ANNs of the Jominy profile predictor Point

Input

Nh

Chemical elements

Predicted hardness values

J1

C, Mn, Si

-

2

J2

C,P,S,Cu

Ĵ1

2

J3

C, Cr, S, Ni

Ĵ1, Ĵ2

3

J4

C, Cu, Mn, Mo

Ĵ2, Ĵ3

3

J5

C, B, Ti, Nb

Ĵ3, Ĵ4

4

C, B, Ti, Nb

Ĵ2, Ĵ5, Ĵ7

4

C, B, Ti, Nb

Ĵ1, Ĵ2, Ĵ5, Ĵ10

4

C, Sn, Al, V

Ĵ4, Ĵ6, Ĵ7, Ĵ9, Ĵ10

4

C, S, P

Ĵ2, Ĵ3, Ĵ5, Ĵ11

4

C, S, P, Cu

Ĵ1, Ĵ2, Ĵ4, Ĵ5, Ĵ11

3

J11

C, S, Cr, Ni

Ĵ1, Ĵ4, Ĵ5, Ĵ12, Ĵ15

3

J12

C, Mn, Ni

Ĵ1, Ĵ3, Ĵ4, Ĵ13

3

J13

C, S, Cr, Ni, Nb,Ti, B

Ĵ1, Ĵ4, Ĵ5

3

J14

C, P, S, Cr

Ĵ1, Ĵ2, Ĵ3, Ĵ5

3

J15

C, P, S, Cr

Ĵ1, Ĵ3, Ĵ5

2

J6 J7 J8 J9 J10

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Scientific papers - Heat Treatments The ANNs are sequentially trained by means of the stan-

its general validity.

dard back-propagation algorithm by exploiting 600 randomly selected samples from an experimental dataset

Results

composed of about 1100 samples, while the other sam-

Simulation results were analyzed to evaluate the good-

ples are used for validation. This dataset was available

ness of resulting Jominy curves in comparison with desi-

before the development of the analysis described in this

red hardness defined in the FCA standard.

paper, namely none of the data which have been used in

For each Jominy position the error was calculated com-

the present work have been exploited in the training of

pared to the average value of the specification limits,

the system. For the purpose of the analysis discussed

consequently the average error, calculated on the whole

here, the ANN-based Jominy predictor has been used

curve, was estimated.

simply as a â&#x20AC;&#x153;virtual hardenability test developerâ&#x20AC;? and the

Thanks to this analysis a series of optimal and promising

goodness of the achieved results further demonstrates

chemical compositions were defined.

Fig.3 - ANN-based prediction of the Jominy curves: best results

To validate results also a series of Jominy tests were per-

Figure 3, and the comparison of calculated errors was

formed in plant. As an example, some of the curves with

reported in Table 4.

the lower average error were reported in Figure 2 and

Tab.4 - Average error estimation for Jominy hardness profiles compared to FCA reference standard Test ID

Average error (HRc) Neural network-based prediction

Average error (HRc) Industrial tests

2

3.8

2.9

12

3.1

2.5

42

3.0

3.2

43

3.6

2.8

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Memorie scientifiche - Trattamenti Termici

Fig.4 - Jominy results on casting with optimal chemical composition predicted by the ANN-based predictor

Conclusion

tor could be a powerful tool to avoid a large number of

The present study was a preliminary investigation with

experimental industrial tests.

the aim of optimizing production process for suitable

3) The error that occurs with the ANN-based prediction

steels for low pressure cementation. Considering the

has the same order of magnitude of experimental data di-

first results the following consideration can be gathered:

spersion on this category of steels.

1) Usual chemical composition ranges are too wide to

4) The presented results lead to the development of ste-

obtain components in accordance with severe automo-

el for the specific application as a subgroup of MnCrB5

tive reference standards, that is why a specialization is

series. These steels have the certain advantage of a gre-

required.

at economical convenience thanks to the absence of

2) A properly trained neural network-based predic-

expensive alloy elements as Nickel or Molybdenum.

References [1]

Morgano E. Trattamenti Termici Innovativi e Processi Meccanici Superficiali. Seminary at University of Pisa. 2017.

[2]

Fiat Group Automobiles standard 52414. Acciai da Cementazione Acciai al Manganese – Cromo. 2010.

[3]

Vermeulen W.G., Van Der Wolk P.J., De Weijer A.P., Van der Zwaag S. Prediction of Jominy hardness profiles of steels using Artificial Neural Networks J. Mater. Engineering, 1996, 57–63.

[4]

Dobrazanski L.A., Sitek W. Application of neural network in modelling of hardenability of constructional steels. Jour. of Mat. Proc. Technology, 1998, 78, 59–66.

[5]

Dobrazanski L.A., Sitek W. The modelling of hardenability using neural networks. Journal of Materials Processing Technology.

1999, 92–93, 8–14.

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Scientific papers - Heat Treatments [6]

Chan B., Bibby M., Holtz N. Predicting hardness with artificial neural networks. Canadian Metallurgic Quaterly. 1995, 34, 353–356.

[7]

Colla V., Reyneri L.M., Sgarbi M. Parametric characterization of hardness profiles of steels with neuro-wavelet networks. Lecture Notes in Computer Science, 1999, 1607, 606–614.

[8]

Colla V., Reyneri L.M., Sgarbi M. Neuro-wavelet parametric characterization of Jominy profiles of steels, Integrated Computer-Aided Engineering, 2000, 7(3), 217–228.

[9]

Valentini R., Colla V., Sgarbi M., Reyneri L.M. Parametric Jominy profiles predictor based on neural networks. Revista de metalurgia, 2005, SPEC. VOL, 314–317.

[10]

Cateni S., Colla V., Vannucci M., Vannocci M. Prediction of steel hardenability and related reliability through neural networks Proc. IASTED Int. Conf. on Artificial Intelligence and Applications, AIA 2013, 2013. 169–174.

[11]

ASTM A255-10. Standard Test Method for Determining Hardenability of Steel. 2018.

[12]

Cateni, S., Colla, V., Vannucci, M. General purpose input variables extraction: A genetic algorithm-based procedure GIVE a GAP, ISDA 2009 - 9th Int. Conf. on Intelligent Systems Design and Applications. 2009, 1278–1283.

[13]

Cateni, S., Colla, V., Vannucci, M. A genetic algorithm-based approach for selecting input variables and setting relevant network parameters of a SOM-based classifier, International Journal of Simulation: Systems, Science and Technology. 2011, 12(2), 30–37.

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Memorie scientifiche - Trattamenti Termici

Heat treatment routes utilizing intercritical annealing of automotive medium-mn steel sheets A. Grajcar, M. Morawiec

Low-carbon medium-Mn steel with an increased aluminium content was subjected to different heat treatments using dilatometry simulations. Samples of the 0.16C-5Mn-1.5Al steel were intercritically annealed between 700 and 800°C for duration of 1 hour. These heat treatment routes were designed to enrich the austenite in carbon to stabilize retained austenite after cooling the samples to room temperature. The austenite decomposition was monitored using the dilatometry. The martensite start temperatures were calculated and determined experimentally when the Ms was higher than the room temperature. The microstructures were investigated using light and scanning electron microscopy techniques. It was found that a heat treatment temperature affects substantially a thermal stability of the austenite and other microstructural features of the multiphase microstructures. The effect of C partitioning during the intercritical annealing on the austenite stability is discussed.

KEYWORDS: INTERCRITICAL ANNEALLING - MEDIUM-MANGANESE STEEL RETAINED AUSTENITE - TRANSFORMATION KINETICS - DILATOMETRY

Introduction

steels the main route for their manufacturing is cold rol-

A new generation of steels for the automotive industry

ling after which a simple heat treatment takes place [5-

is undergoing a dynamic development because of high

8]. A typical Mn content in the cold-rolled medium-Mn

demands on various environmental, safety and fuel eco-

steels is from 7 to ca. 11%. From the industry point of

nomy aspects.

view it increases the price of finished sheet products.

These demands coming from different regulations put a

Moreover, there are various technological problems

pressure on the automotive and steel industries to ma-

related to hot-dip galvanizing, instable flow behavior,

nufacture a wide range of steels with a good combination

weldability, etc.

of mechanical properties [1-4]. The special interest is in advanced high strength steels characterized by a multiphase microstructure. Recently, the focus is on medium-manganese steels, which include from 3 to 12% Mn. The idea for these steels is to obtain a good combination of strength and ductility, together with a relatively low price related to the chemical composition and

Adam Grajcar, Mateusz Morawiec Silesian University of Technology,

Faculty of Mechanical Engineering,

Institute of Engineering Materials and Biomaterials, Gliwice, Poland

adam.grajcar@polsl.pl

manufacturing process. In case of medium-manganese

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Scientific papers - Heat Treatments Far less attention is paid to the hot rolling process for

steel is as follows: 0.16% C, 4.7% Mn, 1.6% Al, 0.2% Mo,

this grade of steels [8-11]. One can assume that the man-

0.2% Si.

ganese addition together with carbon should increase

The steel was hot rolled and air cooled. The initial mi-

the stability of austenite and enhance the TRIP (TRan-

crostructure is martensitic due to the strong hardenabi-

sformation Induced Plasticity) effect. However, some

lity effect of manganese.

authors reported [12-14] that the increase of Mn content

The intercritical annealing was applied for hot rolled and

to the chemical composition can result in an opposite

machined samples.

effect.

The heat treatment was performed by the means of di-

The reason is that the Mn addition decreases the diffu-

latometry, which at the same time gives the information

sional enrichment of the austenite in carbon (a thermo-

on a phase transformation kinetics upon heating and

dynamic effect).

subsequent cooling.

In hot rolled steels it is also very important to know how

The samples in a shape of rods of 5 mm diameter and 10

manganese influences the austenite decomposition ki-

mm length were heated up to the desired temperatures

netics.

in a range from 700 to 800°C at a heating rate of 3°C/s,

From our preliminary tests this is clear that manganese

held for duration of 1 hour and cooled to room tempe-

shifts the ferritic and bainitic regions to very long times.

rature at the rate of 60°C/s. This heat treatment approach

In case of the presented lean-Mn steel obtaining the bai-

is presented in Fig. 1.

nitic transformation during isothermal holding at diffe-

The whole process was carried out in vacuum. After

rent temperatures was impossible even after 5 hours.

the heat treatment the samples were cut in the middle

This means that the redistribution of C to the austenite

part of the rod for the preparation for light and scanning

can not take place. Hence, the different approach of heat

electron microscopy.

treatment for this kind of hot-rolled steel was applied in

The samples were grinded with the use of 220, 500, 800

the present study.

and 1200 grinding papers and polished with the use of 3 and 1 µm diamond solution.

Experimental procedure

Next, the samples were etched using 5% Nital. The Vi-

The chemical composition of the analyzed medium-Mn

ckers hardness was measured with the load of 9.81 N.

Fig.1 - The heat treatment routes applied for the investigated steel

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Memorie scientifiche - Trattamenti Termici Results and discussion

temperature for the samples intercritically annealed at

The first step of the analysis was the dilatometry.

700°C.

The dilatometric results present how the phase transfor-

This is due to the different state of the austenite at the

mation kinetics changes at different intercritical annea-

intercritical region expressed by the lever rule.

ling temperatures.

At 800°C the microstructure contains a large austenite

In case of the presented results, the change of marten-

fraction but with the smallest carbon content.

sitic start temperature indicates the different stability of

Thus, the austenite does not have the sufficient thermal

austenite upon cooling to room temperature. First of all,

stability and transforms to martensite during a cooling

the martensite start temperature of the steel was calcu-

step.

lated using the mucg83 software developed by Bhade-

When the annealing temperature decreases the amount

shia [15].

of the austenite decreases because a larger volume

The calculated temperature was equal to 335°C.

fraction of ferrite is formed from the initial martensite.

This was the same value as determined by the experi-

Hence, the martensite start temperature decreases to

mental approach.

