La
Metallurgia Italiana
International Journal of the Italian Association for Metallurgy
n. 3 Marzo 2019 Organo ufficiale dell’Associazione Italiana di Metallurgia. Rivista fondata nel 1909
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xiii edizione
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La Metallurgia Italiana
La
Metallurgia Italiana
International Journal of the Italian Association for Metallurgy
n. 3 Marzo 2019 Organo ufficiale dell’Associazione Italiana di Metallurgia. Rivista fondata nel 1909
International Journal of the Italian Association for Metallurgy Organo ufficiale dell’Associazione Italiana di Metallurgia. House organ of AIM Italian Association for Metallurgy. Rivista fondata nel 1909
n. 3 Marzo 2019
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
Anno 111 - ISSN 0026-0843
Leghe leggere / Light metals Nitrogen Adjustment in Molten Steel Using RH Vacuum Degasser S. Tanaka, H. Onoda, S. Kimura, K. Semura 5 Study on heat transfer mechanism of pure calcium cored wire in molten steel by feeding rate L. Jingang, S. Shuomeng, W. Weihua, L. Zhanjun, C. Rensheng, H. Ning 13 Mathematical models, algorithms and software for dynamic simulation of ladle treatment technology O.A. Komolova, K.V. Grigorovich 20 Thermodynamics and kinetics for the evolution of non-metallic inclusions in pipeline steels Y. Zhang, Y. Ren, L. Zhang
25
Corrosion / Corrosione Resistance to localized corrosion of lean duplex stainless steels after brief thermal treatments F. Zanotto, V. Grassi, A. Balbo, F. Zucchi, C. Monticelli
35
Attualità industriale / Industry news
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Gestione editoriale e pubblicità Publisher and marketing office: Siderweb spa Via Don Milani, 5 - 25020 Flero (BS) tel. 030 25 400 06 - fax 030 25 400 41 commerciale@siderweb.com - www.siderweb.com La riproduzione degli articoli e delle illustrazioni è permessa solo citando la fonte e previa autorizzazione della Direzione della rivista. Reproduction in whole or in part of articles and images is permitted only upon receipt of required permission and provided that the source is cited. Reg. Trib. Milano n. 499 del 18/9/1948. Sped. in abb. Post. - D.L.353/2003 (conv. L. 27/02/2004 n. 46) art. 1, comma 1, DCB UD Siderweb spa è iscritta al Roc con il num. 26116
Manifestazioni AIM
44
Technology to Control Inclusions in Stainless Steels edited by: S. Lee, E. Jeong, J.-W. Ki, Kisu Kim, J.Choi, S.-Y. Kim 45 Control of large-sized inclusions and macro segregation in stells for high-speed train wheels and axles edited by: X. Wang, M. Jiang, K. Wang, W. Sun 56 Scenari / Experts' Corner Intervista al Dr. Roberto Moreschi
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Challenges for Secondary Metallurgy to meet future needs edited by: M. Dorndorf - C. Schrade 69 Atti e notizie / Aim news Calendario eventi internazionali
72
Norme pubblicate e progetti in inchiesta
73
l’editoriale La Metallurgia Italiana
Dong Joon Min PhD Professor, POSCO Chair Materials Science and Engineering, Yonsei University, Seoul Korea
4
As consumer expectations on steel quality become increasingly stringent, there has been significant focus on research and technology development related to secondary refining and controlling non-metallic inclusions in steels. In addition, new steel product developments for unconventional energy exploration and storage, ultra-light steel applications for transportation, and electrical steels to minimize core loss in power conversion have been some of the factors testing the limits of secondary refining technologies. Many of the global companies that compete with one another have produced similar products within this spectrum and the perspective companies have worked vigorously to distinguish their products through quality, while maintaining cost-competitiveness. Thus, with over 300 million tons of steel over-capacity estimated and the overall structural slowdown with the oversupply that has become the “New Normal�, steel producers must distinguish their products to increase economic returns and maintain market dominance of their products. In this respect, a focus on secondary refining and non-metallic inclusions is an essential area for steelmakers to concentrate their research and technical efforts. Within this special issue, several specific topics to enhance steel quality in secondary steelmaking that encompasses process optimization and theoretical elucidation of inclusion formation and control have been introduced. For process optimization, dynamic ladle process control using software optimization, precise melt temperature control using finite element modeling, optimization of nitrogen at the RH degasser, and stirring practices at the AOD (argon oxygen decarburization) have been examined. In the work of Komolova and Girgorovich, physical and chemical models based on mass and energy conservation with non-equilibrium thermodynamics was utilized for dynamic simulation and optimization of steel ladle treatment, which was found to correlate well with actual plant data. Liu et al. described the effect of calciumsilicon cored wire feeding on the melt temperature and optimizing the feeding practice, which could provide process guidance for calcium treatment in steels. According to Tanaka et al., dissolved nitrogen in case-hardening steels could be further adjusted in the RH-degasser by regulating the gas flow rate and vacuum during degassing, which resulted in the development of a nitrogen control model to accurately predict and control nitrogen. Beyond carbon steels, Lee et al. elaborated on inclusions in stainless steels formed during AOD processing. Their work indicated accurate control of the argon stirring energy could significantly increase the cleanliness and subsequent melt temperature and slag basicity control could also enhance quality. In summary, these time-appropriate contents have a common goal to improve steel cleanliness and considering the limitless boundaries of consumer expectations, quality improvements in secondary steelmaking will continue to be an active area of research and development for steel producers and researchers alike.
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno
Nitrogen adjustment in molten steel using rh vacuum degasser S. Tanaka, H. Onoda, S. Kimura, K. Semura
Technique for adjustment of nitrogen concentration [N] of molten steel in a RH degasser was improved to meet [N] specifications demanded for several kinds of case-hardening steel. To develop this technique, actual plant trials were carried out to investigate patterns of change in [N] under various RH treatment conditions. This investigation focused particularly on Sulfur content because it is a surface active element. The apparent equilibrium of [N] in molten steel increased with an increase in circulating of N2 gas flow rate, a decrease in degree of vacuum in the RH vessel and an increase in sulfur concentration [S]. The rates of [N] desorption and absorption both decreased with an increase in [S]. To determine appropriate RH treatment conditions according to the various [N] specifications, a prediction model was adopted. The calculation model includes various parameters, such as the interface area of circulating N2 or Ar gas bubbles, the surface area of molten steel in the RH vessel when using circulating N2 or Ar gas, and the mass transfer coefficient of [N] in molten steel. These parameter values were estimated by fitting calculated changes in [N] to actual measured data under a vacuum of 30 Torr and 0.5 Torr in the RH vessel. This model was used as guidance for the operators of RH to control [N] onsite. As a result, [N] adjustment could be done with high accuracy and case-hardening steel having the various [N] specifications was successfully manufactured using RH with only N2 gas flow.
KEYWORDS: SECONDARY REFINING – LADLE METALLURGY– NITROGEN – GAS METAL REACTION – PREDICTION MODEL – VACUUM DEGASSER
INTRODUCTION Iron and steelmaking processes in Kobe Works were shut in November 2017 to be consolidated into Kakogawa Works for strengthening of cost competitiveness of special steel wire rods and bars. Therefore, steel blooms and billets manufactured in Kobe Works have to be done in Kakogawa Works. One of these steels is high nitrogen steel, for example case-hardening steel. This is for making gears used for sliding parts, for example transmission gears, differential gears and so on. [N] is added because these parts need to be tough and long-lasting. Molten steel for case-hardening steel was refined by ASEA-SKF in Kobe Works and needed vacuum treatment to decrease [H] to prevent hair cracks. Table 1 shows a comparison of ASEASKF and RH. [N] desorption in the vacuum degasser (VD) process where ASEA-SKF is smaller than RH because the whole surface of molten steel is covered with molten slag and the surface area of the molten steel under vacuum is small. Nitride alloys were used to adjust [N] in molten steel in ASEA-SKF. On the contrary, in the RH vessel the surface area of molten steel exposed to vacuum atmosphere is large because the surface of the molten steel is not covered with molten slag and is bursted by circulating gas. Figure 1 shows the comparison of [N] desorption using ASEA-SKF and RH. [N] is decreased rapidly using RH even when [N] is added in the previous process, for example converter and Ladle Furnace (LF). Therefore, in RH Ar gas was used instead of N2 gas as the circulating gas to adjust [N]. La Metallurgia Italiana - n. 3 2019
In RH there were some research and prediction models of [N] behavior(1-6). However, there were few reports that [N] was controlled for various [N] specifications in RH. In present study, patterns of change in [N] were investigated under various conditions to develop a prediction model of [N] in RH to provide guidance for the operators of RH to control [N] onsite. Among them, high nitrogen steel, for example, case-hardening steel is one of the steel grades that Kakogawa Works had not manufactured.
Shota Tanaka, Hiroyuki Onoda, Sei Kimura, Koichiro Semura
Steelmaking Development Section, Steelmaking Development Department, R&D Laboratory, Iron and Steel Business, Kobe Steel LTD, Japan
5
Secondary refining VD(ASEA-SKF)
RH
Way of [N] addition
Nitride alloys
Nitride alloys with circulating Ar gas
[N] desorption in vacuum degasser process
10 ppm
~ 150 ppm
Apparatus
Fig. 1 – Comparison of [N] desorption using ASEA-SKF and RH EXPERIMENTAL We carried out some actual plant trials to investigate patterns of change in [N] under various RH treatment conditions. Table 2 shows the experimental conditions of RH. Initial [N], circulating N2 gas flow rate, degree of vacuum in RH vessel and [S] were
changed. We focused on [S] content because it is a surface active element(1, 2). Some samples were taken from molten steel in a ladle during RH treatment and [N] was analyzed by insert gas fusion – using the thermal conductivity method.
Tab. 2 – Comparison of [N] desorption using ASEA-SKF and RH
6
Heat size
250 ton
Initial [N] concentration
65 ~ 180 ppm
Circulating N2 gas flow rate
1,500 ~ 5,000 NL/min
Degree of vacuum in RH vessel
0.5 and 30 Torr
[S] concentration
0.003 ~ 0.055 %
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno RESULTS Effect of circulating N2 gas flow rate Figure 2 shows the behavior of [N] when the circulating N2 gas flow rate was varied from 1,500 to 5,000 NL/min under a vacuum of 30 Torr and [S] was 0.008 to 0.010 %. Apparent
equilibrium [N] were about 105 ppm, about 135 ppm and over 150 ppm when circulating N2 gas flow rate was 1,500 NL/min, 3,000 NL/min and 5,000 NL/min respectively. It increased with an increase in circulating N2 gas flow rate because the rate of [N] absorption from N2 gas bubbles increased.
Fig. 2 – Influence of circulating N2 gas flow rate on [N] behavior Effect of vacuum degree Figure 3 shows the influence of [N] when degree of vacuum was 0.5 and 30 Torr where circulating N2 gas flow rate was 3,000 NL/min and [S] was 0.003 - 0007 %. Apparent equilibrium [N] was about 60 ppm when degree of vacuum was 0.5
Torr and about 120 ppm under a vacuum of 30 Torr. It increased with as the degree of vacuum decreased because the rate of [N] desorption from surface of molten steel in the RH vessel decreased.
Fig. 3 – Influence of degree of vacuum on [N] behavior Effect of sulfur concentration Figure 4 shows the influence of [N] when [S] was varied from 0.003 to 0.055 % where circulating N2 gas flow rate was 3,000 NL/min and the degree of vacuum was 0.5 Torr. Apparent equilibrium [N] were about 80 ppm, 110 ppm and 120ppm when [S] was 0.003%, 0.015 % and 0.055% respectively. It increased with an increase in [S]. And the rate of change in [N] La Metallurgia Italiana - n. 3 2019
was decreased with an increase in [S]. These results suggest that the rates of [N] desorption and absorption were decreased because sulfur acted as a surface active element. Figure 5 shows the reaction sites of [N] in RH. The reason why apparent equilibrium [N] was changed by [S] is that balances between the rates of [N] desorption from the surface of molten steel and absorption from N2 gas bubbles is changed. We assumed that 7
Secondary refining the surface area of molten steel in the RH vessel is larger than interface area of N2 gas bubbles and molten steel. Therefore,
the decrease in the rate of [N] desorption is larger than the decrease in the rate of [N] absorption.
Fig. 4 – Influence of [S] on [N] behavior
Fig. 5 – Reaction sites of [N] in RH MODEL OF [N] BEHAVIOUR Details of model [N] needs to be adjusted within a limited time in RH to be in time to start casting. This prediction model was adopted to determin appropriate RH treatment conditions, for example the kind of circulating gas, circulating gas flow rate and degree of vacuum, according to the various [N] specifications, the con-
centration of other components and RH treatment time. The model was based on Nabeshima’s model(3) and can explained as follows. The rates of [N] absorption in the RH vessel and ladle are calculated using equations [1] and [2] when perfect mixing in molten steel is assumed, and the reaction sites of [N] are the 4 sites shown in Fig. 5.
[1]
8
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno [2]
where [N]L : [N] of molten steel in ladle (%) [N]V : [N] of molten steel in the RH vessel (%) VL : volume of molten steel in ladle (m3) VV : volume of molten steel in the RH vessel (m3) t : reaction time (min) RS : the rate of [N] absorption at surface of molten steel in the RH vessel (%/min) RAr : the rate of [N] absorption at interface of circulation Ar gas (%/min) RN2 : the rate of [N] absorption of interface at circulation N2 gas (%/min) Rleak : the rate of [N] absorption of air leak from snorkel (%/min) Q (m3/min) is the circulatiing quantity of molten steel and is calculated using equation [3] proposed by Kuwabara et al.(4).) [3]
where Qg : circulation gas flow rate (Nm3/min) D : diameter of snorkel (m) P0 : pressure at height of nozzle of circulation gas (atm) PV : pressure in the RH vessel (atm) Ď Fe : density of molten steel (ton/m3) RS and RAr are calculated using equation [4] when rate-determining steps are both transferring [N] to the surface of the molten steel and the chemical reaction of N on the surface of the molten steel. [4]
where n : S or Ar of subscript An : reaction interfacial area of surface of molten steel or interface between Ar gas bubbles and molten steel (m2) [N]i,n : initial [N] of surface of molten steel or interface between Ar gas bubbles and molten steel (%) [N]e,n : equilibrium [N] of surface of molten steel or interface between Ar gas bubbles and molten steel (%) km : mass transfer coefficient of [N] of molten steel (m/min) kr : chemical reaction rate constant of [N] (m/(min/%)) RS and RAr are calculated using equation [5] when [N]i,n from equation [4] was deleted. [5]
La Metallurgia Italiana - n. 3 2019
9
Secondary refining RN2 is calculated using equation [6] taking into consideration that N2 partial pressure in gas bubbles is different between the height of the snorkel nozzle and the surface of molten steel. [6]
where h : height between nozzle of snorkel and surface of molten steel (m) AN2 : reaction interfacial area of interface between N2 gas bubbles and molten steel (m) [N]e,N2 : equilibrium [N] of interface between N2 gas bubbles and molten steel (%) [N]e,n and [N]e,N2 are calculated using equations [7] and [8]. [7]
[8]
where PN2,n : N2 partial pressure in the RH vessel or Ar gas bubbles (atm) Pb,N2 : N2 partial pressure in N2 gas bubbles (atm) fN : activity coefficient of [N] R : gas constant (J/(mol•K) Pb,N2(x) is N2 partial pressure in N2 bubbles and is calculated using equation [9]. [9]
where g: acceleration of gravity (m/s2) kr is expressed as equation [10] estimated by Harashima et al.(1).
[10]
where aO : activity of [O] aS : activity of [S] 10
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno AAr and AN2 are considered to be proportional to geometrically 2/3 square of quantity of circulation gas in molten steel and they are calculated using equation [11].
[11]
where m : Ar or N2 of subscript αm : coefficient of Am tb : bubbles residence time (min) Rleak is calculated using equation [12].
[12]
where Qleak : 0.17 Nm3/min(4) MN2 : molecular weight of N2 Parameters fitting Unknown parameter values in the model are km, AS, αAr and αN2. They were estimated by fitting calculated changes in [N] to actual measured data under a vacuum of 30 Torr and 0.5 Torr in the RH vessel. Table 3 shows estimated parameter values. The value of km is assumed to be 10,000 m/min assuming that nitrogen transfers so rapid in molten steel and it doesn’t become mass transfer
rate-determining step. The apparent reaction surface area of molten steel using circulating N2 gas is about 4 times larger than the sectional area of the lower vessel under a vacuum of 30 Torr and about 10 times larger under 0.5 Torr. We assume that the surface area of molten steel in the RH vessel is large because [N] desorption is increased under a high degree of vacuum.
Tab. 3 – Estimated unknown parameter values Degree of vacuum 0.5 Torr Km
30 Torr 10,000 m/min
AS using Ar gas
2
43 m
17 m2
AS using N2 gas
74 m2
16 m2
αAr
38.5
36.5
33.7
32.0
αN2
RESULTS OF NITOROGEN ADJUSTMENT IN RH This model was used as guidance for the operators of RH to control [N] onsite. Both static control which determined RH treatment conditions before starting the treatment and dynamic control in which changing RH treatment conditions on the way to treatment end according to actual analysis of [N] composition were used to adjust [N]. Figure 6 shows the results of
La Metallurgia Italiana - n. 3 2019
[N] control and Figure 7 shows the accuracy of [N] adjustment. [N] was controlled with a high degree of accuracy and an accuracy of 6.3 ppm was achieved for [N] adjustment. Case-hardening steel having various [N] specifications was successfully manufactured using RH with only N2 gas flow. In addition, the consumption of nitride alloys was greatly reduced compared to ASEA-SKF.
11
Secondary refining
a) – [N] absorption control under dynamic control
b) – [N] desorption control under dynamic control
Fig. 6 – Results of [N] adjustment
Fig. 7 – Accuracy of [N] adjustment CONCLUSIONS • [N] behaviors were investigated under various RH treatment conditions. Apparent equilibrium [N] in molten steel increased with an increase in circulating N2 gas flow rate, a decrease of degree of vacuum in the RH vessel and an increase of [S]. Rates of [N] desorption and absorption both decreased with an increase in [S]. • To determine appropriate RH treatment conditions according to the various [N] specifications, prediction model of [N] in RH
was adopted and unknown parameter values in the model were estimated by fitting calculated changes in [N] to actual measured data. • This model is used for guidance to control [N] onsite. As a result, [N] adjustment could be achieved with high accuracy and case-hardening steel having the various [N] specifications was successfully manufactured in Kakogawa Works using RH with only N2 gas flow.
REFERENCES [1]
K. Harashima, S. Mizoguchi, H. Kajioka and K. Sakakura: Tetsu-to-Hagané, 73(1987), 1559.
[2]
S. Mukawa, Y. Mizukami and Y. Ueshima: Tetsu-to-Hagané, 84(1998), 411.
[3]
S. Nabeshima, H. Ogawa and Y. Miki: Tetsu-to-Hagané, 101(2015), 627.
[4]
Y. Kato, T. Kirihara, K. Yamaguchi, T. Fujii and S. Ohmiya: Tetsu-to-Hagané, 83(1997), 18.
[5]
M. Takahashi, Y. Han, M. Sano, K. Mori and M. Hirasawa: Tetsu-to-Hagané, 74(1988), 69.
[6]
K. Kadoguchi, M. Sano and K. Mori: Tetsu-to-Hagané, 71(1985), 70.
[7]
T. Kuwabara, K. Umezawa, K. Mori and H. Watanabe: Trans. Iron Steel Inst. Jpn., 28(1988), 305.
12
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno
Study on heat transfer mechanism of pure calcium cored wire in molten steel by feeding rate L. Jingang, S. Shuomeng, W. Weihua. L. Zhanjun, C. Rensheng, H. Ning
In order to descript heat transfer mechanism of pure calcium cored wire in molten steel accurately and comprehensively, the temperature changes of the pure calcium cored wire fed into the molten steel at different feeding rate were studied using heat transfer finite element model of pure calcium cored wire in molten steel. Experimental results show when the molten steel temperature is 1570˚C, and the feeding rate is 1.50m/s-5.00m/s, the time required for the pure calcium core to reach the melting point is 1.19s-1.43s, and the time to reach the boiling point is 2.75s-3.36s; When the temperature of the calcium core reaches the boiling point of 1484˚C, the steel sheet of the pure calcium cored is 1496 ~1498˚C, and the steel sheet is in the state of non-melting. Through numerical simulation and industrial test, the best feeding rate is 1.75m/s, and the yield of calcium is 26.1%, which provides technical guidance for the research and development of calcium treatment process.