190 and 105°C, respectively for the samples held at 760

Next the dilatometric analysis was carried out to deter-

and 720°C.

mine the influence of intercritical annealing temperature

This is because the lower intercritical temperatures al-

on the M s temperature.

low for the higher carbon content to be enriched in the

The change in M s temperature indicated a change in the

remaining austenite.

stability of the austenite.

This took place due to an enrichment of the austenite in carbon during ferrite formation in the intercritical region. These dilatometric results are presented in Fig. 2.

It can be seen that the M s temperature decreases gradually together with the decrease of the intercritical temperature from 800 to 700°C. For example, the M s = 245°C at 800°C and is below room

Fig.2 - Dilatometric analysis during cooling to room temperature after intercritical annealing at 700-800°C

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Scientific papers - Heat Treatments

Fig.3 - LM microstructures after intercritical annealing at: a) 800°C, b) 760°C, c) 720°C, d) 700°C

The sufficient enrichment of the austenite in carbon took place only at the temperature of 700°C. For this treatment there was no martensite start in the dilatometric curve (Fig. 2).

Fig.4 - SEM microstructures after intercritical annealing at: a) 800°C, b) 760°C, c) 720°C, d) 700°C; F - ferrite, M - martensite, RA - retained austenite

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Memorie scientifiche - Trattamenti Termici After the dilatometric tests, the microstructure investi-

The most important phase present in the microstructure

gation was carried out. The microstructures from the li-

is retained austenite, which fraction increases together

ght microscopy (LM) are presented in Fig. 3. It is clear

with decreasing the intercritical temperature. At 700°C

that the microstructures inherit the martensite morpho-

there is no fresh martensite in the sample (Fig. 4d). Hen-

logy. It is difficult to assess if the martensite transformed

ce, this is an optimal heat treatment temperature.

into austenite. The difference in a ferrite content can

The last step of the analysis was hardness measuremen-

only be visible in these LM micrographs. It is difficult to

ts. Six tests were carried out on each sample, after whi-

separate ferrite at 800°C (Fig. 3a). This phase is clearly

ch the average value was calculated. The results of the

seen at 760°C as thin layers between martensite laths

hardness analysis are presented in Fig. 5. The hardness

(Fig. 3b). At 720 and 700°C the amount of the ferrite is

increases at the higher annealing temperatures. This is

larger. Moreover, some ferrite grains are more granular.

in good correlation with the dilatometric and micro-

The higher fraction of the ferrite allowed to enrich the

structure investigations. At 800°C, the highest hardness

austenite in carbon.

was measured. This corresponds to the large fraction

A higher magnification in Fig. 4 allowed for more detai-

of fresh martensite formed between 245 and 20°C. The

led analysis of the microstructure. It can be seen that the

fresh martensite formed during cooling does not under-

presented laths of ferrite are very thin; some of them be-

go any tempering. This is expressed as its highest hard-

low 100 nm in width. The width of the austenite and ferri-

ness (ca. 460 HV). When the temperature of the annea-

te decreases as the intercritical temperature is reduced.

ling decreases the amount of the intercritical austenite

The separation of individual laths of austenite and ferrite

and resulting fresh martensite decrease. Moreover, the

is also stronger at the lower intercritical temperatures.

amount of retained austenite increases with the decre-

Moreover, the martensite morphology is not the same

asing temperature. At 700°C no fresh martensite is for-

in all samples. The various volume fraction of fresh mar-

med during cooling to room temperature, as the dilato-

tensite can be observed at different temperatures. Ac-

metry test shows (Fig. 2).

cording to the dilatometric analysis in a 720-800°C range

Hence, the lamellar mixture of the highest volume

some poor-enriched carbon austenite transforms into

fraction of ferrite and austenite shows the lowest hard-

martensite during cooling (it is called fresh martensite).

ness (ca. 310 V).

Fig.5 - Hardness for different intercritical annealing temperatures

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Scientific papers - Heat Treatments Conclusions

gest ferrite fraction and resulting carbon enrichment of

Presented results showed that intercritical annealing is

the austenite. At higher temperatures triple ferrite-au-

a simple heat treatment, which can be applied for lean

stenite-fresh martensite microstructures were formed.

medium-Mn Al-alloyed steels. The dilatometric analy-

The martensite start temperatures were above room

sis indicated that it is possible to stabilize the austeni-

temperature being the reason for the fresh martensite

te during the intercritical annealing through its carbon

formation. The hardness was the highest for the higher

enrichment from the formed ferrite. The microstructu-

annealing temperature, i.e., for the corresponding lar-

res consist of different portions of lamellar ferrite, fresh

gest volume fraction of fresh martensite. The lowest

martensite and retained austenite. The ferrite fraction in-

hardness showed the sample with the duplex ferrite-au-

creased as the interctitical temperature decreased from

stenite microstructure, with the highest potential for the

800 to 700°C.

automotive sheet steel application.

The beneficial very thin duplex ferrite and austenite mi-

Further studies are needed to assess a volume fraction

crostructure was only obtained at 700°C due to the lar-

of retained austenite and its carbon content.

Acknowledgement The financial support of the National Science Center, Poland, is gratefully acknowledged, grant no. 2017/27/B/ST8/02864.

References [1]

Lehnhoff GR, Findley KO, De Cooman BC. The influence of silicon and aluminium alloying on the lattice parameter and stacking fault energy of austenitic steel. Scripta Materialia 2014;92:19-22.

[2]

Zhao JZ, Mesplont C, De Cooman B.C. Quantitative analysis of the dilatation during an isothermal decomposition of austenites.

Materials Science and Engineering A 2002;332:110-6.

[3]

Steineder K, Krizan D, Schneider R, Beal C, Sommitsch C. On the microstructural characteristics influencing the yielding behavior of ultra-fine grained medium-Mn steels. Acta Materialia 2017;139:39-50.

[4]

Kang S, Speer JG, Krizan D, Matlock DK, De Moor E. Prediction of tensile properties of intercritically annealed Al-containing 0,19C-4.5Mn (wt%) TRIP steels. Materials and Design 2016;97:138-46.

[5]

Dutta A, Ponge D, Sandlobes S, Raabe D. Understanding hot vs. cold rolled medium mangenese steel deformation behavior using in situ micoscopic digital image correlation. Materials Science Forum 2018;941:198; doi:10.4028/www.scientific.net/ MSF.941.198.

[6]

Benzing JT, Kwiatkowski da Silva A, Morsdorf L, Bentley J, Ponge D, Dutta A, Han J, McBride JR, Van Leer B, Gault B, Raabe D, Witting JE. Multi-scale characterization of austenite reversion and martensite recovery in cold-rolled medium-Mn steel. Acta Materialia 2019;166:512-30.

[7]

Cheng P, Hu B, Liu SL, Guo H, Enomoto M, Shang CJ. Influence of retained austenite and Cu precipitates on the mechanical properties of a cold-rolled and intercritically annealed medium Mn steel. Materials Science and Engineering A 2019;746:41-9.

[8]

Liu C, Peng Q, Xue Z, Wang S, Yang C. Microstructure and mechanical properties of hot-rolled and cold-rolled medium-Mn trip steels. Materials 2018;11:2242; doi:10.3390/ma11112242.

La Metallurgia Italiana - January 2020

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Memorie scientifiche - Trattamenti Termici [9]

Timokhina IB, Hodgson PD, Pereloma EV. Effect of microstructure on the stability of retained austenie in transformation-induced-plasticity steels. Metallurgical and Materials Transactions A 2004;35A:2332-41.

[10]

Liu C, Peng Q, Xue Z, Deng M, Wang S, Yang C. Microstructure-tensile properties relationship and austenite stability of a Nb-Mo micro-alloyed medium-Mn trip steel. Metals 2018;8:615; doi:10.3390/met8080615.

[11]

Qi X, Du L, Hu J, Misra RDK. Effect of austenite stability on toughness, ductility and work-hardening of medium-Mn steel. Materials Science and Technology 2018; doi:10.1080/02670836.2018.1522088.

[12]

Grajcar A, Kilarski A, Kozłowska A, Radwański K. Microstructure evolution and mechanical stability of retained austenite in thermomechanically processed medium-Mn steel. Materials 2019;12:501; doi:10.3390/ma12030501.

[13]

Grajcar A, Kilarski A, Kozłowska A. Microstructure-property relationships in thermomechanically processed medium-Mn steels with high Al content. Metals 2018;8:929; doi:10.3390/met8110929.

[14]

Haupt M, Dutta A, Ponge D, Sandlobes S, Nellessen M, Hirt G. Influence of intercritical annealing on microstructure and mechanical properties of a medium manganese steel. Procedia Engineering 2017;207:1803-8.

[15]

Bhadeshia HKDH. Thermodynamic analysis of isothermal transformation diagrams. Metal Science 1982;16:159-65.

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Industry news - Heat Treatments

New environmental-friendly developments and applications of the pulsed plasma heat treatment technology in the field of engine, transmission and power train engineering A. Gebeshuber, V. Strobl, T. Müller

Nowadays nitrided gears have been successfully introduced for applications such as steering gears in engines, internal gears in transmissions and other not very highly loaded applications. In future, the fast growing market for top performance gears in the vehicle industry with its high load demands, is a particularly promising expansion area. Especially here is a need for reliable fatigue data from optimized material-plasmanitriding-combinations. To generate such fatigue data a step by step approach with specially designed specimens and test procedures has been chosen. This paper discusses present results and gives an outlook on the steps ahead. The first part of this investigation focuses on bending fatigue strength of the gear tooth, while other failure modes, such as pitting, will be tested in the second part of the project.

KEYWORDS: PLASMA NITRIDING - FATIGUE DATA - GEAR DESIGN RESIDUAL COMPRESSIVE STRESSES - DISTORTION

Introduction

has induced the development of new technologies in

Over the last 15 years, there has been a constant rise in

classical carburizing, such as low pressure carburizing

the amount of hard machining - mainly grinding of tooth

with gas quenching. But all these technologies require

flanks - required after carburizing. At the transmission factory of VW Kassel for example, there was an increase from 2 % in 1998 to as much as 14 % in 2005 [1]. The cost of such post heat treatment machining may thus amount up to 40% of total unit costs.

The gear industry approaches this problem with diffe-

Andreas Gebeshuber, Volker Strobl, Thomas Müller Rübig GmbH&Co.KG, Wels, Austria

rent strategies. The cost saving pressure in gear production does not only affect the machining processes, but also calls for new possibilities of heat treatment, and

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Attualità industriale - Trattamenti Termici high temperatures and martensitic transformation as a

additional new tools and machining methods have been

metallurgical mechanism in order to increase both har-

developed, so that the availability of high strength steels

dness and fatigue life.

suitable for nitriding and the production process invol-

Both factors - high temperatures combined with fast

ved are no longer constraints upon nitriding in the gear

quenching, and volume change by martensitic transfor-

industry.

mation - lead to considerable distortion and therefore

Because of the intensive hard machining after carburi-

to an inevitable machining operation when high profile

zing required for high performance gears, and even ta-

and surface quality and narrow geometric margins need

king into account the higher alloyed steels, soft machi-

to be met.

ning of high strength steels and longer process times in nitriding, considerable savings in the overall costs can

Nitriding is a hardening process which does not use

still be achieved if the hard machining after carburizing

martensitic transformation, but precipitation hardening

is no longer necessary and can thus be dropped.

as a mechanism.

Furthermore, nitriding is carried out at fairly low tem-

Investigation procedure

peratures in the range of 500 to 600°C, while fast quen-

This paper presents some results of a systematic ap-

ching in a fluid medium is not necessary.

proach to generate reliable data on plasmanitrided high

Plasma nitriding is a special form of nitriding, where the

strength steels for gear and drive application.

necessary free nitrogen is generated by a physical process - the low pressure plasma, whereas the dissocia-

Potential for plasmanitriding gears

tion of cyanates in saltbath nitriding, or of ammonium in

The potential of plasmanitriding regarding load capacity

the gas nitriding process, are chemical reactions. This

of highly stressed gears can be seen from Schlöterm-

difference in generating the free nitrogen has several ad-

ann´s investigations [2].

ditional advantages, such as

Extensive fatigue tests with gas or plasmanitrided gears

• environmental friendliness (green process)

have been conducted with different steels and material

• easy integration into production line

conditions and have been compared with the fatigue

• high accuracy and process stability

data according to DIN 3990 [3].