KEYWORDS: PURE CALCIUM CORED WIRE – FEEDING RATE – MELTING POINT – BOILING POINT – CALCIUM TREATMENT PROCESS
INTRODUCTION In the process of molten steel refining, Calcium silicon cored wire are adopted in most of the plants. The core of the calcium silicon wire is filled with powdered calcium silicate powder, and the exterior is wrapped with steel sheet. Silicon calcium powder has the drawback of uneven and unstable. Moreover, the granular material has large specific surface area and large gas gap. The calcium in it is easily oxidized during storage and transportation, so that calcium is wasted [1-3] in vain, resulting in low calcium yield. In addition, the calcium silicon cored wire also adds more silicon in the molten steel, which makes it difficult to control the silicon content produced in the low silicon content steel grade. Besides, during the use process, because of the low strength of calcium silicon cored wire, it is easy to slip and run away when feeding wire, which brings adverse effects to the precise control of calcium content and stable continuous production. Pure calcium cored wire with external steel sheet and internal high density solid cored pure calcium rod high-pressure pulled out, being not slippery, without deviation while wire feeding, can feed the calcium into deep position of molten steel, without adding silicon, or other impurities, so as the raising of varieties of steel smelting requirements, in order to improve the effect of calcium treatment, and increase the purity of molten steel[4,5], the study of pure calcium cored wire replacement calcium silicon cored wire is started. Calcium is a flammable and readily oxidizable metal, which can not be
La Metallurgia Italiana - n. 3 2019
observed at normal temperature. After the feeding of molten steel, the melting and gasification process of pure calcium cored wire cannot be observed. The change of pure calcium cored wire temperature and the melting and gasification behavior after entering molten steel cannot be directly measured. Therefore, the finite element method is used to simulate the melting and gasification mechanism of pure calcium cored wire feeding into molten steel at different feed speed, which provides a theoretical basis for the specification model of pure calcium cored wire and the wire feeding process optimization.
Liu Jingang, Wang Weihua, Li Zhanjun, Chu Rensheng, Hao Ning
ShouGang Research Institute of Technology, Beijing, 100043, China; Beijing key Laboratory of Green Recyclable Process for Iron & steel Production Technology, Beijing, 100043, China; Beijing Engineering Research Center of Energy Steel, Beijing, 100043, China
Sun Shuomeng
Ningbo Shougang auto parts co., Ltd, Ningbo, 315336, China
13
Secondary refining Numerical model establishment Selection of model parameters The diameter of the pure calcium cored wire is 8.82mm, the diameter of the pure calcium core is 7.7mm, and the thickness of the steel sheet is 0.56mm.
Calcium content of the pure calcium cored wire is more than 95%, core weight: sheet weight =3:7 The physical parameters of the metal Ca are shown in Table 1, and the physical properties of steel are shown in Fig. 1.
Tab. 1 – Physical properties of metal calcium melting point (˚C )
boiling point (˚C )
atomic quantity
density (kg/m3)
specific heat (J/Kg.K)
thermal conductibity ( W/m.K )
calcium core weight ( g/m )
839
1484
40.08
1550
632
201
67
(a)
(b)
(c)
(a) Density (b)Thermal conductivity (c)Equivalent specific heat Fig. 1 – Physical properties of steel (1) The initial condition and boundary condition of the model Set 1570˚C in refining end of low carbon steel as the initial temperature of the molten steel, room temperature of 20˚C as pure calcium cored wire initial temperature, the heat transfer between molten steel and pure calcium cored wire as the third
boundary condition, pure calcium cored wire enter into the molten steel at a certain speed that is strong for the convective heat transfer boundary condition. The convection heat transfer coefficient between the molten steel and the pure calcium cored wire is calculated by Eq.1.
[1]
where A−Comprehensive heat transfer coefficient, 3.6~6.8×10-3,take 4.17×10-3; D−Pure calcium cored wire diameter, 0.882×10-2m; V−the speed of pure calcium cored wire entering into molten steel, 1.50~5.00m·s-1; P−Steel density, 7×103Kg·m-3; μ−Dynamic viscosity of molten steel, 5.5×10-3Kg·s·m-2; C−Specific heat of molten steel, 842 J·Kg-1·℃-1. λ−Thermal conductivity of molten steel,20~37 W·m-1·℃-1. The convective heat transfer coefficient of its corresponding molten steel and pure calcium cored wire is shown in Table 2.
14
La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno Tab. 2 – Convective heat transfer coefficient of molten steel and pure calcium cored wire at different feeding rate. Feeding rate, m/s
1.50
1.75
2.00
3.00
4.00
5.00
Convective heat transfer coefficienth, W/m2.˚C (×103)
12.67
14.33
15.96
22.06
27.77
33.20
In order to simplify the calculation, the steel shell condensed on the outer surface of pure calcium cored wire when the pure calcium cored wire just enters the molten steel is not considered. While the calcium core is simulated, when the calcium core reaches or exceeds the boiling point, it is still not treated with gasification for convenience of calculation. (2)The establishment of the model Since pure calcium cored wire length is far greater than the
cross-sectional size, heat transfer is basically in the two-dimensional space, the heat transfer along the length direction of the transfer wire is negligible. So the 2D finite element method is adopted in this model pure calcium cored wire model is set by two dimensional modeling and divided into a finite element mesh, model element number is 2060, the number of nodes is 2109, the finite element model of heat transfer wire as shown in Fig. 2.
Fig. 2 – The finite element model of heat transfer of pure calcium cored wire
Fig. 3 – Temperature distribution of cross section of pure calcium cored wire at different feeding rate
Based on Fig. 2, combined the diameter of the pure calcium core is 7.7mm, so the calcium core is within 3.85mm of the center, and from 3.85mm to 4.41mm part is the steel sheet.
thermal conductivity of steel sheet. The steel sheet becomes a restrictive part of heat transfer, so the heating rate is different, and the heating rate of calcium core with relatively fast heat conduction speed is not very different. It is worth noting that there is an obvious temperature gradient due to the contact heat transfer between two substances at the junction surface of 3.85mm calcium core and steel sheet, and the temperature gradient increases with the increase of wire feeding rate. Because the outer layer is steel sheet and its melting point is high, when the pure calcium core reaches the melting temperature of 839˚C, the steel sheet is far below its melting temperature, that is, it is still solid. In order to obtain the time required for the melting point of the calcium core at different feeding rate, the temperature changes of the center and edge of the pure calcium cored wire with time are analyzed, as shown in Fig. 4.
INFLUENCE OF MELTING OF CALCIUM CORE AT DIFFERENT FEEDING RATE First, the temperature of the cross section of the pure calcium cored wire is calculated when the pure calcium cored wire is inserted into the molten steel at different feeding rate to reach the melting temperature. The results are shown in Fig. 3. From Fig. 3, it can be seen that the temperature difference inside the calcium core (0-3.85mm,the same below) is very small at different feeding rate, and the temperature of the steel sheet (when the distance is greater than 3.85mm, the same below) increases with the increase of feeding rate. This is because the thermal conductivity of calcium core is much larger than the
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(a) Pure calcium cored wire center (b) Pure calcium cored wire edge Fig. 4 – The temperature change of pure calcium cored wire with time at different feeding rate From Fig. 4, it is known that the temperature at both the center and the edge of pure calcium cored wire increases with time, and the heating rate decreases with the increase of time. The higher the feeding rate is, the higher the temperature of the center and the edge of the pure calcium cored wire is. From Fig. 4 (a), it is known that when the feed speed is 1.5m/s, 1.75m/s, 2m/s, 3m/s, 4m/s and 5m/s, the time of pure calcium core temperature rises to the melting point, that is, the melting time of pure calcium cored wire core is 1.43s, 1.40s, 1.38s, 1.29s, 1.38s, 1.29s, 1.23s and 1.19s.
INFLUENCE ON CALCIUM CORE GASIFICATION AT DIFFERENT FEEDING RATE When the calcium core reaches gasification temperature, The existence state of steel sheet in the outer layer of pure calcium cored wire has a great influence on the pure calcium cored wire, In the case of liquid, the calcium core can break through the liquid steel sheet and react with the molten steel directly. If the steel sheet is still solid, the core should be ejected from the end to react with the molten steel. Therefore, in order to get the existence state of the steel sheet when the calcium core reaches the gasification temperature, the temperature change of the cross section of pure calcium cored wire at different feeding rate when it reaches the gasification temperature is calculated, and the results are shown in Fig. 5 to 9.
Fig. 5 – Temperature distribution of cross section of pure calcium cored wire at different feeding rate
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Metallurgia fuori forno Fig. 5 shows the temperature distribution of cross section of pure calcium cored wire at different feeding rate. It is known that the temperature change is the same as that of the Fig3, but the temperature difference between the calcium core and the steel sheet is smaller when the calcium core reaches the gasification temperature. When the feeding rate is 1.5m/s,
1.75m/s, 2m/s, 3m/s, 4m/s and 5m/s, the temperature difference between the core edge and the center is 11.7˚C, 12.1˚C, 12.3˚C, 13˚C, 13.7˚C, 14.3˚C, respectively. The temperature change of the center and edge of the pure calcium cored wire at different feeding rate is shown in Fig. 6
(a) Pure calcium cored wire center (b) Pure calcium cored wire edge Fig. 6 – The change of pure calcium cored wire temperature with time at different feeding rate From Fig. 6, we know that the temperature at both the edge and center of the pure calcium cored wire increase with time, while the heating rate decreases with time. The higher the feeding rate is, the higher the temperature of the center and the edge of the pure calcium cored wire is. The center temperature of the pure calcium cored wire rises to the boiling point at 1484˚C, that is, the gasification temperature, is taken as the standard. It can be get that when the wire feeding rate are 1.5m/s, 1.75m/s, 2m/s, 3m/s, 4m/s, 5m/s, The time of the pure calcium cored wire to reach the gasification temperature correspondingly are: 3.36s, 3.28s, 3.20s, 2.99s, 2.85s, 2.75s. The steel of the outer layer of pure calcium cored wire (the steel sheet) is SPCC. According to its chemical composition, the solidus temperature is 1474˚C, and the liquidus temperature is 1524˚C. In the mushy zone, the dividing wire of the solid frac-
(a) 2.53s
(b) 2.53s
tion fs is 0.7, and the corresponding temperature is 1509˚C. When the solid fraction fs is less than 0.7, it is considered that the mixed area of solid liquid phase is pure liquid zone, that is, the melting temperature of outer skin is 1509˚C. According to Fig. 6, when the calcium core temperature rises to 1484˚C, the temperature of steel sheet is 1496˚C~1498˚C, so it doesn't melt. INDUSTRIAL TEST VERIFICATION OF PURE CALCIUM CORED WIRE MELTING MECHANISM In order to verify the calculation results of the melting mechanism of pure calcium cored wire, an industrial trial was carried out. The appearance of pure calcium cored wire pulled out from molten steel in different time is shown in Fig. 7. The appearance of the pure calcium cored wire is shown in Fig. 8.
(c) 2.93s
(d) 3.20s
Fig. 7 – The appearance of pure calcium cored wire pulled out from molten steel in different time La Metallurgia Italiana - n. 3 2019
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Fig. 8 – Pure calcium cored wire melting appearance From Fig. 7, based on the analysis of residual pure calcium cored wire end, when the pure calcium cored wire enters the molten steel, firstly, the pure calcium cored wire end is quickly heated to the liquid. At this time, the temperature of the steel sheet did not reach the melting point. Therefore, the calcium core was in a liquid state in the tube formed by the steel sheet, and then dropped off after leaving the molten steel and had strong chemical reaction with oxygen in the air, as shown in Fig. 7 (a) and (b). If the time of the pure calcium cored wire in the molten steel is prolonged, the calcium core end is continuously heated in the molten steel and turns into a gaseous state. The gaseous calcium will be ejected from the pure calcium cored wire end to form calcium vapor and directly react with oxygen in the air, as shown in Fig. 7 (c). The calcium vapor floats in the molten steel and partially melted in the molten steel during the floatation. The non melting part is lost in the air, and the gaseous calcium will accelerate the melting of the steel sheet during the ejection process. When the critical time is exceeded, the calcium core will be vaporized and the steel sheet will be melted completely, that is, the pure calcium cored wire is melted into the molten steel, as shown in Fig. 7 (d). In Fig. 8, the main component of the white substance on the end of the pure calcium cored wire is calcium compounds, the pure calcium cored wire is wrapped with slag in the other parts except the end. After removing the slag from the pure calcium cored wire, the surface is covered with a layer of steel shell. It is known that when the pure calcium cored wire is inserted into the molten steel, the molten steel contacted with the pure calcium cored wire is cooled and adhered to the pure calcium
cored wire. This act prevents the pure calcium cored wire except the end from being molten or even vaporized firstly, In other words, the melting and gasification of the pure calcium cored wire start from the end and gradually move back. This helps to insert the pure calcium cored wire into the molten steel deeper and increases the yield of calcium. FEEDING RATE CONTROL The feeding rate of the pure calcium cored wire should have an appropriate value. If the feeding rate is too high, the calcium in the ladle will be too concentrated, After the pure calcium cored wire is completely melted, a large amount of calcium vapor is produced locally, forming a large number of "bubbles". In the process of floatation, a large number of "bubbles" can not be absorbed by molten steel and a large amount of molten steel escapes, resulting in seething and splashing of molten steel, and reducing the yield of calcium. If the feeding rate is too slow and the time of calcium gasification is prolonged, the pure calcium cored wire will be centralized and vaporized at the bottom, which will also cause the above situation and reduce the yield of calcium. The depth of the molten steel is calculated by 3m, and the pure calcium cored wire is not vertically downward after entering the molten steel, but it will spiral or bend down. So the path of the pure calcium cored wire reached at the bottom is estimated as 6m, which is two times of the depth of the steel. It is known from table 3 that the best feeding rate is 1.75-2.0m/s, which makes sure the pure calcium cored wire reach the bottom without easily heaping up.
Tab. 3 – Gasification time and the time used by the pure calcium cored wire reaching the bottom of the molten steel at different feeding rate THE TIME USED BY THE PURE CALCIUM CORED FEEDIN GRATE, GASIFICATION WIRE REACHING THE BOTTOM OF THE MOLTEN m/s TIME, s STEEL ,s
18
1.50
3.36
4.00
1.75
3.28
3.43
2.00
3.20
3.00
3.00
2.99
2.00
4.00
2.85
1.50
5.00
2.75
1.20 La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno In order to verify this result, a comparative experiment was carried out on the yield of calcium element of pure calcium cored wire at different feeding rate. The results are shown in Fig. 9.
Fig. 9 – The relationship between the different feeding rate of pure calcium cored wire and the yield of calcium It is known from Fig. 9, when the feeding rate is 1.50m/s, 1.75m/s, 2.00m/s, 3.00m/s, 4.00m/s, the average yield of calcium is 24.7%, 26.1%, 22.1%, 18.4%, and 15.8%. So, the industrial test of different feeding rate shows that the optimum feeding rate is 1.75m/s, which is in accordance with the numerical simulation results. CONCLUSION Using heat transfer 2D finite element model of pure calcium cored wire in molten steel, the temperature changes of pure calcium cored wire in molten steel at different wire feeding rate are calculated, The rules of melting and gasification of calcium core is analyzed, and the optimum feeding rate of pure calcium cored wire is determined combined with industrial test, the
conclusions as follows: 1)The junction point of the calcium core and the outer steel sheet is the turning point of the heat transfer, and it is also the limiting step of the heat transfer. 2)When the molten steel temperature is 1570˚C, and the feeding rate is 1.50m/s-5.00m/s, the time required for the pure calcium cored wire to reach the melting point is 1.19s-1.43s, and the time to reach the boiling point is 2.75s-3.36s. 3)When the temperature of the calcium core reaches the boiling point 1484˚C, and the temperature of the steel sheet of the pure calcium cored wire is 1496˚C ~1498˚C, the steel sheet is in the state of non-melting. 4)The optimum feeding rate is 1.75m/s, and the yield of calcium is 26.1%.
REFERENCES [1]
SUN Yan-hui, FANG Zhong-qiang. Formation of intermediate products during calcium treatment and modification routes of alumina inclusions by intermediate products[J]. Journal of University of Science and Technology Beijing ,2014, 36(12): 16151625.
[2]
Pretorius E B, Oltmann H G, Cash T. The Effective Modification of Spinel Inclusions by Ca Treatment in LCAK Steel [J]. Iron & Steel Technology, 2010, (8): 31-44.
[3]
ZHANG Cai-jun, CAI Kai-ke, YUAN Wei-xia. Study on Sulfide Inclusions and Effect of Calcium Treatment for Pipeline Steel[J].Iron and Steel,2006, 41(8): 31-33.
[4]
Zhang Xiaobing. Thermodynamic modeling for controls of deoxidation and oxide inclusions in molten steel [J]. Acta Metallurgica Sinica. 2004, 40(5): 509-514.
[5]
LUO Lei, SUN Yanhui, CHEN Yong, WU Guorong. Effect of calcium treatment on the non-metallic inclusions of pipeline steel [J]. Iron and Steel.2013, 48( 1): 42-45.
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Secondary refining
Mathematical models, algorithms and software for dynamic simulation of leadle treatment technology O.A. Komolova, K.V. Grigorovich
The original software for dynamic simulation of ladle treatment of steel was developed. Physical and chemical models based on the conservation law of mass and energy, as well as the principles of non equilibrium thermodynamics were used. All the stages of the process (zones) were taken into account in this software. This software takes into consideration input data such as: ladle equipment facilities, initial temperature, slag and metal mass compositions, input time and mass of additives, blowing, electrical and time regimes, thermodynamic database, thermal, physical and chemical databases for additives and inert gas, production database (for statistics). This software allows us to calculate the main characteristics of the ladle treatment such as temperatures and chemical composition of slag and the steel melts. For validation of the software, the results of ladle treatment of real heats of steel and sampling control results were used. This software can be used in online calculations and control of process parameters during ladle treatment, simulation and optimization of ladle treatment technology, teaching and training of steelmaking staff.
KEYWORDS: MATHEMATICAL MODELS – LADLE TREATMENT – NON EQUILIBRIUM THERMODYNAMICS
INTRODUCTION The production technology of modern steel grades is based on the production of metal with narrow intervals of chemical composition, alloying elements, modifiers, reducing the content of harmful impurities and non-metallic inclusions. Achieving these parameters requires fine-tuning of steelmaking technologies at each stage, taking into account changes in the temperature and composition of the steel melt and slag and the of additives introduction mode. Modern metallurgical technologies of the XXI century provide various methods of ladle processing to control the quality of steels and alloys. All industrial experiments on technology optimization are complex and extensive. The best way is a computational modeling of metallurgical technologies. The modelling of metallurgical processes is a difficult problem that requires the development of physical - chemical models and mathematical algorithms, allowing adequate description of high-temperature processes occurring in open non-equilibrium systems. The most of computer software, that modelling the real metallurgical process are based on approximating and statistical models demanding enormous numbers of experimental data [1-3]. This fact essentially limits possibilities of the software, which aren’t capable to sufficiently react to various disturbance and random processes in a wide range of parameters change. Using of computational models that adequately simulate the processes during ladle treatment of steel allows us to calculate the optimal technology for the production of a certain steel grades simulates this on a computer without a series of costly industrial experiments, develop new technolo20
gies for steel production and identify new factors that affect product quality. The aim of this study was to develop of mathematical models, algorithms and software for dynamic simulation of steel treatment technology in ladle furnace (LF). METHODS AND MODELS Software for dynamic simulation of ladle treatment technology was based on the physical and chemical models and thermodynamics models [4-5]. The target of this software was modelling and on-line control of steel temperature and chemical composition of slag and steel melt during steelmaking processes (ladle-furnace). Physical and chemical models based on mass and energy conservation law and principals of nonequilibrium thermodynamics were used. All process stages (zo-
O.A. Komolova, K.V. Grigorovich
Baikov Institute of Metallurgy and Material Science Ras, Moscow, Russia
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Metallurgia fuori forno nes) were taken into account in this software. It was assumed that the metallurgical systems do not reach equilibrium and are in non equilibrium steady states. In accordance with the L. On-
sager, it was proposed that the reaction rate is proportional to the gradient of the chemical potential according the formulae:
[1]
where Vi – reaction speed of i -component, mol/s; S- interaction surface, m2; L- Onsager’s coefficient, mol2/(J·s·m); gradμi- gradient of the chemical potential of i component, J/(mol·m). All components in the interaction zone in the slag-metal system are equal to the turbulent mass transfer conditions. Therefore it was assumed that the surface area of interaction, Onsager coefficients, temperature and boundary layer thickness - δ are the same for all reactions. If the coefficient ß is: [2]
Than reaction speed of i –component is: [3]
Where Ke and Kr are equilibrium and real reaction constants. To calculate the reaction rates of interactions between components of slag-metal system was developed by an iterative algorithm. Model defines a direction of chemical reactions for metal-slag system witch presented as a matrix of k reactions and takes into consideration mass and energy balance equations. All interaction zones are described by deterministic rather than statistical dependencies, models are stable over a wide range of variables and it is stable even after changes of technology. This software takes into consideration input data such as: temperature, slag and metal mass and compositions, input time and mass of additives, blowing, electrical and time regimes. Additional data which is to be used in calculations are ladle equipment facilities (ladle geometry, transformer parameters, electrode consumption, number of lances, type of refractory materials), thermodynamic database, thermal, physical and chemical databases for additives and inert gas, production database (for statistics). It was demonstrated that software is stable even after changes in technological scheme.
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Mathematical model consists of the following blocks: - Calculation the speed of interaction between the components in the slag-metal system; - Calculating the amount of metal and slag in the interaction zone depending on the power of stirring of the bath; - Calculation of the mass of metal and slag; - Calculation of the chemical composition and temperature of the slag and metal bath. Calculations of energy balance for metal-slag system and in all areas including arc heating and takes into consideration heats of chemical reactions; Calculations of heat of metal and slag melts, alloying elements and fluxes; Heat loss calculations through the lining by radiation, for heating the inert gas and the reacting components at the boundary of the slag - metal lining ; Calculations of nonmetallic inclusions formation and removal. The Figure 1 represents of the calculating scheme of the LadleFurnace software package.