• low gas consumption

The conclusion of the tests is that the norm values are

• possibility of easy partial nitriding by mechani-

very low in comparison to the test results.

cal masking.

A test to compare case hardened and nitrided gears was The low distortion of the nitriding process in general

conducted by H. Mallener [4].

and the above mentioned additional advantages make

This showed that nitrided gears present reduced tooth

the plasma nitriding process one of the most promising

root strength, especially in the low cycle area.

alternatives in heat treatment for high performance ge-

Another investigation concerning the load carrying ca-

ars, when low overall costs, high quality in production,

pacity of nitrided steel gears by L. Albertin and H. Winter

integration into the machining floors and pollution pre-

conducted at the Gear Research Center (FZG, Munich)

vention are an issue.

has produced similar results [5].

However, nitriding has never been used on a broader

A recent research project about bending fatigue stren-

range in the production of highly stressed gears.

gth of case hardened and nitrided gears by Prof. P. Da-

The reason is not only the nitriding process itself, but

voli [6] showed a different result. His concluding sta-

also the higher costs of the alloyed steel required, and

tement: “The damage line of the two nitriding steels is

the need of machining of teeth up to a strength level of

quite high, not far from the Wöhler line; which implies

1500 MPa and even more. In recent years, new steels and

a good overload capacity for these two nitriding steels.

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Industry news - Heat Treatments As a final statement, the good bending load carrying ca-

core strength, surface hardness, white layer thickness,

pacity of properly heat treated nitriding steels has been

nitriding depth and residual compressive stress profiles

confirmed by the tests.“

were applied to four different high strength steels in the

By matching material selection with proper heat treat-

course of plasmanitrided processes especially develo-

ment and optimized plasmanitriding, it should be pos-

ped for such applications.

sible to obtain good fatigue values in the low cycle area as well.

According to DIN 3990-5 the tooth bending fatigue strength ranges from 270 up to 550 MPa (Fig. 1), depen-

Results

ding on material and applied surface treatment (nitriding

The scope of the investigations was the determination of

or case hardening).

fatigue behaviour of different high strength plasmanitri-

Schlötermann already showed in 1988 [2], that nitrided

ded steels.

steel can yield higher values than case-hardened steel -

During the research program, different combinations of

in contradiction to the values acc. to DIN 3990-5.

Fig.1 -Bending fatigue strength: comparison of literature data and development goal RÜBIG

This was the motivation for Rübig to start further inve-

were compared with those of case hardened specimen.

stigations with the goal of narrowing the scatter band by rising the lower band (Fig. 1).

The results show, that proper combination of steel and plasma-nitriding parameters result in bending fatigue

In a first step, during which the influence of material

values that are at least equivalent to case-hardened ste-

choice and plasma nitriding was to be studied, all fatigue

el (see Fig. 2).

tests were performed at FZG in Munich on gears with a modulus of 2,75.

So it was possible to determine the tooth bending fatigue of a typical gear tooth geometry. The fatigue data

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Attualità industriale - Trattamenti Termici

Fig.2 -Bending fatigue strength: comparison of results RÜBIG and literature data [8]

From our investigations it is obvious that the residual

in the diffusion zone.

compressive stress profile is the dominating factor on fatigue life in all cases where surface hardness compri-

The next steps in the project are:

ses a range between 700 and 1000 HV and core strength

• Conduct fatigue tests on high performance plasmani-

between 1000 to 1600 MPa.

trided gears to evaluate flank bearing properties compared to case-hardened gears (in progress)

Just as with all the other surface hardening processes,

• Evaluation of the economic benefit due to shorter

such as carburiz-ing or induction hardening, it is the ma-

manufacturing process (pic. 3) of plasmanitrided gears

ximum value of compressive stress in the surface area

compared to case-hardened gears (in progress

and the depth of the stress profile that influences the fatigue limit positively.

The latter step is of great interest since the plasmanitriding process features very little distortion due to the

Generally, the fatigue data for nitrided steels were in the

hardening mechanism by precipitation hardening.

same range as for comparable carburized samples, with high compressive residual stresses the fatigue data are

Interim results show, that there is a high potential for the

even exceeding the values for carburized sam-ples.

substitution of case-hardening by plasmanitriding and can lead to a reduction of total costs for gear manufa-

So far we could not establish any difference in fatigue

tur-ing.:

life values that would relate to the microstructure of the white layer or to the microstructure in the diffusion zone.

This may be due to the fact that all samples were, just as for the data reported in literature, nitrided with little or no white layer and with hardly any visible precipi-tation

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Industry news - Heat Treatments

Fig.3 -Manufacturing process of gears: “case hardening” vs. “plasmanitriding”

We are convinced that upon completion of this com-

rial-plasmanitriding combinations to match maximum

prehensive investigation, we will be able to give ba-

gear load ca-pacity for a wide range of standard and spe-

sic rec-ommendations for choosing optimized mate-

cial gear applications.

References

[1]

Weißner, R., Fenstermann, J.: Anforderungen, Stand der Technik und Perspektiven von Verzahnungen im Automobilbau. Dresden; Germany: 2. Deutsch-Italienische Near-Net-Shape-Tagung, 2005.

[2]

Schlötermann, K.: Auslegung nitrierter Zahnradgetriebe, Untersuchungen zu den Auswirkungen unterschiedlicher Nitrierparameter auf den Werkstoffzustand und die Tragfähigkeit von Zahnrädern. Dissertation, RWTH Aachen, 1988.

[3]

Norm DIN 3990 Teil 5 1984: Tragfähigkeitsberechnung von Stirnrädern, Dauerfestigkeitswerte und Werkstoffqualitäten

[4]

Mallener, H., Schulz, M.: Nitrieren von Zahnrädern als Alternative für das Einsatzhärten?. HTM Z . Werkst. Wärmebeh. Fertigung 48, (1993) 3, pp. 166 – 171

[5]

Albertin, L., Frolich, R. L., Winter, H.& Höhn B.-R., Michaelis, K.: Load Carrying Capacity of Nitrided Gears. AGMA FTM; 94FTM4, 1993

[6]

Davoli, P., Boniardi, M., Longoni, C.: Bending Fatigue Resistance of Case Hardened and Nitrided Gears. München; Germany: International Conference on Gears, 2005

[7]

Trubitz, P., Kaminsky, T., Brusky, U., Kaiser, B., Adelmann, J.: Schwingfestigkeit nitrierter Proben. Vorhaben Nr.160, Lebensdauervorhersage nitrierter bauteilähnlicher Proben mit Hilfe normierter Wöhlerstreubänder – Einfluss der Randschicht. Forschungsheft FKM; Heft 197, 1995

[8]

Stenico, A.: „Werkstoffmechanische Untersuchungen zur Zahnfußtragfähigkeit einsatzgehärteter Zahnräder“, Dissertation, TU München, 2007

La Metallurgia Italiana - January 2020

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AttualitĂ industriale - Trattamenti Termici

Liquid Nitrocarburizing treatment on valves in internal combustion engines M. Legrand, V. Monteux

In a strong position because of more than 60 yearsâ&#x20AC;&#x2122; experience in the area of thermochemical treatments, HEF Group, world leader in liquid ionic nitrocarburizing processes, has developed and optimized in that field his own treatments family. They are commercialized under the name CLIN (Controlled Liquid Ionic Nitriding) keeping as standard a commune philosophy: offering a solution in continuous development to support its customers with highest quality and robustness, still in advance in front of environment constraints.

Thus, liquid nitrocaburizing became and remains an essential surface treatment solution, globally established, well known and recognized as only-one efficient solution for specific automotive components such as valves.

On that specific component, the new environmental standards for a better engine efficiency and the latest constraints on surfaces technologies such as hard chromium coating, lead to new developments and designs for moving parts in internal combustion engines. These activities are the challenges of engine manufacturers.

Liquid nitrocaburizing treatment is now considered as industrial mature technology of treatment of new automotive valves design, with pre-localized induction hardening, in ferrous-chromium alloys and / or nickel base alloy known for their high resistance to thermal oxidation.

Furthermore, this treatment is nowadays required as a technically and economically competitive alternative to hard chrome plating treatment. The latter is strongly penalized by the restrictions imposed by the new environmental laws (REACH ...) and risk eventually being no longer competitive (high cost price).

KEYWORDS: NITRIDING TREATMENT - SALT BATH - LIQUID IONIC - ENGINE EFFICIENCY - FATIGUE RESISTANCE - ALTERNATIVE TO CHROMIUM PLATING - VALVE

Introduction In the recent years, automotive industries are facing a complete revolution in combustion engine design to face high constraints on their efficiency - CO2 emission reduction - Downsizing,

M. Legrand, V. Monteux HEF Durferrit, Andrezieux Boutheon, France

- Hybrid version development.

To withstand these constraints, the valves designs

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Industry news - Heat Treatments have been reviewed taking into consideration

- Higher stress on tip due to repetitive contact with

- Highest temperature into chamber (the valve seat

cam, tappets, follower…. Induction hardening of the

must be able to resist the corrosive effects of hot

tip become standard and obvious.

exhaust gases from the engine). New material have been successfully tested such as high alloyed iron base material or nickel base material.

INLET VALVE • • •

Larger head diameter Lower operating temperature than exhaust (500-800°C) Material choice - Low carbon martensitic steel - X45 Cr Si 9 3 (9% Cr) can be hardened for durability on tip

EXHAUST VALVE • • • •

Higher operating temperature (800-1200°C) Corrosive environment Lubrication issue Material choice - usually bi metallic valves - X45 Cr Si 9 3 (9% Cr) (can be hardened for durability on tip) on tip and stem side - Stainless steel 21% Cr on head, - Or Stainless 30% Cr 15% Ni on head - Or nickel base material on head.

Meanwhile, their standard surfaces treatments techno-

experience in the area of thermochemical treatments,

logies, such as chromium plating, used for decade are

HEF Group, world leader in liquid ionic nitro carburizing

suffering of new environmental pressure, inviting them

processes, has developed his own new formulation cal-

to consider alternatives in their new challenges.

led ARCOR® DT.

To succeed in the replacement of chromium plating on valves made of high alloyed elements, nitriding te-

This process reaches unexpected results without any

chnologies players have to adjust temperature of their

pre acid etching, with a high efficiency (short treatment

nitriding atmosphere in the lowest side of their opera-

time although low temperature), homogeneity, repro-

ting range (approx. 500-540°C to limit relieving of the

ducibility, robustness (whatever the material is) and cost

pre-localized induction hardening), keeping efficiency

efficiency (limited consumables ratio on total operating

in reduction of passive surfaces on such material (to

cost by ensuring low consumption, although low tem-

allow homogeneous and quick growing of the nitride

perature).

layers).

ARCOR® DT is a low temperature nitro carburizing proARCOR DT: new formulation development

cess, operating in a temperature range (500°C to 540°C)

In a strong position because of more than 60 years’

lower than tip induction hardening temperature. Hence

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Attualità industriale - Trattamenti Termici re-hardening after nitriding is no more required contrary

Results and conclusion

to some others options than TUFFTRIDE, TENIFER or

ARCOR DT consists of the well-known standard se-

others nitro carburizing technologies.

quence: - Alkaline degreasing and rinsing operations,

Its efficiency comes from its special chemistry offering

- Preheating into air circulated oven

lower melting point (hence higher fluidity. Furthermore,

- Nitriding at low temperature (500°C to 540°C) for a tre-

its optimized oxo-basicity offers high nitriding rate even

atment time lower than 3 hours

at lower temperatures and even on high alloyed steel

- Cooling in industrial water,

and nickel base material.

- Washing in industrial water - Drying,

Finally, following advantages have to be pointed out

- Mechanical finishing such as blasting using thin metal-

- Cleaner valves after nitriding

lic powder.

- requiring only low pressure shot blasting after nitriding - ideally suited for hollow-stem valves, thin-stem and

In the following table, metallographic observation un-

valves with low ‘stem-bowing’ tolerance.

der optical microscopy are being given on different ma-

- greener technology.

terials.