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Fig. 1 – Calculating scheme of the Ladle Furnace software package RESULTS AND DISCUSSION For testing of Software and validation of the model, the results of ladle treatment of 25 real heats of steel for pipe line and sampling control results were used. Comparative results of calculated values obtained by the software and results of chemical composition control of metal melt during treatment at the ladle furnace 165 t., presented on the Fig. 2 Fig 3 presented of comparative results of chemical composition control of the samples of molten metal during processing at the 355 tonn ladle furnace. It was shown that Software designed allows us to make dynamic simulation and optimization of ladle
22
treatment technology. It was established that the software to adequately describe the dynamic changes of the basic characteristics of the metal, slag and reaction of system on the process control feedback. It was shown that the software developed for the dynamic simulation of ladle treatment of steel and fractional gas analysis method allows to optimize the secondary treatment technology. This software can be used in online calculations and control of process parameters during ladle treatment, modelling and optimization of ladle treatment technology, teaching and training of steelmaking staff.
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Metallurgia fuori forno
Fig. 2 – Comparative results of calculations by the LF software and obtained results of chemical composition control of metal melt during ladle treatment at the ladle furnace 165 tonn
Fig. 3 – Comparative results of calculations by the LF software and obtained results of chemical composition control of metal melt during ladle treatment at the ladle furnace 355 tonn La Metallurgia Italiana - n. 3 2019
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Secondary refining CONCLUSIONS The original software for dynamic simulation of ladle treatment of steel was developed. Physical and chemical models based on the conservation law of mass and energy, as well as the principles of non equilibrium thermo-dynamics were used. All the stages of the process (zones) were taken into account in this
software. Software for dynamic simulation of ladle treatment technology to increase the steels quality was used in simulating of real in-dustrial hits. It was shown that software designed allows us to provide the dynamic simulation of ladle treatment technology, to optimize one and to lead the process within an optimal way.
REFERENCES [1]
Tsymbal V.P. Mathematical modeling of complex systems in metallurgy: a textbook for high schools. - Kemerovo; M .: Publishing Association “Russian Universities”: Kuzbassvuzuzdat-ASTSh, 2006. - 431 p.
[2]
D'yachko A.G. Mathematical and simulation modeling of production systems. –Moscow: MISiS, 2007, 538 p.
[3]
Sovetov B.YA., Yakovlev S.A. Modeling systems. - M .: Higher School, 2001. 343 p.
[4]
Komolova, O.A., Modeling of the components interaction of slag and metal phases in the production of steel, development of algorithms and software for the processes description, Extended Abstract of Cand. Sci. (Tech.) Dissertation, Moscow: Moscow Inst. Steel Alloys, 2014
[5]
Konstantin Grigorovich, Olga Komolova, Darina Terebikina: "Analysis and optimization of ladle treatment technology of steels processing", Journal of Chemical Technology and Metallurgy, 50, 6, 2015, 574-580
24
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Metallurgia fuori forno
Termodynamics and kinetics for the evolution of non-metallic inclusions in pipeline steel Y. Zhang, Y. Ren, L. Zhang
In the current study, the thermodynamics for the complex deoxidation of the pipeline steels was established and discussed, comparing to the deoxidation of pure iron. A remarkable difference in the deoxidation curve bretween real pipeline steel and pure iron was found since for real steels complex inclusions, such as calcium-aluminate, alumina and sulfide, were generated during deoxidation rather than pure alumina with the increasing of the dissolved aluminum in steel. A kinetic model was developed to predict the composition variation with time of non-metallic inclusions, steel and slag during LF refining of pipeline steels. Reactions between alloy elements and steel generating inclusions, slag and steel, between slag and lining refractory, between steel and lining refractory were considered in this kinetic model, considering air absorption and the removal of inclusions. The effects of inclusion size and gas flow rate during LF refining on the composition variation with time of inclusions, steel and slag were investigated. The kinetic model showed a good agreement with industrial measurement of the composition of inclusions, steel and slag.
KEYWORDS: THERMODYNAMICS – KINETICS – EVOLUTION – NON-METALLIC INCLUSIONS – PIPELINE STEELS
INTRODUCTION The composition, size, number, and distribution of inclusions in linepipe steels have a huge influence on the performance the linepipe steels. [1, 2] The solid inclusions may cause the clogging of the submerged entry nozzle.[3-5] The indeformable inclusions in linepipe steels can lead to the hydrogen induced cracking.[6-8] The control of inclusions is the key task of the steelmakers. There have been a large number of studies on the control of inclusions by lab experiments and plante trials.[9] Traditionally, it is suggested that inclusions in linepipe steels should be modified to liquind calcium aluminates at steelmaking temperature.[10, 11] Meanwhile, the modification machnism of the Al2O3 inclusions in Al-killed steel was widely investigated. It is found that the Al2O3 inclusions were directly reduced by dissolved calcium or indirectly reduced via CaO formed from [Ca] and [O]. Inversely, it is also reported that CaS phase in inclusions are transiently formed immediately after the calcium addition, then decreases and disappears. [12-16] However, it is found that the fully liquid CaO-Al2O3 are in large size and easier to be rolled to stringer shaped inclusions.[17] Thus, to avoid the formation of the liquid CaO-Al2O3 inclusions, it is proposed that adding the insufficient or superfluous calcium can modify inclusions to partially liquid CaO-Al2O3 inclusions, which can hardly lead to the nozzle clogging and the formation of the hydrogen induced cracking. To achieve the better controlling of inclusion composition after the reaction equilibium, many efforts
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has been carried out to predict the deoxidation and inclusion formation in liquid pure iron by thermodynamic calculation.The deoxidation equilibrium relationship by single deoxidizer could be maturely predicted by thermodynamic calculation, such as Al-O[18, 19], Mg-O[20], Ca-O[21, 22], etc. The thermodynamic data of alloy elements in the molten steel were summarized and suggested by JSPS.[23] To predict the formation of complex inclusions such as Mg-Al-O, Al-Ti-O, Al-Ca-O, Al-Si-O, Mg-AlCa-O, etc., the experimets and thermodynamic calculation for the pure iron with the addition of multi-deoxidizers have been conducted.[24-29] Especially, the application of thermodynamic databases was introduced to steelmaking process[30, 31],
Ying Zhang, Ying Ren, Lifeng Zhang
School of Metallurgical and Ecological Engineering, University of Science and Technology Beijing Beijing 100083, China email: zhanglifeng@ustb.edu.cn
25
Secondary refining which significantly improve the accuracy of the thermodynamic prediction. However, most of the previous studies focus on the inclusion formation in the liquid iron, insteady of the real steel. Thus, the thermodynamics for the complex deoxidation of the pipeline steels was established and discussed in the current study. To precisely predict composition variation with time of non-metallic inclusions, steel, and slag, some kinetic model has been developed. Harada et al[32-34] developed a kinetic model to calculate composition changes in the molten steel, the slag, and the inclusion in ladle refining. The coupled reaction model was applied to the reaction between the molten steel/slag phase and the molten steel/inclusion originating from the slag. Empirical equations were used to calculate the rate of MgO dissolution in the slag from the refractory. Peter et al developed a model to model and simulate the ladle refining at two different LMF’s, refining Al-killed and Si-deoxidized steels using the process simulation program Metsim and the thermodynamical program FactSage.[35] A comprehensive model for inclusion development in gas stirred ladles developed by the Shu and Scheller, which is a useful tool for simulation and optimization of ladle metallurgical treatments in industry.[36] Jung et al applied a the effective equilibrium reaction zone model using the FactSage macro processing code to develop a kinetic LF process model.[37] Shin et al developed a refractory–slag– metal–inclusion multiphase reaction model by integrating the refractory–slag, slag–metal, and metal–inclusion elementary reactions in order to predict the evolution of inclusions during the secondary refining processes.[38] Piva et al simulated steelslag and steel-inclusion reaction kinetics in silicon-manganese killed steels using FactSage macros.[39] Ying et al developed kinetic models using FactSage Macro Processing to simulate
the composition evolutions of slag, steel, and inclusions during the calcium treatment, slag refining, and reoxidation processes. [40-43] In the current study, a kinetic model was developed to predict the composition variation with time of non-metallic inclusions, steel and slag during LF refining of pipeline steels with the consideration of reactions between alloy elements and steel generating inclusions, slag and steel, between slag and lining refractory, between steel and lining refractory
(a) MgO·Al2O3
(b) CaO·Al2O3
THERMODYNAMICS FOR INCLUSIONS IN PIPELINE STEELS The formed typical inclusions in linepipe steels are shown in Fig.1. During the production prcess of linepipe steels, the Al2O3 inclusions formed after the Al deoxidation. As the reduction of MgO in slag and refractory, the Mg transfers to the moten steel and react with the inclusions. The formed typical MgO·Al2O3 inclusion is shown in Fig.1(a). After the modification of the slag refining and the calcium treatment, the inclusions are modified to liquid calcium aluminate at steelmaking temperature, as shown in Fig.1(b). Fig.2 shows the equilibrium relation of oxygen and aluminum contents in steel at 1873 K. In Fig.2(a) shows the Al-O curve in pure iron at 1873 K. With an increase of Al in the liquid iron, the Al content decreases first and then goes up. The minimum concentration of O in the liquid iron can be lowered to 3 ppm. Fig.2(b) shows the Al-O curve in the linepipe steel at 1873 K. The oxygen content can be decreased to several ppm with the addition of Al. The formed inclusions are solid 2CaO·SiO2 with less than 0.01% Al, while the inclusions are liquid at steelmakine temperature with more than 0.01% Al. There is an obvious difference in liquind iron and the linepipe steels, as shown in Figs.2(a) and (b).
Fig. 1 – The formed MgO·Al2O3 and CaO·Al2O3 inclusions in linepipe steels. 26
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(a) Al-O curve in pure iron. [9]
(b) Al-O curve in linepipe steels. [15]
Fig. 2 – The equilibrium relation of oxygen and aluminum contents in pure iron and linepipe steels at 1873 K. Fig.3 (a) shows the equilibrium relation of oxygen and calcium contents in pure iron under various S concentrations at 1873 K. With the increase of Ca in liquid iron, the oxygen content decrease first, then goes up. The oxygen content can be lowered to 2 ppm. The addition of S lowers the deoxidation ability of the Ca in liquid iron. Fig.3 (b) is the equilibrium of Ca-O curve in linepipe steels at 1873 K using Factsage with the FactPS and FToxid and FTmisc databases. The evolution route of oxide pha-
se is CaO·6Al2O3 → CaO·2Al2O3 → liquid inclusion → CaO. The oxygen can be decreased to roughly 5 ppm. With S increasing from 0 ppm to 200 ppm, there is a slight decrease of the oxygen content due to the transformation of the formed oxides. Since the interaction between O and S is much smaller than that between Ca and O , the relationship of the equilibrated oxygen content and the calcium content could hardly be influenced by S content.[44]
(a) Ca-O curve in pure iron. [9]
(b) Ca-O curve in linepipe steels. [15]
Fig. 3 – The equilibrium relation of oxygen and calcium contents in pure iron and linepipe steels at 1873 K. The effect of oxygen content on the equilibrium of sulphur and calcium contents in steel at 1873 K is shown in Fig.4. Fig.4 (a) shows the Ca-S curve in pure iron with various O. It can be seen that the S content decreases with the increasing Ca in liquind iron. Meanwhile, the S content obviously increases with the increasing O content in the liquid iron. Meanwhile, the Ca-S curve in pure iron is shown in Fig.4 (a). In linepipe steels, the La Metallurgia Italiana - n. 3 2019
addition of Ca can significantly decrease the S content. As the dissolved oxygen increases, the desulphurization ability of calcium is weakened obviously. The inflection points mainly caused by the oxide phase transformation. It can be inferred that the O content has a huge influence on the equilibrium relationship between sulphur and calcium contents in the molten steel.
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(a) Ca-S curve in pure iron. [9]
(b) Ca-S curve in linepipe steels. [15]
Fig. 4 – The equilibrium relation of oxygen and calcium contents in pure iron and linepipe steels at 1873 K. The calculated stability diagram of Al-Ca-O system in the pure iron at 1873K is shown in Fig.5(a).[28] The formation of Al2O3, CaO·6Al2O3, CaO·2Al2O3, CaO·Al2O3, liquid calcium aluminates, and CaO nclusion phases are considered in the current calculation. The Al2O3 and CaO are severally formed in the pure iron with additions of superfluous O and Ca. The calcium aluminates are formed between Al2O3 and CaO inclusions region. Fig.4 (b) shows the calculated Al-Ca-O inclusion diagram in linepipe steels at 1873K.[15] With the addition of Ca in line-
pipe steels, the inclusions evolve from Al2O3 to CaO·6Al2O3 to CaO·2Al2O3, liquid calcium aluminates. The CaS and 2CaO·SiO2 are formed with the superfluous Ca. A remarkable difference in the deoxidation curve bretween real pipeline steel and pure iron was found since for real steels complex inclusions, such as calcium-aluminate, alumina and sulfide, were generated during the deoxidation rather than pure alumina with the increasing of the dissolved aluminum in steel.
(a) Pure iron[28]
(b) Linepipe steels[15]
Fig. 5 – Stability diagram of inclusions in Al-Ca-killed steel. KINETICS FOR EVOLUTION OF INCLUSIONS IN PIPELINE STEELS The basic reactions between steel-slag-refractory-alloy-inclusion-air during refining process was decribed in the previous study, as the schematic diagram shown in Fig. 6.[45] The reactions of stee-slag, slag inclusions, alloy-steel, air-steel, refrac28
tory-steel, refractory-slag are considered in the current model. The reactions and the thermodynamic data can be found in our previous work. The reactions at interface were assumed to be equilibrated at any time. Then, the mass transfer is the ratecontrolling step. Therefore, the reaction rate is mainly determined by the mass transfer coefficients. La Metallurgia Italiana - n. 3 2019
Metallurgia fuori forno
Fig. 6 – Schematic diagram of multiphase reactions during refining process[45] In the current kinetic model, the empirical equations of the mass transfer coefficient calculation were corrected as following Eqs. [1]-[4].
[1] [2] [3]
[4]
where km and ks are the mass transfer coefficients of molten steel and slag, respectively, m/s; km_inc and kinc are the mass transfer coefficients of steel and inclusion during steelinclusion reaction, m/s; ɛ is the effective stirring power, (W/t); T is the temperature of molten steel and inclusions, (K); D is the diffusion coefficient of elements in the molten steel, m2/s; uslip respects the relative velocity of steel and inclusion, m/s; dp is the inclusion diameter, μm; TS and TL are severally the solidus and liquidus of the inclusion, K. A plant trial of ladle refining process of linepipe steels was simulated using the current developed model. The initial slag and steel are mainly CaO-Al2O3-SiO2-MgO and Al-killed linepiLa Metallurgia Italiana - n. 3 2019
pe steels. The initial inclusions in the molten steel are 2μm solid Al2O3-rich inclusions. During the refining process, the lime was added to increase the CaO in the refining slag. To validate the accuracy of the current model, the calculated and the measured T.S evolutions were compared in Fig.7. The predicted results shows a good agreement with the experimental ones, indicating that the current model can be used to predict the reactions during the ladle refining process. Fig.8 shows the predicted evolution of slag compositions during the ladle refining process. It can be seen that there is an obvious increase in CaO content with the addition of lime. Fig.9 shows the volution of T.Ca and T.Mg during the ladle re29
Secondary refining fining process. It can be seen that the addtion of CaO leads to the increase of T.Ca. Then, the T.Ca and T.Mg slowly goes down during the refining process, which may be caused by the floatation of Mg-containing inclusions. Fig.10 shows the evoltion of the CaO and
Al2O3 in inclusions in the molten steel. The initial inclusions are mainly Al2O3. With the slag refining process, the Al2O3 inclusions are modified to liquid calcium aluminate. It can be inferred that the slag refining has an obvious influence on the inclusions composition.
Fig. 7 – Schematic diagram of multiphase reactions during refining process
Fig. 8 – Predicted evolution of slag composition during the ladle refining process
Fig. 9 – Predicted evolution of steel composition during the ladle refining process 30
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Fig. 10 – Predicted evolution of inclusion composition during the ladle refining process To investigate the effect of inclusion size on inclusion composition evolution during the ladle refining process, the evolutions of 4.5μm, 15.5μm, and 95.5μm inclusions are compared in Fig.11. From the evolusions of Al2O3, CaO, and CaS in inclusions. It can bee seen that the small inclusions are modified and almost unchanged after roughly 900 seconds. However, the 95.5μm inclusions can hardly reach the equilibrium after
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2400 seconds. Meanwhile, the CaS can form by the slag refining without the calcium treatment. The MgO in inclusions slightly increases with the dissolution of MgO refractory. It is indicated that the small inclusions can be easily infuenced by the slag-steel-inclusion reactions, while the large size inclusions would keep their main composition during the refining process.
31
Secondary refining
Fig. 11 – Predicted effect of inclusion size on inclusion composition evolution during the ladle refining process
CONCLUSIONS (1) A remarkable difference in the deoxidation curve bretween real pipeline steel and pure iron was found since for real steels complex inclusions were generated during deoxidation rather than pure alumina with the increasing of the dissolved aluminum in steel. (2) A kinetic model was developed to predict the composition variation with time of non -metallic inclusions, steel and slag during LF refining of pipeline steels. The reactions between alloy elements and steel generating inclusions, slag and steel, between slag and lining refractory, between steel and lining refractory were considered. (3) The small inclusions can be easily infuenced by the slagsteel-inclusion reactions, while the large size inclusions would
32
keep their main composition during the refining process. while the large size inclusions would keep their main composition during the refining process. ACKNOWLEDGEMENTS The authors are grateful for support from the National Natural Science Foundation of China (Grant No. 51725402 and No. 51704018), National Postdoctoral Program for Innovative Talents (Grant No. BX201700028), Young Elite Scientists Sponsorship Program By CAST (No. 2017QNRC001), China Postdoctoral Science Foundation (No.2017M620016) and the High Quality steel Consortium (HQSC) at University of Science and Technology Beijing (USTB), China.
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Secondary refining [24] Itoh, H., M. Hino, and S. Ban-Ya, Thermodynamics on the Formation of Spinel Nonmetallic Inclusion in Liquid Steel. Metallurgical and Materials Transactions B, 1997. 28(5): p. 953-956. [25] Jung, I.-H., et al., Thermodynamic Modeling of the Al2O3–Ti2O3–TiO2 System and Its Applications to the Fe–Al–Ti–O Inclusion Diagram. ISIJ International, 2009. 49(9): p. 1290-1297. [26] Itoh, H., M. Hino, and S. Banya, Thermodynamics on the Formation of Non-metallic Inclusion of Spinel (MgO·Al2O3) in Liquid Steel. Tetsu-to-Hagane, 1998. 84(2): p. 85-90. [27] Taguchi, K., et al., Complex Deoxidation Equilibria of Molten Iron by Aluminum and Calcium. ISIJ International, 2005. 45(11): p. 1572-1576. [28] Ren, Y., et al., Detection of Non-metallic Inclusions in Steel Continuous Casting Billets. Metallurgical and Materials Transactions B, 2014. 45(4): p. 1291-1303. [29] Kang, Y.-B. and H.-G. Lee, Inclusions Chemistry for Mn/Si Deoxidized Steels: Thermodynamic Predictions and Experimental Confirmations. ISIJ International, 2004. 44(6): p. 1006-1015. [30] Jung, I.-H., S. Decterov, and A. Pelton, A thermodynamic model for deoxidation equilibria in steel. Metallurgical and Materials Transactions B, 2004. 35(3): p. 493-507. [31] Jung, I.-H., Overview of the applications of thermodynamic databases to steelmaking processes. CALPHAD: Computer Coupling of Phase Diagrams and Thermochemistry, 2010. 34(3): p. 332-362. [32] Harada, A., et al., Kinetic Analysis of Compositional Changes in Inclusions during Ladle Refining. ISIJ International, 2014. 54(11): p. 2569-2577. [33] Harada, A., et al., A Kinetic Model to Predict the Compositions of Metal, Slag and Inclusions during Ladle Refining: Part 2. Condition to Control the Inclusion Composition. ISIJ International, 2013. 53(12): p. 2118-2125. [34] Harada, A., et al., A Kinetic Model to Predict the Compositions of Metal, Slag and Inclusions during Ladle Refining: Part 1. Basic Concept and Application. ISIJ International, 2013. 53(12): p. 2110-2117. [35] Peter, J., et al. Experimental Study of Kinetic Processes During the Steel Treatment at two LMF's. in AISTech 2005 Proceeding. 2005. Charlotte, NC. [36] Scheller, P.R. and Q. Shu, Inclusion Development in Steel During Ladle Metallurgical Treatment- A Process Simulation ModelPart: Industrial Validation. Steel Research International, 2014. 85(8): p. 1310-1316. [37] Van Ende, M.-A. and I.-H. Jung, A Kinetic Ladle Furnace Process Simulation Model: Effective Equilibrium Reaction Zone Model Using FactSage Macro Processing. Metallurgical and Materials Transactions B, 2017. 48(1): p. 28-36. [38] Shin, J.H., Y. Chung, and J.H. Park, Refractory–Slag–Metal–Inclusion Multiphase Reactions Modeling Using Computational Thermodynamics: Kinetic Model for Prediction of Inclusion Evolution in Molten Steel. Metallurgical and Materials Transactions B, 2017. 48(1): p. 46-59. [39] Piva, S.P.T., D. Kumar, and P.C. Pistorius, Modeling Manganese Silicate Inclusion Composition Changes during Ladle Treatment Using FactSage Macros. Metallurgical and Materials Transactions B, 2017. 48(1): p. 37-45. [40] Ren, Y., L.-f. Zhang, and Y. Zhang, Modeling reoxidation behavior of Al–Ti-containing steels by CaO–Al2O3–MgO–SiO2 slag. Journal of Iron and Steel Research International, 2018: p. 1-11(Published online). [41] Ren, Y. and L. Zhang, Modeling Inclusion Evolution in Al-Ti-killed Steels during Ladle Mixing Process. Ironmaking & Steelmaking, 2017: p. Accepted. [42] Huang, F., et al., Kinetic Modeling for the Dissolution of MgO Lining Refractory in Al-Killed Steels. Metallurgical and Materials Transactions B, 2017. 48(4): p. 1-12. [43] Ren, Y., Y. Zhang, and L. Zhang, A kinetic model for Ca treatment of Al-killed steels using FactSage macro processing. Ironmaking & Steelmaking, 2017. 44(7): p. 497-504. [44] Taguchi, K., et al., Deoxidation and Desulfurization Equilibria of Liquid Iron by Calcium. ISIJ Iniemational, 2003. 43(11): p. 1705–1709. [45] Zhang, Y., Y. Ren, and L. Zhang. Kinetic variation of the composition of inclusions, steel and slag during refining of pipeline steel. in AISTech 2017 Iron and Steel Technology Conference, May 8, 2017 - May 11, 2017. 2607-2613.