Material

Metallographic observation under optical microscopy

Applications

for monometallic inlet valve and on tip side on bi metallic exhaust valve

X 45 Cr Si 9 3 (QT)

Nital etching, magnification x500

(9%Mn ; 4,5%Ni ; 21%Cr ; Fe balance) Used on seat side for bi metallic exhaust valve

X 50 CrMnNiNbN 21 9

Marble etching, magnification x1000

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Industry news - Heat Treatments

(30%Ni, 15%Cr, Fe balance) Used on seat side for bi metallic exhaust valve

ECMS 3015

Nital etching, magnification x500

(30%Cr, 3%Fe, Ni balance) Used on seat side for bi metallic exhaust valve

Nimonic 80A

Marble etching, magnification x500

As a conclusion, the new formulation of Liquid Nitrocar-

- Liquid nitrocarburizing can be operated at a temperature,

burizing ARCOR DT becomes quickly the relevant solu-

which allow keeping the hardness, brought by induction

tion, showing the highest advantages for treatment on new

hardening, without acting as a stress relieving and without

generation valves:

negative consequence on produc-tion costs (limited consumables consumption; short treatment time).

- Liquid nitrocaburizing treatment improves the resistance to wear and the anti-friction qualities by in-creasing the surface hardness along the valve stem, while retaining the corrosion resistance of the material at the valve seat.

- Liquid nitrocarburizing is able to treat simultaneously the different material even ones with higher al-loyed element contents.

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AttualitĂ industriale - Trattamenti Termici

Non-destructive hardness / microstructure testing of heat-treated parts by mass production, with multiple frequency magnetic induction method A. Horsch

Increasing quality demands, new productliability regulations, as well as international market networks force manufacturers to take special measures encompassing thefield of material testing. Nowadays, specified tolerances are extremely small. Therefore processes have to include the conducting of 100% material, structure and hardening tests on a fast, reliable and simple basis. The technology applied must be the latest state of art, it must comply with maximum safety requirements and be economical.

Introduction

stations with 8 standard test frequencies and up to 16

Today, every company carries out spot checks during

test positions have significantly increased the efficiency

manufacturing, in order to guarantee the quality of pro-

of this test method.

ducts. Within the framework of control one of the most

The field of application ranges from testing the micro-

important problems is to ensure repeatability of quality.

structure, hardening and case depth to material min-

Quality control faces major problems within this area;

gle. One of the outstanding advantages of modern eddy

these might start with supply of the material or products.

current testing systems using PMFT (Preventive Multi

Parts properties; that is structure, hardness or case dep-

Frequency Testing) is the ability to detect unexpected

th, which are determined by the heat treatment process,

faults.

are subject to variations. To ensure the correct heat treatment results, a fast and efficient test method is nee-

Magnetic test methods have been used since the be-

ded, it should be non-destructive.

ginning of the 1940s. Springer [1] and Maercks [2] have

It must be clear that heat treatment is not a standard sta-

already made their publications in 1941. Figures 1 and 2

tistical process, each heat treatment batch is a separa-

show the permeability measuring device he uses and a

te production that is not related to the previous one. In

magnetic coil sorting device.

a continuous belt furnace a quasi-statistical process is given, but here too a spot check can only provide limited information about the entire production. The heat treatment company provides only a process which can be verified by spot check. This does not say whether all

A. Horsch Arnold Horsch e.K., Remscheid, Germany

parts also comply the specification.

Computer based multi-frequency eddy current testing

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Industry news - Heat Treatments

Fig.1 - Magnetic test system from Springer, laboratory setup [1]

Fig.2 - Magnetic test system from Springer, industrial test system [1]

Both describe successful sorting tests for the sorting

checks in order to conclude the general nature of the

of differences in hardness of high-speed steel and har-

quality from the test data. Table 1 shows typical tested

dness depth differences after case hardening on gears.

parts [3].

WHY 100% TEST? Variation in product usually is in accord with statistical analysis and can be predicted or estimated. For this reason, it is sufficient to make a certain number of spot

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AttualitĂ industriale - Trattamenti Termici Tab.1

Application Examples of Material Testing by Means of Eddy Current Preventive Multi-Frequency Testing (PMFT) Component f. e.

Parameters to be tested

Bearing parts

case depth, structure and hardness pattern

Axle components

case depth, structure and hardness pattern

Pinion pins and axles

case depth, structure and hardness pattern

Linear guide components

case depth, structure and hardness pattern

Steering racks and the like

case depth, structure and hardness pattern

Gears

case depth, structure and hardness pattern

Cam Shafts

case depth, structure and hardness pattern

Bolts, Screws, Nuts

Structure, hardness, decarburisation

and much more

Problems develop if during manufacturing some occur-

unexpectedly. Table 2 lists possible hardening faults [3].

rences arise which are not subject to standard statistical

We also have to keep in mind that materials are not per-

distribution. First, one needs to know what may happen

fect and can have flaws nobody expected.

Tab.2

Possible hardening errors during Heat Treatment Which parameter was incorrect?

Austenitizing temperature

Austenitizing time

Quenching

In which way was it incorrect?

What are the effects of this?

too high

Overhardening, incorrect structure martensite + residual austenite

too low

Underhardening, incorrect structure martensite + bainite + ferrite

too long

Overhardening, case too high, incorrect structure martensite + residual austenite

too short

Underhardening, shallow case, incorrect structure martensite + bainite + ferrite

too fast

Incorrect structure, cracks martensite + residual austenite

too slow

Incorrect structure martensite + bainite + ferrite

formation of vapor bubbles too low

soft spots not defined Hardness too high

too short -

hardness too low

too long

hardness too high

Tempering time

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Industry news - Heat Treatments

too slow

shallow case, misplaced case, austenitizing time too short

too fast

case too high, misplaced case, austenitizing time too long

Rate of feed by induction hardening

Damaged inductor

undefined

undefined

Mal-positioning by induction hardening

undefined

unsymmetrical hardening pattern, overheating, melting

How does eddy current testing works [4]

the transport from the hardening station to the test sta-

Eddy current testing work with coils, generators,

tion takes only a few seconds.

ac-current and ac-voltage, frequencies, field strength

Faulty parts, are immediately detected and separated,

and induction law. Eddy current test examine them for

providing an enormous savings in time and costs!

their metallurgical microstructure, thus for their me-

The energy for eddy current testing is very small, in the

chanical features like hardness, case depth or alloy.

milliwatt range. Field strength is low and permeability is

The eddy current test does not provide absolute values

in the range of the initial permeability. Test frequencies

(e.g. “56 HRC” or “2.6 mm case depth”). The eddy cur-

ranging from some Hz to some hundreds of kHz provide

rent test does detect fine differences in micro-structu-

information on undesired structures via the frequency

re with high sensitivity. In the production line, within a

dependent penetration depth of the eddy current and on

fraction of a second, a non-destructive 100 % test for

the formation of permeability. Very small electrical si-

micro structure

gnals require a very precise evaluation in order to assure

• case depth

their differentiation from ambient interferences. A small

• hardness run out,

drift from variation in temperature and high long term

• hardness pattern

stability are absolutely necessary. Digitization of input

• tensile strength

voltages immediately at the front end of the electronic

• carbon content

evaluation is of huge advantage.

• soft spots

Relative permeability [µr] is strongly affected by heat

• surface decarburisation

treatment. The amount of carbon and other alloy elements significantly influence the size and course of

is completed, and thus quick corrective reactions to any

the permeability curve. Generally, hardened (stressed)

variance from the specified structure can be realized.

structures have lower permeability than soft structures

With a suitable mechanical part handling arrangement,

(refer to graph of 0.78% C), see Figure 3.

Fig.3 - Relative permeability of different materials

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Attualità industriale - Trattamenti Termici Electrical conductivity [σ] is influenced only a little by

it. It is a „comparative test“. Values of OK parts (referen-

structure changes and alloy differences. Whereas, the

ce parts) which were presented to and stored in the in-

conductivity of carbon steel is < 10MS/m high-alloyed

strument beforehand are compared with the values of

chrome-nickel is about 1.3 MS/m. However, the tem-

currently produced parts.

perature coefficient is about 4 to 5 % per 10°C. Thus,

How can one get values? The alternating current i which

the temperature of the test part has influence on the test

flows through the red coil created a magnetic flux [B]

result and should be allowed to vary only negligibly (±

through the test part (grey in the sketch). The size of the

5°C).

magnetic flux and thus the size of the voltage u indu-

How can magnetic and electrical changes that correlate

ced in the blue coil is directly dependant on the electri-

with the mechanical properties be quickly and reliably

cal conductivity [σ] and on the magnetic „conductivity“

tested non-destructively?

(permeability [µr ]) of the test part in the coil, Figure 4.

Eddy current testing has proven to be well qualified for

Fig.4

The test part by ist σ and µr strongly influences the cou-

real faults never look like artificially created faults, fur-

pling between sender coil (red) and receiver coil (blue).

thermore only this one specially calibrated fault can be

Thus voltages induced in the receiver coil imply the

found. Another fault will not be found and go into pro-

structure, e.g. to verify correct hardness, case depth,

duction and will be installed e.g. in an engine.

core hardness as well as alloy. This complex signal is di-

An expected error will not be sorted out, the processes

splayed twodimensionally as vector (complex number).

must be adjusted so that this error cannot occurs. Figure 5 shows the possible distributions of OK. parts, artificial

Check against known faults

NOK. parts, real NOK parts and unexpected fault parts,

First of all, we have to determine how faults occur in the

clearly showed the not all faults will be sorted out by the

heat treatment process. In principle, they do not arise

method with expected artificial faults. The real NOK par-

as expected faults, but they always arise as unexpected

ts lift a larger spread than the artificially manufactured

faults. If an eddy current test system is set to an expected

parts, that means not all NOK parts will be sorted out.

fault, these setting master must be manufactured under

The unexpected fault will not be found.

controlled conditions. This means that the unexpected

La Metallurgia Italiana - gennaio 2020

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Industry news - Heat Treatments

Fig.5

If permeability is tested with a single magnetic field

What occurs if another material, e.g. X40Cr13 is mixed

strength HS, test results may be unreliable as soon as

up unexpectedly and tested? The example in

other types of mixed parts become involved. Let us take

Figure 6 b shows that the permeability curve of C45

a typical mixed part test of two kinds of steel: C45 and

crosses the curve of X40Cr13 at HS (red circle).

23NiCrMo2. The largest difference in permeability is at

A differentiation with one field strength (frequency) is

HS. This test can be done with a single frequency test

not possible. Both remain mixed up in spite of the eddy

instrument set up with that field strength, because this

current test [4]. That means there is no separation.

setting provides the largest difference between both, Figure 6a.

Fig.6a

Fig.6b

Check against unknown faults

ced worldwide to test all kinds of material, including the

If we want to sort against an unexpected fault, we need

determining of microstructure and hardening characte-

a completely different approach to the testing task. We

ristics. There are numerous applications for such a sy-

need a technique which -

stem.

• has the ability, to calibrate only with OK parts • a broad spectrum of tests and many test frequencies

How it works?

• fast enough

The use of several test frequencies to induce different

• provides maximum warranty to find unexpected errors

field strengths (H1-8) in the coil means that all these dif-

Eddy current testing, especially Preventive Multi Fre-

ferences are detectable or viewable. Different structures

quency Testing (PMFT) offer this and is applicable for

from different heat treat methods (and by heat treat pro-

a large range of uses. It has been successfully introdu-

cess faults) create different permeability curves which

La Metallurgia Italiana - January 2020

pagina 87

Attualità industriale - Trattamenti Termici are detected. Thus the eddy current testing becomes

The test take place at different field strength, so to say

reliable, expected and unexpected wrong structures are

“preventively” over a larger area.

detected. Different to Figure 6a+b, Figure 7 shows the “curtains” be moved aside in order to widen the view.