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Corrosione
Resistance to localized corrosion of lean duplex stainless steels after brief thermal treatments. F. Zanotto, V. Grassi, A. Balbo, F. Zucchi, C. Monticelli
Lean duplex stainless steels are biphasic steels with a low nickel and/or molybdenum content, which represent an interesting and cost effective alternative to the employment of the most common austenitic stainless steels (such as AISI 304 and 316) or traditional more expensive duplex steels (such as DSS 2205). However, during production and welding operations, a relatively brief permanence (a few minutes or tens of minutes) in the critical temperature range between 650 and 950 °C, induces the formation of undesirable secondary phases (mainly chromium and molybdenum carbides and nitrides) which influence the alloy corrosion resistance. This paper compares the pitting corrosion resistance of three lean duplex alloys, LDSS 2101, DSS 2304 and LDSS 2404, before and after heat treatments of 10 min at 650, 750 and 850 °C and analyses the correlated pitting propagation modes. With this aim, critical pitting temperature (CPT) measurements were performed in 0.1 M NaCl solution. The results show that concerning the solubilized alloys their localized corrosion resistance is in good agreement with PRENMn indications. Thermal treatments, and in particular that at 750 °C, reduce the pitting corrosion resistance of the alloys, mainly in the case of LDSS 2404 and 2101, due to the impoverishment in passivating elements (especially chromium and molybdenum) in regions close to precipitates. The heat treated DSS 2304 alloy shows a more limited reduction of its performances in comparison to the solubilized material and its CPT does not vary significantly with the treatment temperature. This phenomenon can be linked to the low nitrogen content of the alloy, which limits the volume of precipitates during heat treatments, allowing to reduce the extent of impoverishment in passivating elements.
KEYWORDS: STAINLESS STEEL – LEAN DUPLEX – THERMAL TREATMENT – PITTING CORROSION – CRITICAL PITTING TEMPERATURE
INTRODUCTION Duplex stainless steels (DSS) are steels with austeno-ferritic microstructure with a relevant applicative interest as they have a favourable combination of mechanical properties, weldability and high corrosion resistance in different environments. These characteristics make these alloys appropriate for different field conditions: from the paper industry to the petrochemical one, from the construction sector to that of nuclear energy production [1]. In the last 20 years, the research of new duplex alloys followed two main directions: from one side the improving of corrosion resistance by increasing the alloyed content of chromium, molybdenum and nitrogen (superduplex and hyperduplex stainless steels) and, from the other side, the market introduction of a new duplex family, called lean duplex stainless steels (LDSS) [2], still rather corrosion resistant, but more cost effective. The first lean duplex placed in the market was DSS 2304 (UNS 32304), with a low molybdenum content, which permitted both the prevention of detrimental secondary phases formation (χ and σ phases) during welding operations and production cost savings [1]. Subsequently, since the years 2000, the development of new lean duplex alloys focused on the reduction of nickel content, as this element, like molybdeLa Metallurgia Italiana - n. 3 2019
num, is significantly subject to market price fluctuations. With this aim, alloys were developed with nitrogen contents close to the element solubility limit and with high manganese levels, added to increase nitrogen solubility and to stabilize austenite phase. In this context, LDSS 2101 (UNS S32101) was introduced in the market with the aim of advantageously substituting the most frequently used austenitic stainless steels (AISI 304 e
F. Zanotto, V. Grassi, A. Balbo, F. Zucchi, C. Monticelli
Centro di Studi sulla Corrosione e Metallurgia “Aldo Daccò”Università degli Studi di Ferrara, via Saragat 4A, 44122 Ferrara
35
Corrosion 316), in field conditions where high mechanical properties and significant localized corrosion performance are required [3-6], but also carbon steels where maintenance costs are relevant [1]. In more recent years (year 2010), again in the field of low nickel and high manganese content duplex stainless steels, LDSS 2404 (UNS S82441) was developed. This alloy, containing higher chromium, molybdenum and nitrogen levels with respect to LDSS 2101 and DSS 2304, is expected to show a better localized corrosion resistance compared to the latter ones. In particular, the PREN (Pitting resistance Equivalent Number) value of LDSS 2404, and also its mechanical properties, are very similar to those of the traditional DSS 2205 (UNS S32205), with lower costs. Currently, DSS 2205 is widely used in aggressive environments, such as chemical and desalination plants and in the petrochemical industry, but it is also successfully used in architectural and civil engineering works, such as the Marina Bay pedestrian bridge in Singapore [7] or the Millenium Bridge in York [8]. In general, the possibility of reducing maintenance costs, guaranteeing the aesthetic appearance of the material for a very long time, lightening the structures by maintaining or even improving the mechanical resistance and, finally, the possibility of buying, transporting, welding and assembling more limited quantities of material, as a result of weight reduction, so allowing manpower cost savings, make LDSS a beneficial alternative to the use of the austenitic stainless steels and carbon steels most commonly used in construction field and civil engineering works [7] The optimization of the mechanical properties and the corrosion resistance of duplex stainless steels, is obtained by ensuring a ferrite (α)/austenite (γ) ratio close to 1 and avoiding detrimental secondary phases formation [9]. However, during production and welding operations, or under high temperature conditions, excessive permanence in the critical temperature range between 650 and 950 ° C, can determine (depending on the chemical composition of the alloy) the growth of χ and σ phases and/or carbides and nitrides mainly of chromium. These phases can have a negative effect on the mechanical performances and localized corrosion resistance of the alloy [10].
The authors of this research study carried out several studies dealing with the resistance to localized corrosion and stress corrosion cracking of LDSS, before and after thermal treatments between 650 and 850 ° C [11-18]. In particular, it was found that aging for 5-30 min in the 650 - 850 °C temperature range causes the formation of mainly chromium nitrides at LDSS 2101 grain boundaries, due to the high nitrogen content (0.22 wt%) of the alloy [13-15]. Instead, in DSS 2304 aging for up to 60 min in the same temperature range essentially determines chromium (and molybdenum) carbide precipitation, because the nitrogen content is much lower (0.1 wt%) [16,18]. These microstructural modifications tend to reduce the alloy localized corrosion resistance, as they can cause depletion in passivating elements (such as chromium and molybdenum) around the precipitates. This paper aims at comparing the effect of 10 min aging in the temperature range between 650 and 850 ° C on the microstructure and pitting corrosion resistance of LDSS 2101, DSS 2304 and LDSS 2404 in a 0.1 M NaCl solution. The different microstructures and phase compositions also affected the pitting attack morphology. MATERIALS AND METHODS The experimental tests were carried out on LDSS 2101, DSS 2304 and LDSS 2404 alloys (supplied by Outokumpu in the form of sheets under solubilized conditions) with the nominal chemical compositions shown in Table 1. In the same table, the PREN values of each alloy are showed. This index, specifically developed to predict the pitting corrosion resistance of stainless steels [19], is considered an easy way to estimate the corrosion resistance of stainless steels from their chemical composition and currently it is also applied to various DSS. [20]. Usually, the PREN is calculated on the basis of the Cr, Mo and N contents of the alloys, by using the formula PREN =% Cr + 3.3 ·% Mo + 16 ·% N [19]. The PREN values in Table 1 suggest that LDSS 2404 is the most pitting corrosion resistant alloy, due to its high Cr, Mo and N contents, while DSS 2304 and LDSS 2101 are expected to afford almost equivalent resistance to localized corrosion.
Tab. 1 – Nominal chemical composition (wt.%) and PREN values (PREN = %Cr + 3.3·%Mo + 16·%N [32]) of the studied lean duplex stainless steels.
DSS
C
Mn
Cr
Ni
Mo
N
Fe
PREN
LDSS 2404
0.02
3.0
24
3.6
1.6
0.27
bal.
34
DSS 2304
0.02
-
23
4.8
0.3
0.10
bal.
26
LDSS 2101
0.03
5.0
21
1.5
0.3
0.22
bal.
26
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Corrosione The electrodes for electrochemical measurements had a surface area of about 0.45 cm2. They were cut from sheets of the three alloys. Some of them were thermally aged for 10 min at 650, 750 and 850 °C and then cooled in air. The microstructures, before and after the heat treatments, were documented by optical microscope (OM) observations, after metallographic etching with the Beraha’s reagent, and by scanning electron microscope (SEM), using backscattered electrons (BSD). Finally, the elemental microanalysis with EDS microprobe were used to study the composition of the secondary phases formed during the heat treatments. The pitting corrosion resistance was assessed by CPT (critical pitting temperature) measurements in 0.1 M NaCl solution. Before immersion of the working electrode, the solution was thermostated at 5 ° C. Then, in order to increase the reproducibility of the measurements [21], the working electrode was cathodically polarized at -0.9 VSCE for 5 min to reduce the air-formed surface oxide film. Subsequently, the electrode was left at the open circuit potential (EOCP) for 30 minutes and finally it was
anodically polarized to + 0.75 mVSCE, while the electrolyte temperature was increased by 1 °C / min [22]. The CPT was evaluated as that temperature at which the current exceeded 100 μA / cm2 and the test was stopped when the current reached values of 250÷300 μA / cm2. Each mean CPT value was determined as the average of 3 trials. The pitting corrosion morphology was observed by SEM with BSD electrons. RESULTS Microstructure The microstructures of as-received LDSS 2101, DSS 2304 e LDSS 2404 samples are very similar. As an example, the microstructure of the short transversal section of DSS 2304, obtained with OM after Beraha’s reagent etching, is shown in Figure 1. The image evidences the ferritic matrix (darker phase) embedding flattened austenitic islands (lighter phase), produced by sectioning the elongated austenitic grains (rolling direction perpendicular to the shown section).