Fig.7

If several frequencies of alternating current are used, a

ty OK parts only are needed to setup the instrument and

voltage vector is obtained for each frequency in the im-

to form the tolerance zones (calibration).

pedance, and there is a locus curve. Of course, other

A challenge test with NOK parts (e.g. not-hardened, in-

factors also influence the induced voltage.

correctly quenched, austenizing temperature not rea-

The position of the test part in the coil, geometric va-

ched, too short or too long tempering, annealing tem-

riances in the test part and other factors alter the recei-

perature too high or too low, etc.) can be done, but is

ved voltages.

not needed.

These factors must therefore be controlled. Neverthe-

The test system will reliably detect faulty parts with both

less, received voltage values of several OK parts will

known and unknown defects.

always vary slightly.

The multi frequency test method works reliably for all

They are subject to scattering. The vector tips of the vol-

kinds of defects which may happen during heat treat-

tages form a cloud (refer to the green dots at 4 kHz). If

ment of steel It works like a fingerprint comparison, only

the vector tips are now enclosed with an elliptic toleran-

when all parameters are 100% fit, a part will be recogni-

ce zone, the test can be reduced to a comparison of the

zed as OK, Figure 8 b [5].

vectors to inside (OK) or outside (NOK) of the tolerance zone.

Testing current production with such a frequency band (eight frequencies), Figure 8 a, and the comparison with the previously created tolerance zones made with good parts has become well known in professional circles as Preventive Multi Frequency Test (PMFT).

A quite wide frequency band (at least 1:1000) is used in order to detect all abnormal structures detectable by eddy current as faulty and to sort them out. Ten to twen-

La Metallurgia Italiana - gennaio 2020

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Industry news - Heat Treatments

Fig.8a

Fig.8b

Only if all tolerance fields are matched part will be OK, if

included during calibration blows up the automatically

only one tolerance field is not matched, the part is sor-

formed tolerance zones, and the instrumentâ&#x20AC;&#x2122;s sensitivi-

ted as NOK. [5]

ty is deteriorated. Vectors of bad parts must be deleted

Careful choice of OK parts is a precondition to reliable

before the actual test is started.

testing. One or several NOK parts (red dot), Figure 9,

Fig.9

Practicle examples Used test system Modern test equipment, which are fulfilling all the described requirements are few. Only one test equipment meets the requirements of Preventive Multi Frequency Testing â&#x20AC;&#x201C; PMFT perfectly, Figure 10 a + b shows this equipment.

Fig.10a

La Metallurgia Italiana - January 2020

Fig.10b

pagina 89

Attualità industriale - Trattamenti Termici Inspection of induction hardend cam shafts

Checked for:

Forged camshafts are inductively hardened in the series.

• case depth

In this example, the camshafts are automatically tested

• structure

in a test machine in the production line Figure 11.

• hardness pattern • material mix

Fig.11

Inspection of screws

In this system, with a turntable, with newly develo-

Screws and other fasteners are extremely important hi-

ping U-coils, screws can be checked very sensitivly for

gh-security components throughout different industrys

unexpected various faults, as:

and construction industries. The failure of a screw can lead

• strength

to production stoppages or serious accidents. In the faste-

• structural analysis

ner industry is the permissible error quantity less than 1

• carburizing

ppm.

• decarburizing

For this reason, more and more testing machines will

• material mix

be equipped with Preventive Multi Frequency Testing -

• aso.

PMFT. Figures 12 a + b shows such a high speed testing

Such testing technology can be adapted at any time to exi-

system for screws, the test speed is appr. 10-15 parts/sec..

sting testing machines.

La Metallurgia Italiana - gennaio 2020

pagina 90

Industry news - Heat Treatments

Fig.12a

Fig.12b

Inspection of balls

• different heat treatments in one material

Balls ,Figure 13 a, are used everywhere in the industry,

• test speed 17 parts/sec.

not just in ball bearings. Here balls of 0.6 - 4.5mm are

must be sorted against each other, there is no mingling

tested in a flexible system.

allowed. Figure 13 b -+ c, shows the realised machine,

The goal is not to deliver wrong balls.

with 4 test channels, each test channel works with 8 test

The challenge was:

frequencies.

• 28 different materials

Fig.12a

Fig.12b

Fig.12a

Inspection Of Distance Pins

Frequency Testing – PMFT.

The typical fault of this part is a wrong hardness pattern,

The fault by Part 1 is a classical decarburization and the

Figure 14, against this fault the parts are tested.

fault by Part 2 is first also a classical decarburization, fol-

The Material is a Carbonsteel wit 0,35 % C content, tole-

lowed by a subsequent faulty carburization to reversed the

rance SHD 0,5+1mm, ≥50HRC.

decarburization.

During the series test parts were found which failed during the eddy current test.

A macroscopic test with etching did not reveal any errors.

A microscopic test together with a hardness depth measurement gave the hardness profiles shown in Fig. 15 and 17.

Part 1, Figure 16, had a decarburisation and Part 2, Figure 18, a ferrite space within the hardness zone.

This bug was certainly found with the Preventive Multi

La Metallurgia Italiana - January 2020

pagina 91

AttualitĂ industriale - Trattamenti Termici

Fig.14

Fig.15 - Hardness trace with decarburizing

Fig.16 - Bainite, Martensite and Ferrite 500:1

Fig.17 - Hardness trace with Ferrite space under surface

Fig.18 - Bainite, Martensite and Ferrite space under surface 25:1

La Metallurgia Italiana - gennaio 2020

pagina 92

Industry news - Heat Treatments Inspection of special bolts

terial,

This component, Figure 19 requires uniform hardening

• batch 1, 275-285 HB, microstructure Figure 20

in a very narrow tolerance band. The drawing tolerance is

• batch 2, 315-330 HB, microstructure Figure 21

260-340 HB, but all parts of a hardening batch must be in a

both batches are in the drawing tolerance but the spread

tolerance of 20HB.

in the heat treat batch is too high. With Preventive Multi

A high risk is a material-batch mixing that leads to different

Frequency Testing – PMFT, established a save test system

hardnesses within the drawing tolerance. In this example,

which made sure that no batch mixing or other faults are

we have two mixed material batches from the same ma-

in production.

Fig.19

Fig.20 - 275-285 HB 500:1

Fig.21 - 315-330 HB 500:1

La Metallurgia Italiana - January 2020

pagina 93

Attualità industriale - Trattamenti Termici Summary

heat treatment defects / material mingling. Thus, the

The test against unknown heat treatment defects / ma-

eddy current test is perfectly suited for a 100% test for

terial mingling by means of preventive “PMFT” eddy

medium and large quantities in order to ensure pro-

current test gives the heat treatment department a safe

duction quality.

test method with the highest possible safety. Due to the

The older testing technique of a single frequency eddy

integration in production lines, very low test costs are

current test against expected errors with artificially ma-

possible, with enormous security.

nufactured adjusting parts is not a reliable test to find

In addition to high test sensitivity for microstructural

heat treatment defects / material mingling, at all.

changes, hardness faults and case depth changes, PMFT provides a very high level of test safety for unexpected

References [1]

Dipl. Ing. H Springer, Über eine neue Praktische Möglichkeit, die Gefügeänderung bei der Wärmebehandlung von Werkzeugstählen messtechnisch zu erfassen, Härterei Technische Mitteilungen, 1941, Band 1, Seite 61-74, Germany

[2]

Dr. O. Maercks, Magnetische Prüfverfahren zur Werkstoffkontrolle und zur Einsatztiefenmessung, Härterei Technische Mitteilungen, 1941, Band 1, Seite 156-165, Germany

[3]

A. Horsch, U. Bardelmeier, Einsatz eines Wirbelstromprüfgerates mit Mehrfrequenztechnik in einer Lohnhärterei, Hanser Verlag, Harterei Technische Mittelungen 49, 1994, Seite 237-244, Germany

[4]

Herbert Baumgartner, ibg NDT-Group, Eddy Current Test, Company publication, 04.Dec.2012

[5]

A. Horsch, 100% Eddy Current Test for hardness testing and material mix, 26th Hungary Heat Treat Forum, 8-10 October 2014, Balatonfüred, Hungary

La Metallurgia Italiana - gennaio 2020

pagina 94

Atti e notizie - AIM news

Eventi AIM / AIM events

WINTER SCHOOL OF METALLURGY “METAL ADDITIVE MANUFACTURING” Corso > Trento, 26-29 gennaio TECNOLOGIA ED INNOVAZIONE NEI FORNI AD ARCO Giornata di Studio > Dalmine c/o ABB, 24 febbraio TENACITÀ E FATICA Corso > febbraio

www.aimnet.it

ANALISI CHIMICHE Corso > febbraio/marzo

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

LEGHE PER ALTA TEMPERATURA PRODOTTE CON TECNOLOGIE ADDITIVE Giornata di Studio > Firenze, c/o Baker Hughes, 17 marzo IL MONDO INDUSTRIAL: ASPETTI METALLURGICI E METODOLOGIE DI CONTROLLO Giornata di Studio > Torino c/o CNH Industrial Village, 19 marzo

MACCHINA FUSORIA Corso Itinerante > 22-23-28-29 aprile - 6-7 maggio RACCORDERIA E PRESSATURA. GLI ASPETTI METALLURGICI Giornata di Studio > Suzzara c/o Intertraco Italia, aprile FAILURE ANALYSIS 11a ed.

Corso > primavera-estate LO SVILUPPO TECNOLOGICO DEGLI IMPIANTI DI PRODUZIONE A CALDO DELLE VERGELLE ALTO CARBONIO E ALTO LEGATE E PER FILO SALDANTE Giornata di Studio > Lecco c/o Caleotto, primavera 27° CONVEGNO NAZIONALE TRATTAMENTI TERMICI Genova, 6-7 maggio ECCC 2020 - 10TH EUROPEAN CONTINUOUS CASTING CONFERENCE Bari, 17-19 giugno RIVESTIMENTI DECORATIVI AL SERVIZIO DELL’ESTETICA DEL PRODOTTO Giornata di Studio > Firenze, 25 giugno STAMPAGGIO: PROPRIETÀ MECCANICHE, TRATTAMENTO TERMICO E MECCANISMO DI DANNEGGIAMENTI Giornata di Studio > Ivrea c/o Confindustria Canavese, giugno PRESS & SINTER: EVOLUZIONE DI PRODOTTI, MATERIALI E TECNOLOGIE NEL SETTORE AUTOMOTIVE Giornata di Studio > Brescia, giugno HTDC 2020 - 7TH INTERNATIONAL CONFERENCE HIGH TECH DIE CASTING Vicenza, 1-3 luglio 38° CONVEGNO NAZIONALE AIM 16-18 settembre 2020, Napoli TRATTAMENTI TERMICI E MODELLAZIONE Giornata di Studio > Brescia, settembre ACCIAI AD ALTO TENORE DI CARBONIO Corso > settembre PRESSOCOLATI IN ZAMA Giornata di Studio > Brescia, autunno MATERIALI INNOVATIVI PER L’ADDITIVE MANUFACTURING Giornata di Studio > Torino c/o FCA, novembre RIVESTIMENTI - III MODULO: RIVESTIMENTI PER VIA UMIDA Corso modulare > novembre 1ST INTERNATIONAL MEETING RAW MATERIALS & RECYCLING Bergamo, 3-4 dicembre

La Metallurgia Italiana - January 2020

pagina 95

Atti e notizie - AIM news

Corso // Le Prove non Distruttive Milano, 20-21 novembre 2019

Il 20 e 21 novembre scorsi l’Associazione Italiana di Metallurgia -

(SMT) sulle prove non tradizionali in campo, Luca Bertolini (libero

AIM - ha organizzato a Milano, al Centro Congressi FAST, la quarta

professionista) su getti e fucinati, Stefano Biagioni (Trenitalia) sui

edizione del proprio corso sui controlli non distruttivi, sponsoriz-

materiali rotabili.

zato da Förster e TICAM e patrocinato da AIPnD.

Giuseppe Torsello, ex CISE, oggi TICAM, ha spiegato la possibi-

Hanno preso parte all’attività formativa circa 60 partecipanti (metà

lità di fare le prove meccaniche non distruttive. Ad arricchimento

dall’Università e metà dall’Industria) con background molto diver-

della lezione frontale, TICAM ha esposto per entrambe le giornate

si, dai completamente digiuni di PND ai pluricertificati di livello 3.

attrezzature e campioni da dimostrazione.

Per i primi il sottoscritto ha tenuto una lezione introduttiva di 45

Infine Daniele Bisi (CRF) ha affascinato l’uditorio narrando filoso-

minuti in cui ha raccontato la filosofia delle prove ed i principi, le

fia ed operatività dei CND nell’automobile ed a completamento

finalità, i vantaggi e le limitazioni dei metodi più diffusi, per tutti gli

Bernardo Barile (ArcelorMittal) ha illustrato la straordinaria capa-

altri i relatori hanno illustrato poi lo stato dell’arte dei metodi più

cità dell’industria siderurgica di raccogliere in produzione un’in-

innovativi e dei più importanti campi di applicazione.

finità di dati sulle lamiere da stampaggio e metterli a disposizione

Sui metodi sono intervenuti Umberto Papponetti (2P) con una panoramica sugli ultrasuoni di ultima generazione, Marco Castagna

dell’utilizzatore in tutte le fasi della successiva lavorazione, grazie ad una identificazione in continuo delle lamiere medesime.

(Gilardoni) sui controlli radioscopici dell’era digitale, Rainer Sailer

Il corso si è concluso con i saluti finali del sottoscritto e di una

(Förster) con una chiarissima esposizione in inglese sulle novità

folta rappresentanza del direttivo AIPND, a partire dal presidente

delle correnti indotte e Stefano Benuzzi (Tec Eurolab) sugli stu-

Ezio Tuberosa, al vicepresidente Daniele Bisi, ai consiglieri Fede-

pefacenti risultati della tomografia.

rico Raggio ed Oliviero Oldani.

Dopo uno stimolante intervento accademico del prof. Carboni (Politecnico di Milano) su efficacia diagnostica e difetto minimo rilevabile, si è passati ai settori di applicazione, dove storici docenti come Marco Casaril (OMECO) sulle saldature, Cosimo Carnovale (TenarisDalmine) sui tubi senza saldatura, Mario Cusolito (AIM) sui prodotti lunghi e Ezio Trentini (OMECO) su gru, carriponte e funi, si sono avvicendati a giovani specialisti come Igor Giroletti

La Metallurgia Italiana - gennaio 2020

Con la passione che li contraddistingue ciascuno di loro ha lanciato un messaggio molto significativo: le PND hanno bisogno di gente seria, preparata e generosa ed il mondo dell’industria e delle costruzioni in genere ha bisogno di loro per garantire la sicurezza dei lavoratori, degli utenti e della popolazione in generale.

Ezio Trentini

Coordinatore del corso assieme a Mario Cusolito

pagina 96

Atti e notizie - AIM news

Normativa / Standards in fabbrica con tubi di servizio in acciaio,

UNI CEN ISO/TS 35105:2019

AIM - UNSIDER - Norme pubblicate e

isolamento termico in poliuretano e guaina

Industrie del petrolio e del gas naturale -

progetti in inchiesta (aggiornamento

esterna in polietilene.

Operazioni artiche - Requisiti dei materiali per le operazioni artiche.

30 novembre 2019) UNI EN 253:2019

Norme pubblicate e progetti allo stu-

Tubi per teleriscaldamento - Sistemi bloc-

UNI EN ISO 35106:2019

dio (elenco)

cati monotubo per reti di acqua calda inter-

Industrie del petrolio e del gas naturale -

rate direttamente - Produzione di tubi di

Operazioni artiche - Dati oceano-meteoro-

Norme UNSIDER pubblicate da UNI

servizio in acciaio, isolamento termico in

logici, sul ghiaccio e sui fondali marini.

nel mese di novembre 2019

poliuretano e guaina esterna in polietilene.

UNI EN ISO 35101:2019

UNI EN ISO 3183:2019

UNI EN 1753:2019

Industrie del petrolio e del gas naturale -

Industrie del petrolio e del gas naturale -

Magnesio e leghe di magnesio - Lingotti e

Operazioni artiche - Ambiente di lavoro.

Tubi di acciaio per i sistemi di trasporto per

getti di leghe di magnesio.

UNI EN ISO 19905-3:2019

mezzo di condotte.

UNI EN 17248:2019

Industrie del petrolio e del gas naturale -

UNI EN ISO 20074:2019

Sistemi di tubazioni per teleriscaldamento

Valutazione specifica del sito di unitĂ mo-

Industrie del petrolio e del gas naturale -

e raffrescamento - Termini e definizioni.

bili in mare aperto - Parte 3: UnitĂ galleggiante.

Sistemi di trasporto mediante condotte - Gestione dei rischi geologici per le con-

UNI EN 15698-2:2019

dotte terrestri.

Tubi per teleriscaldamento - Sistemi bloc-

UNI EN ISO 18647:2019

cati di tubi doppi per reti di acqua calda in-

Industrie del petrolio e del gas naturale -

UNI EN 489-1:2019

terrate - Parte 2: Raccordi e valvole isolati

Impianti di perforazione modulari per piat-

Tubazioni per teleriscaldamento - Sistemi

termicamente in fabbrica con tubi di servi-

taforme fisse in mare aperto.

bloccati di tubi singoli e doppi per reti di

zio in acciaio, isolamento termico in poliu-

acqua calda interrate direttamente - Parte

retano e una guaina esterna in polietilene.

UNI EN ISO 10893-3:2019 Controlli non distruttivi dei tubi di acciaio -

1: Ripristino del giunto esterno ed isolamento termico per reti di acqua calda se-

UNI EN 15698-1:2019

Parte 3: Controllo automatizzato mediante

condo EN 13941-1.

Tubi per teleriscaldamento - Sistemi bloc-

flusso disperso sullâ&#x20AC;&#x2122;intera superficie di tubi

cati di tubi doppi per reti di acqua calda

di acciaio ferromagnetico, senza saldatura

UNI EN 488:2019

interrate - Parte 1: Assemblaggio di tubi

e saldati (eccetto quelli ad arco sommer-

Tubi per teleriscaldamento - Sistemi bloc-

doppi isolati termicamente in fabbrica con

so), per la rilevazione di imperfezioni lon-

cati monotubo per reti di acqua calda in-

tubi di servizio in acciaio, isolamento ter-

gitudinali e/o trasversali.

terrate direttamente - Valvole in acciaio

mico in poliuretano e una guaina esterna in

isolate termicamente in fabbrica per tubi di

polietilene.

Industrie del petrolio e del gas naturale -

servizio in acciaio, isolamento termico in poliuretano e guaina esterna in polietilene.

UNI EN ISO 35103:2019

UNI EN 14419:2019

Operazioni in clima artico - Monitoraggio

Tubi per teleriscaldamento - Sistemi bloc-

ambientale.

UNI EN 448:2019

cati con tubi singoli o doppi per reti di ac-

Tubi per teleriscaldamento - Sistemi bloc-

qua calda interrate direttamente - Sistemi

UNI EN ISO 19903:2019

cati monotubo per reti di acqua calda in-

di sorveglianza.

Industrie del petrolio e del gas naturale -

terrate direttamente - Raccordi realizzati

La Metallurgia Italiana - January 2020

Strutture di calcestruzzo in mare aperto.

pagina 97

Atti e notizie - AIM news UNI EN ISO 10426-3:2019

flusso disperso sullâ&#x20AC;&#x2122;intera superficie di tubi

UNI EN 1753:2002

Industrie del petrolio e del gas naturale -

di acciaio ferromagnetico, senza saldatura

Magnesio e leghe di magnesio - Lingotti e

Cementi e materiali per la cementazione

e saldati (eccetto quelli ad arco sommer-

getti di leghe di magnesio.

dei pozzi - Parte 3: Prove della formulazio-

so), per la rilevazione di imperfezioni lon-

ne dei cementi per pozzi in acque profon-

gitudinali e/o trasversali.

de.

UNI 8810:1987 Analisi chimica dei materiali ferrosi.

UNI EN 15698-1:2009

Determinazione dello stagno negli acciai.

Norme UNSIDER ritirate da UNI

Tubazioni per teleriscaldamento - Sistemi

Metodo per spettrofotometria di assorbi-

nel mese di novembre 2019

bloccati a doppio tubo preisolati per reti di

mento atomico nella fiamma.

acqua calda interrate direttamente - Parte UNI EN ISO 3183:2018

1: Assemblaggio di tubi doppi per mezzo

UNI 8808:1987

Industrie del petrolio e del gas naturale -

di tubi di servizio di acciaio, isolamento

Analisi chimica dei materiali ferrosi.

Tubi di acciaio per i sistemi di trasporto per

termico in poliuretano e tubi di protezione

Determinazione del molibdeno negli ac-

mezzo di condotte.

esterna di polietilene.

ciai. Metodo per spettrofotometria di assorbimento atomico nella fiamma.

UNI EN 488:2016

UNI EN 489:2009

Tubazioni per teleriscaldamento - Sistemi

Tubazioni per teleriscaldamento - Sistemi

UNI 8567:1984

bloccati di tubazioni preisolate per reti di

bloccati di tubazioni preisolate per reti di

Analisi chimica dei materiali ferrosi.

acqua calda interrate direttamente - As-

acqua calda interrate direttamente - As-

Determinazione del tungsteno negli acciai.

semblaggio di valvole per tubi di servizio

semblaggio della giunzione per tubi di ser-

Metodo per spettrometria di assorbimento

di acciaio con isolamento termico di po-

vizio di acciaio con isolamento termico di

atomico nella fiamma.

liuretano e tubo di protezione esterna di

poliuretano e tubo esterno di polietilene.

polietilene.

UNI 8566:1984 UNI EN 14419:2009

Analisi chimica dei materiali ferrosi.

UNI EN 253:2016

Tubazioni per teleriscaldamento - Sistemi

Determinazione del cobalto negli acciai e

Tubazioni per teleriscaldamento - Sistemi

bloccati di tubazioni preisolate per reti di

nelle ghise. Metodo per spettrofotometria

bloccati di tubazioni preisolate per reti di

acqua calda interrate direttamente - Siste-

di assorbimento atomico nella fiamma.

acqua calda interrate direttamente - As-

mi di sorveglianza.

semblaggio di tubi di servizio di acciaio,

UNI 8363:1982

isolamento termico a base di poliuretano

UNI EN ISO 19903:2007

Analisi chimica dei materiali ferrosi.

e tubi di protezione esterna di polietilene.

Industrie del petrolio e del gas naturale -

Determinazione del calcio negli acciai.

Strutture in mare fisse di calcestruzzo.

Metodo per spettrofotometria di assorbi-

UNI EN 15698-2:2015

mento atomico nella fiamma.

Tubazioni per teleriscaldamento - Sistemi

UNI EN ISO 10426-3:2005

bloccati a doppio tubo preisolati per reti di

Industrie del petrolio e del gas naturale -

UNI 7842:1978

acqua calda interrate direttamente - Parte

Cementi e materiali per la cementazione

Analisi chimica dei materiali ferrosi.

2: Assemblaggio di raccordi e valvole per

dei pozzi - Parte 3: Prove della formulazio-

Determinazione dello zinco negli acciai.

tubi di servizio di acciaio, isolamento ter-

ne dei cementi per pozzi in acque profon-

Metodo per spettrofotometria di assorbi-

mico in poliuretano e tubi di protezione

de.

mento atomico nella fiamma.

EC 1-2004 UNI EN 1753:2002

UNI 7726:1977

UNI EN ISO 10893-3:2011

Magnesio e leghe di magnesio - Lingotti e

Analisi chimica dei materiali ferrosi. Deter-

Controlli non distruttivi dei tubi di acciaio -

getti di leghe di magnesio.

minazione dellâ&#x20AC;&#x2122; antimonio negli acciai. Me-

esterna in polietilene.

Parte 3: Controllo automatizzato mediante

La Metallurgia Italiana - gennaio 2020

todo spettrofotometrico al verde brillante.

pagina 98

Atti e notizie - AIM news UNI 7715:1977

UNI 7330:1974

UNI 6651:1970

Analisi chimica dei materiali ferrosi. Deter-

Analisi chimica dei materiali ferrosi. Deter-

Analisi chimica dei minerali di manganese.

minazione del piombo negli acciai inossi-

minazione del rame negli acciai e nelle ghi-

Determinazione del rame. Metodo colori-

dabili. Metodo per spettrofotometria di as-

se. Metodo fotometrico.

metrico.

UNI 7089:1972

UNI 6650:1970

UNI 7714:1977

Analisi chimica dei materiali ferrosi. De-

Analisi chimica dei minerali di manganese.

Analisi chimica dei materiali ferrosi.

terminazione del rame negli acciai e nelle

Determinazione del rame. Metodo gravi-

Determinazione del piombo negli acciai

ghise. Metodo elettrolitico a potenziale

metrico.

non legati. Metodo per spettrofotome-

controllato.

sorbimento atomico nella fiamma.

tria di assorbimento atomico nella fiamma (metodo previa estrazione del ferro).

UNI 6649:1970 UNI 6708:1970

Analisi chimica dei minerali di manganese.

Analisi chimica dei minerali di manganese.

Determinazione del fosforo. Metodo colo-

Determinazione del titanio.

rimetrico.

Determinazione del cobalto negli acciai.

UNI 6707:1970

UNI 6648:1970

Metodo potenziometrico.

Analisi chimica dei minerali di manganese.

Analisi chimica dei minerali di manganese.

Determinazione dell’ ossido di magnesio

Determinazione del fosforo. Metodo volu-

UNI 7337:1974

sotto forma di pirosolfato. Metodo gravi-

metrico.

Analisi chimica dei materiali ferrosi. Deter-

metrico.

UNI 7493:1975 Analisi chimica dei materiali ferrosi.

minazione dell’ alluminio negli acciai. Metodo gravimetrico.

UNI 6647:1970 UNI 6705:1970

Analisi chimica dei minerali di manganese.

Analisi chimica dei minerali di manganese.

Determinazione del fosforo. Metodo gravi-

UNI 7336:1974

Determinazione dell’ ossido di calcio.

metrico.

Analisi chimica dei materiali ferrosi.

Metodo volumetrico al permanganato.

Determinazione dell’ azoto negli acciai.

Metodo spettrofotometrico.

UNI 6644:1970 UNI 6704:1970

Analisi chimica dei minerali di mangane-

Analisi chimica dei minerali di manganese.

se. Determinazione del manganese totale.

UNI 7333:1974

Determinazione dell’ossido di calcio.

Metodo volumetrico al permanganato di

Analisi chimica dei materiali ferrosi.

Metodo gravimetrico all’ ossalato.

potassio.

nelle ghise. Metodo spettrofotometrico al

UNI 6702:1970

UNI 6643:1970

fenilfluorone.

Analisi chimica dei minerali.

Analisi chimica dei minerali di mangane-

Campionamento.

se. Determinazione del manganese totale.

Determinazione dello stagno negli acciai e

UNI 7332:1974

Metodo volumetrico al bismutato.

Analisi chimica dei materiali ferrosi. De-

UNI 6653:1970

terminazione del wolframio negli acciai e

Analisi chimica dei minerali di manganese.

UNI 6642:1970

nelle ghise. Metodo spettrofotometrico al

Determinazione del piombo.

Analisi chimica dei minerali di manganese.

tiocianato di potassio.

Metodo volumetrico.

Determinazione dell’ ossido di alluminio.

Metodo gravimetrico all’ ossichinolina

UNI 7331:1974

UNI 6652:1970

previa separazione dell’ alluminio sotto

Analisi chimica dei materiali ferrosi. Deter-

Analisi chimica dei minerali di manganese.

forma di fosfato.

minazione del cromo negli acciai e nelle

Determinazione del piombo. Metodo gra-

ghise. Metodo fotometrico.

vimetrico.

La Metallurgia Italiana - January 2020

pagina 99

Atti e notizie - AIM news UNI 6639:1970

terminazione del ferro totale nei minerali di

UNI 3658:1955

Analisi chimica dei minerali di manganese.

cromo. Metodo volumetrico per disgrega-

Metodi di analisi chimica dei minerali. De-

Determinazione del nichelio.

zione con perossido di sodio.

terminazione del titanio nei minerali di fer-

Metodo colorimetrico.

ro. Metodi fotometrico e colorimetrico per UNI 3893:1957

disgregazione con perossido di sodio.

UNI 6638:1970

Metodi di analisi chimica dei minerali. De-

Analisi chimica dei minerali di manganese.

terminazione del cromo nei minerali di

UNI 3657:1955 + FA119:1983

Determinazione del nichelio. Metodo gra-

cromo. Metodo volumetrico per disgrega-

Metodi di analisi chimica dei minerali. De-

vimetrico.

zione con perossido di sodio.

terminazione della silice, dell’ ossido di bario, dell’ allumina, della calce e della ma-

UNI 6636:1970

UNI 3892:1957

Analisi chimica dei minerali di manganese.

Metodi di analisi chimica dei minerali. De-

Determinazione del ferro totale.

terminazione dell’ ossido di titanio nei

UNI 3623:1955

Metodo al dicromato di potassio.

minerali di titanio. Metodo gravimetrico al

Metodi di analisi chimica dei minerali. De-

cupferron.

terminazione del fosforo nei minerali di

UNI 6634:1970

gnesia nei minerali di ferro.

ferro. Metodo alcalimetrico, per analisi

Analisi chimica dei minerali di manganese.

UNI 3870:1957

Determinazione della silice (di ossido di si-

Metodi di analisi chimica dei minerali. De-

licio). Metodo gravimetrico per insolubiliz-

terminazione del piombo, del rame e dello

UNI 3619:1955

zazione con acido cloridrico.

zinco nei minerali di ferro. Metodo elettro-

Metodi di analisi chimica dei minerali. De-

litico e gravimetrico per analisi correnti.

terminazione dell’ acqua combinata nei mi-

UNI 6631:1970

correnti.

nerali di ferro.

Analisi chimica dei minerali di manganese.

UNI 3869:1957

Campionamento del minerale caricato su

Metodi di analisi chimica dei minerali. De-

UNI 3618:1955

carri ferroviari.

terminazione del piombo, del rame e dello

Metodi di analisi chimica dei minerali. De-

zinco nei minerali di ferro. Metodo gravi-

terminazione dell’ arsenico nei minerali di

metrico ed elettrolitico.

ferro. Metodo per riduzione con cloruro

UNI 6459:1969 Analisi chimica dei materiali ferrosi. Deter-

stannoso e titolazione con iodio.

minazione del fosforo negli acciai e nelle

UNI 3665:1955

ghise. Metodo alcalimetrico.

Metodi di analisi chimica dei minerali. De-

UNI 3617:1955

terminazione degli alcali totali nei minerali

Metodi di analisi chimica dei minerali. De-

UNI 3896:1957

di ferro. Metodo gravimetrico, per analisi

terminazione dell’ arsenico nei minerali di

Metodi di analisi chimica dei minerali. De-

correnti.

ferro. Metodo gravimetrico.

UNI 3664:1955

UNI 3349:1953

Metodi di analisi chimica dei minerali. De-

Metodi di analisi chimica dei materiali me-

UNI 3895:1957

terminazione dello zolfo nei minerali di fer-

tallici. Determinazione del silicio nel fer-

Metodi di analisi chimica dei minerali. De-

ro. Metodo gravimetrico.

ro-manganese e nel manganese metallo.

terminazione della calce e della magnesia nei minerali di cromo.

terminazione della silice nei minerali di

Metodo gravimetrico.

cromo. Metodo per disgregazione con pe-

UNI 3660:1955

rossido di sodio.

Metodi di analisi chimica dei minerali. De-

UNI 3308:1953

terminazione del comportamento alla cal-

Metodi di analisi chimica dei materiali me-

cinazione all’ aria dei minerali di ferro.

tallici. Determinazione del fosforo nel fer-

UNI 3894:1957 Metodi di analisi chimica dei minerali. De-

La Metallurgia Italiana - gennaio 2020

ro-fosforo. Metodo gravimetrico.

pagina 100

Atti e notizie - AIM news UNI 3306:1953

Metodi di analisi chimica dei materiali me-

rements.

Metodi di analisi chimica dei materiali me-

tallici. Determinazione del cobalto negli

tallici. Determinazione del nichel nel fer-

acciai. Metodo ponderale all’ ossido di zin-

Progetti UNSIDER in inchiesta prEN e

ro-nichel e nel nichel metallo. Metodo

co e all’ alfa-nitroso- beta-naftolo.

ISO/DIS – dicembre 2019 prEN – progetti di norma europei.

gravimetrico.

Norme UNSIDER pubblicate da CEN e UNI 3177:1952

ISO nel mese di novembre 2019

Metodi di analisi chimica dei materiali me-

prEN ISO 15663 Petroleum, petrochemical and natural gas

tallici. Determinazione del vanadio negli

EN 14901-2:2019

industries - Life cycle costing (ISO/DIS

acciai. Metodo per riduzione con solfato

Ductile iron pipes, fittings and accesso-

15663: 2019).

ferroso e titolazione con permanganato

ries - Requirements and test methods for

potassico.

organic coatings of ductile iron fittings and

prEN 14772

accessories - Part 2: Thermoplastic acid

Flanges and their joints - Quality assurance

modified polyolefin coating (TMPO).

inspection and testing of gaskets in accor-

UNI 3106:1950 Metodi di analisi chimica dei materiali me-

dance with the series of standards EN 1514

tallici. Determinazione del nichel negli ac-

EN 14901-1:2014+A1:2019

ciai. Metodo ponderale alla dimetilgliossi-

Ductile iron pipes, fittings and accesso-

ma.

ries - Requirements and test methods for

prEN 13555

organic coatings of ductile iron fittings and

Flanges and their joints - Gasket parame-

UNI 3104:1950

accessories - Part 1: Epoxy coating (heavy

ters and test procedures relevant to the

Metodi di analisi chimica dei materiali me-

duty). .

design rules for gasketed circular flange

tallici. Determinazione del molibdeno ne-

and EN 12560.

connections

gli acciai. Metodo ponderale all’ alfa-ben-

ISO 22055:2019

zoinossima.

Switch and crossing rails.

EN 1514-2:2014/prA1:2019 Flanges and their joints - Gaskets for

UNI 2964:1949

ISO 6892-1:2019

PN-designated flanges - Part 2: Spiral

Metodi di analisi chimica dei materiali me-

Metallic materials - Tensile testing - Part 1:

wound gaskets for use with steel flanges

tallici. Determinazione del tungsteno nel

Method of test at room temperature.

ferro-tungsteno. Metodo ponderale alla cinconina.

prEN 1515-4

Progetti UNSIDER messi allo studio

Flanges and their joints - Bolting - Part 4:

dal CEN (Stage 10.99) – dicembre 2019

Selection of bolting for equipment subject

UNI 2963:1949

to

the

Pressure

Equipment

Directive

Metodi di analisi chimica dei materiali me-

prEN 10250-2 rev

tallici. Determinazione del silicio totale nel

Open die steel forgings for general engine-

ferro-silicio.

ering purposes - Part 2: Non-alloy quality

prEN 993-10

and special steels.

Methods of test for dense shaped refrac-

UNI 2960:1949

2014/68/EU.

tory products - Part 10: Determination of

Metodi di analisi chimica dei materiali me-

prEN 10250-3 rev

permanent change in dimensions on hea-

tallici. Determinazione del titanio negli

Open die steel forgings for general engine-

ting.

acciai e nelle ghise. Metodo al cupferron,

ering purposes - Part 3: Alloy special steels.

ISO/DIS – progetti di norma interna-

fotometrico e colorimetrico (assenza di tungsteno, tantalio e niobio).

prEN 10250-1 rev

zionali

Open die steel forgings for general engiUNI 2959:1949

La Metallurgia Italiana - January 2020

neering purposes - Part 1: General requi-

ISO/DIS 22605

pagina 101

Atti e notizie - AIM news Refractories - Determination of dynamic

ISO/PRF 11971

UNI EN 14901:2014

Young’s modulus(MOE) at elevated tem-

Steel and iron castings - Visual testing of

Tubi, raccordi e accessori in ghisa sferoi-

peratures by impulse excitation of vibra-

surface quality.

dale - Rivestimento epossidico (rinforzato)

tion.

dei raccordi e degli accessori in ghisa sfeISO/FDIS 9647

roidale - Requisiti e metodi di prova

ISO/DIS 19905-3

Steels - Determination of vanadium con-

Petroleum and natural gas industries - Si-

tent - Flame atomic absorption spectro-

Norme UNSIDER pubblicate da CEN e

te-specific assessment of mobile offshore

metric method (FAAS).

ISO nel mese di dicembre 2019

ISO/FDIS 4987

EN ISO 6892-1:2019

ISO/DIS 15663

Steel castings - Liquid penetrant testing.

Metallic materials - Tensile testing - Part 1:

Petroleum, petrochemical and natural gas

ISO/FDIS 4986

Method of test at room temperature (ISO

industries - Life cycle costing.

tic particle testing.

6892-1:2019)

ISO/DIS 15590-2

ISO/FDIS 439

EN 10216-2:2013+A1:2019

Petroleum and natural gas industries - Fac-

Steel and cast irons — Determination of si-

Seamless steel tubes for pressure purpo-

tory bends, fittings and flanges for pipeline

licon content — Gravimetric method.

ses - Technical delivery conditions - Part 2:

units - Part 3: Floating units.

Steel castings - Magne-

transportation systems - Part 2: Fittings.

Non-alloy and alloy steel tubes with speci-

AIM - UNSIDER - Norme pubblicate e ISO/DIS 15590-3

progetti in inchiesta (aggiornamento

Petroleum and natural gas industries - Fac-

31 dicembre 2019).

tory bends, fittings and flanges for pipeline transportation systems - Part 3: Flanges.

fied elevated temperature properties.

ISO 643:2019 Steels - Micrographic determination of the

Norme pubblicate e progetti allo stu-

apparent grain size

dio (elenco).

Progetti UNSIDER messi allo studio

ISO/DIS 4948.2 Classification of steel based on chemical

Norme UNSIDER pubblicate da UNI

composition.

nel mese di dicembre 2019

dal CEN (Stage 10.99) – gennaio 2020 Progetti UNSIDER in inchiesta prEN e

Progetti UNSIDER al voto FprEN e

UNI EN 14901-1:2019

ISO/DIS – gennaio 2020

ISO/FDIS - dicembre 2019

Tubi, raccordi e accessori in ghisa sferoi-

prEN – progetti di norma europei

FprEN - progetti di norma europei

dale - Requisiti e metodi di prova per rivestimenti organici di raccordi ed accessori

prEN ISO 15761

FprEN ISO 10113

in ghisa sferoidale - Parte 1: Rivestimento

Steel gate, globe and check valves for

Metallic materials - Sheet and strip - De-

epossidico (rinforzato)

sizes DN 100 and smaller, for the petro-

termination of plastic strain ratio (ISO/FDIS 10113:2019).

leum and natural gas industries (ISO/DIS EC 1-2019 UNI EN ISO 15630-3:2019

15761:2019).

Acciaio per calcestruzzo armato e calce-

ISO/FDIS - progetti di norma interna-

struzzo armato precompresso - Metodi di

prEN 14772

zionali

prova - Parte 3: Acciaio per calcestruzzo

Flanges and their joints - Quality assurance

armato precompresso

inspection and testing of gaskets in accor-

ISO/FDIS 35102

dance with the series of standards EN 1514

Petroleum and natural gas industries - Ar-

Norme UNSIDER ritirate da UNI nel

ctic operations - Escape, evacuation and

mese di dicembre 2019

rescue from offshore installations.

La Metallurgia Italiana - gennaio 2020

and EN 12560. prEN 13555

pagina 102

Atti e notizie - AIM news Flanges and their joints - Gasket parame-

FprEN ISO 10113

ters and test procedures relevant to the

Metallic materials - Sheet and strip - De-

design rules for gasketed circular flange

termination of plastic strain ratio (ISO/FDIS

connections.

10113:2019)

EN 1514-2:2014/prA1:2019

ISO/FDIS - progetti di norma interna-

Flanges and their joints - Gaskets for

zionali

PN-designated flanges - Part 2: Spiral wound gaskets for use with steel flanges.

ISO 21809-3:2016/FDAmd 1 Petroleum

and

natural

gas

industries

prEN 1515-4

- External coatings for buried or sub-

Flanges and their joints - Bolting - Part 4:

merged pipelines used in pipeline tran-

Selection of bolting for equipment subject

sportation systems - Part 3: Field joint

to

coatings - Amendment 1: Introduction of

the

Pressure

Equipment

Directive

2014/68/EU

mesh-backed coating systems.

ISO/DIS - progetti di norma internazio-

ISO/FDIS 20321

nali

Petroleum, petrochemical and natural gas industries - Safety of machineries - Powe-

ISO/DIS 15663

red elevators.

Petroleum, petrochemical and natural gas industries - Life cycle costing.

ISO/FDIS 20257-1 Installation and equipment for liquefied na-

ISO/DIS 15590-2

tural gas - Design of floating LNG installa-

Petroleum and natural gas industries - Fac-

tions - Part 1: General requirements.

tory bends, fittings and flanges for pipeline transportation systems - Part 2: Fittings.

ISO/FDIS 19959 Steels, nickel alloys and cobalt alloys in-

ISO/DIS 15590-3

vestment castings - Visual testing of surfa-

Petroleum and natural gas industries - Fac-

ce quality.

tory bends, fittings and flanges for pipeline transportation systems - Part 3: Flanges.

ISO/FDIS 4992-1 Steel castings - Ultrasonic testing - Part 1:

ISO/DIS 15177

Steel castings for general purposes.

Steel sheet, hot-rolled twin-roll cast, of commercial quality.

ISO/FDIS 4992-2 Steel castings - Ultrasonic testing - Part 2:

ISO/DIS 3087

Steel castings for highly stressed compo-

Iron ores - Determination of the moisture

nents.

content of a lot

Progetti UNSIDER al voto FprEN e ISO/FDIS - gennaio 2020 FprEN â&#x20AC;&#x201C; progetti di norma europei

La Metallurgia Italiana - January 2020

pagina 103

Atti e notizie - AIM news

Comitati tecnici / Study groups CT AMBIENTE E SICUREZZA (AS) (riunione del 4 dicembre 2019)

Consuntivo di attività svolte • Settore sicurezza: il corso itinerante “Metallurgia Sicura” si è svolto nei giorni 30 ottobre prezzo Cogne Acciai Speciali – Aosta, il giorno 6 novembre 2019 presso AFV Acciaierie Beltrame - Vicenza e il giorno 13 novembre presso Almag – Roncadelle (BS). I partecipanti al corso hanno segnalato una serie di argomenti da approfondire in prossime manifestazioni: si tratta di una quindicina di argomenti che potranno essere trattati nelle giornate di studio del 2020.

Manifestazioni in corso di organizzazione • Settore Ambiente: la GdS “Le emissioni in atmosfera nel settore metallurgico” è confermata per il 5 dicembre a Verona.

Iniziative future • Il premio sicurezza 2020 sarà presentato durante il 38° convegno AIM di Napoli a settembre 2020: è stato richiesto alla segreteria AIM di pubblicare il bando entro fine gennaio.

• In tema di ambiente, i temi per futuri manifestazioni potrebbero essere EPD certificazione prodotto, bilancio di sostenibilità, Lyfe Cycle Assessment, Carbon Footprint ecc.

Stato dell’arte e notizie • Sulla scia di precedenti conversazioni, viene presentato e discusso il caso di un grave infortunio in azienda per fornire spunti di riflessione e di miglioramento.

CT ACCIAIERIA (A) CT FORGIATURA (F)

(riunione congiunta del 11 dicembre 2019) Consuntivo di attività svolte • Si è svolto il corso “Forgiatori” organizzato dal CT Forgiatura nelle date 2-3-9-10 ottobre 1019: l’elaborazione dei questionari di gradimento è stata fatta separatamente per le prime due giornate, dedicate a fabbricazione e trattamenti termici, e singolarmente per le altre due giornate, effettuate con visite presso impianti di produzione. I risultati sono comunque positivi, e il giudizio generale delle differenti valutazioni si mantiene sempre ad alti livelli, con una percentuale di “buono” e “ottimo” che, sommate, arrivano in un caso al 96%. Anche la documentazione distribuita, solitamente un fattore critico in queste manifestazioni, ha soddisfatto i partecipanti (“ottimo” + “buono” oltre 80% e fino al 95%).

Manifestazioni in corso di organizzazione • Per il 24 febbraio 2020 il CT Acciaieria sta organizzando la GdS “Tecnologia e innovazione nei forni ad arco”. L’evento si svolgerà a Dalmine; il programma è stato definito.

Iniziative future Iniziative CT Acciaieria • Il programma del corso “Acciai ad alto carbonio”, organizzato da CT A, è quasi pronto; la manifestazione si svolgerà nel 2020, in primavera o in autunno, probabilmente a Verona.

• Nei mesi di aprile e maggio 2020 si svolgerà il corso itinerante “Macchina Fusoria”. Sono in fase di definizione le aziende che ospiteranno l’evento e i docenti.

Iniziative CT Forgiatori • In collaborazione con il CT Materiali per l’energia si pensa di organizare un corso per i valvolieri, dal titolo “Valvole e leghe speciali per le valvole”, sulla scia della GdS sulle leghe di nichel, che ha avuto ottimo successo. Una bozza di programma sarà pronta a inizio 2020.

Stato dell’arte e notizie • Mapelli, presidente del CT Acciaieria, conferma che nei giorni 17-19 giugno 2020 si terrà a Bari il Convegno Europeo sulla Colata Continua ECCC 2020, organizzato dalla Segreteria AIM. Sono stati raccolti oltre 170 abstract, con buone richieste di partecipazione anche dall’Asia. Diversi sponsor ed espositori hanno manifestato interesse.

• Due nuovi membri di provenienza industriale sono presenti alla riunione e vengono accettati nel CT Acciaieria.

La Metallurgia Italiana - gennaio 2020

pagina 104

Atti e notizie - AIM news

Eventi nazionali ed internazionali / National and international events

2020 March 9-11, Köln, Germany Electrical Engineering of Arc Furnaces April 20-23, Duisburg, Germany Cokemaking April 20-26, Milano, Italy Finiture Green Experience April 26-29, Köln, Germany Refractory Technology May, 4-5, Köln, Germany Hydrogen-based Reduction of Iron Ores May, 6-7, Genova, Italy 27° CONVEGNO NAZIONALE TRATTAMENTI TERMICI May 25-29, Seoul, South Korea MOLTEN - International Molten Slags, Fluxes and Salts conference June 17-19, Bari, Italy ECCC 2020 - 10th European CONTINUOUS CASTING Conference June 23-25, Rome, Italy WMM’20 - 9th International Conference Magnetism and Metallurgy July 1-3, Vicenza, Italy HTDC 2020 - 7th International Conference High Tech Die Casting (ML-P) August 24-26, Shenyang, China TMP - 6th International Conference on ThermoMechanical Processing September 8-12, Cavtat - Dubrovnik, Croatia ICTPMS - 5th International Conference on Thermal Process Modelling and Computer Simulation September 16-18, Napoli, Italy 38° CONVEGNO NAZIONALE AIM October 14-17, Milano, Italy 32.BI-MU October 18-22, BEXCO, Busan, Republic oh Korea WFC - 74th Worls Foundry Congress November, 4-5 Milano, Italy P&E (Polveri&ecoCoating)

2021 22 - 26 March 2021, Bremen, Germany 8th ECIC European Coke and Ironmaking Congress & 9th ICSTI International Conference on Science and Technology of Ironmaking

La Metallurgia Italiana - January 2020

pagina 105

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