Fig. 1 – OM microstructure of the transversal section (perpendicular to rolling direction) of the as-received DSS 2304 (etching with Beraha’s reagent).
As explained in the introduction, the chemical composition of the studied alloys determines the nature of the secondary phases which can grow during the thermal treatments, while the times and temperatures (between 650 and 850 °C) of aging influences their dimension and distribution at the grain boundaries. In general, with a 10 min heat treatment a more significant precipitation is detected by increasing the ageing temperature from 650 to 850 °C [13-16,18], so that the precipitates are better distinguished at the highest temperature. After 10 min at 850 °C, the microstructures of the three studied lean duplex alloys are shown in Figure 2. The precipitates, highlighted by red arrows, are more abundant in LDSS 2101 and 2404 than La Metallurgia Italiana - n. 3 2019
in DSS 2304 and are localized at both α/α grain boundaries and α/у interphases. These precipitates are mainly constituted by chromium nitrides in the case of LDSS 2101 and essentially chromium (and molybdenum) carbides for DSS 2304 [1316,18]. In the case of LDSS 2404, Figure 3 presents the EDS elemental line analysis through these secondary phases which are evidently constituted by chromium nitrides. Therefore, the different volume fraction of precipitates in the alloy microstructures of Figure 2 is connected to the different N content in the alloys which is lower in 2304, where mainly carbides are formed, and higher in LDSS 2101 and 2404, where abundant nitride precipitation is observed. 37
Corrosion
Fig. 2 – SEM-BSD microstructures of the transversal sections of LDSS 2101, DSS 2304 e LDSS 2404, after aging of 10 min at 850 °C.
Fig. 3 – SEM-BSD microstructure of the transversal section of LDSS 2404 and EDS profile-line analysis of chromium, carbon and nitrogen.
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Corrosione CPT results Figure 4 shows the current density/temperature curves recorded on as-received and 850 °C aged LDSS 2101, DSS 2304 and LDSS 2404 electrodes, during polarization at +0.75 VSCE in 0.1
M NaCl solution, while the histogram comparing all the average CPT values obtained from the current density/temperature curves, is reported in Figure 5.
Fig. 4 – Current density vs. temperature curves obtained in 0.1 M NaCl solution on LDSS 2101, DSS 2304 and LDSS 2404 both asreceived and aged 10 min at 850 °C.
Fig. 5 – Average CPT values and standard deviations determined in 0.1 M NaCl solution for LDSS 210, DSS 2304 and LDSS 2404, before and after thermal aging for 10 min at 650, 750 and 850 °C.
Concerning the as-received alloys, LDSS 2404 presents the highest CPT value, i.e. for this grade a higher solution temperature is necessary to induce pitting corrosion at an applied potential of +0.75 VSCE. This suggests, in agreement with the high alloy PREN value (Table 1), that LDSS 2404 has a higher resistance to pitting corrosion in comparison to 2101 and 2304 grades. However, even if the PREN values of DSS 2304 and La Metallurgia Italiana - n. 3 2019
LDSS 2101 are equivalent (Table 1), the average CPT of LDSS 2101 is significantly lower (of about 15 °C) than that of DSS 2304. This result is in agreement with SCC susceptibility data of these as received alloys [17] and is likely linked to the significant Mn content (5 wt.%) in LDSS 2101, which negatively affects the resistance to pitting corrosion of stainless steels, according to some authors [23,24]. Therefore, an alternative 39
Corrosion PREN formula has been proposed including the detrimental effect of Mn on localized corrosion resistance. The modified PREN formula (PRENMn = %Cr + 3,3%Mo + 30%N – 1%Mn) correctly evidences that LDSS 2101 (PRENMn = 24) is less resistant to pitting corrosion than DSS 2304 (PRENMn = 27) and much less resistant than LDSS 2404 (PRENMn = 34). In the latter alloy, the negative effect of the high Mn content (3% by weight) is counterbalanced by the higher content of Cr and especially Mo (1.6% by weight), which is particularly useful for increasing the pitting corrosion resistance of stainless steels [24]. After the aging at 650 °C, a decrease in CPT values of about 8-10 °C was observed for all three alloys. By increasing the treatment temperature to 750 °C, a further worsening of the localized corrosion resistance (CPT lowered of about 8 °C) was detected for LDSS 2101 and LDSS 2404, whereas the average CPT of DSS 2304 remained fairly constant with respect to that obtained after aging at 650 °C. This effect is likely related to the low nitrogen content in this alloy which limits the volume of precipitates during the thermal treatments (Figure 2), so allowing to reduce the phenomena of passivating element depletion at the grain boundaries. After the thermal treatment at the highest temperature (850 °C), LDSS 2101 and LDSS 2404 alloys show an improvement in pitting corrosion resistance, with CPT values increased by about 6 °C compared to those obtained after ageing at 750 °C. This effect is most likely due to the rediffusion of chromium and molybdenum in impoverished areas, linked to the higher diffusion rate of these atoms at the higher temperature. On the other hand, DSS 2304 aged at 850 °C presents a CPT value quite comparable with those obtained at 650 and 750 °C. In this alloy, a partial recovery of the corrosion resistance at 850 ° C was only detected after longer treatment times (60 min) [18], suggesting that in DSS 2304 Cr and Mo rediffusion are likely characterized by slower diffusion kinetics.
40
As far as the pit morphologies are concerned, Figure 6 collects the BSD-SEM micrographs obtained on the low-N DSS 2304 (Figure 6a,c) and on one of the high-N alloys (LDSS 2404, Figure 5b,d) at the end of the CPT tests (after reaching a more or less constant maximum current of 250-300 μA / cm2). Each Figure exhibits the pit morphologies developed on the alloys under both as-received and 750 °C aged conditions. The pitting morphology of LDSS 2101 is quite similar to that on LDSS 2404. The as-received DSS 2304 (Figure 6a) shows a large pit with no evidence of preferential propagation into the ferrite phase. Instead, in the as-received LDSS 2404 (Figure 6b) and also in LDSS 2101, the pit seems to propagate mainly in the ferrite phase, which, in these alloys, is less corrosion resistant than the austenite one [25]. This is reasonable, because the high nitrogen content of these alloys, is mostly concentrated in the austenite phase, so determining a marked difference in PREN values of the austenite and ferrite phases. The 650 °C treatment reduced the pitting corrosion resistance of all asreceived alloys, meaning that chromium- and molybdenumdepleted areas originated in the proximity of the precipitates during the thermal treatment. However, the pit morphologies (not presented in this paper) remained quite similar to those in the corresponding as-received samples (Figures 6a,b), likely due to a limited secondary phase formation which determined small variations in alloyed element distribution. The pit morphologies changed after 750 and 850 °C thermal treatments. As examples, Figures 5c and d obtained on DSS 2304 and LDSS 2404, respectively, after 750 °C ageing, show that the pits strictly propagated along α/γ (mainly) and α/α grain boundaries, close to chromium- (and molybdenum-) rich precipitates, suggesting the presence of more continuous precipitates than at 650 °C, inducing passivating element depletion at intergranular regions.
La Metallurgia Italiana - n. 3 2019
Corrosione
Fig. 6 – SEM-BSD images obtained after CPT tests on the transversal sections of DSS 2304 (left) and LDSS 2404 (right) as-received (a and b) and aged 10 min at 750 (c and d). CONCLUSIONS · In the absence and in the presence of thermal treatments, LDSS 2404 alloy exhibits a higher localized corrosion resistance in comparison to DSS 2304 and LDSS 2101 alloys. LDSS 2101 is the most susceptible to localized corrosion due to its significant Mn content. · Brief thermal treatments (10 min) between 650 and 850 ° C determine the formation of secondary phases at α/α and α/γ grain boundaries: mainly chromium nitrides in the case of LDSS 2101 and LDSS 2404 and chromium and molybdenum carbides in the case of DSS 2304. The Cr and Mo depletion in the areas adjacent to these precipitates causes a reduction in pitting cor-
rosion resistance compared to the as-received samples. · By increasing the treatment temperature from 650 to 750 and then to 850 °C, LDSS 2101 and LDSS 2404 show an initial increase and then a decrease in the susceptibility to pitting corrosion. Instead, the same heat treatments determine a worsening of DSS 2304 pitting resistance which is more or less independent of the temperature in the 650-850 °C range. · On as received LDSS 2101 and 2404, pits mainly propagate in the ferrite phase, while on DSS 2304 they propagate in both ferrite and austenite phases. This pit morphology changes after heat treatments at 750 and 850 °C and becomes essentially intergranular.
REFERENCES [1] Charles J, Chemelle P. The history of duplex developments, nowadays DSS properties and duplex market future trends. Proc. 8th “Duplex Stainless Steel conference”, 2010 October 13–15, Beaune, France, p. 1-50. [2] Nilsson J-O, Chai G, Kivisäkk U. Recent development of duplex stainless steels, Proc. 6th “European stainless steel conference”, 2008 June 10-13, Helsinki, Finland, p. 585–590. [3] Olsson J, Snis M. Duplex – a new generation of stainless steels for desalination plants. Desalination, 2007; 205:104–113. La Metallurgia Italiana - n. 3 2019
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[6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
[23]
[24] [25]
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Merello R, Botana FJ, Botella J, Matres MV, Marcos M. Influence of chemical composition on the pitting corrosion resistance of non-standard low-Ni high- Mn–N duplex stainless steels. Corros Sci, 2003; 45:909–921. Deng B, Jiang Y, Juliang X, Sun T, Gao J, Zhang L, Zhang W, Li J. Application of the modified electrochemical potentiodynamic reactivation method to detect susceptibility to intergranular corrosion of a newly developed lean duplex stainless steel LDX2101. Corros Sci, 2010; 52:969–977. Wei Z, Laizhu J, Jincheng H, Hongmei S. Study of mechanical and corrosion properties of a Fe-21.4Cr-6Mn-1.5Ni-0.24N-0.6Mo duplex stainless steel. Mat Sci and Eng, 2008; A497:501–504. Materiali per idee vincenti. Acciai duplex Outokumpu. 1434IT:2. Lönnberg, Finland, November 2010, www.outokumpu.com, 1-24. Alvarez-Armas I, Degallaix-Moreuil S. Duplex Stainless Steels. ISTE Ltd and John Wiley and Sons, Inc., Great Britain and USA, 2009. Zhang Z, Wang Z, Jiang Y, Tan H, Han D, Guo Y, Li J. Effect of post-weld heat treatment on microstructure evolution and pitting corrosion behavior of UNS S31803 duplex stainless steel welds. Corros Sci, 2012; 62:42–50. Zhang Z, Zhao H, Zhang H, Yu Z, Hu J, He L, Li J. Effect of isothermal aging on the pitting corrosion resistance of UNS S82441 duplex stainless steel based on electrochemical detection. Corros Sci, 2015; 93:120–125. Zanotto F, Grassi V, Zucchi F. Pitting corrosion and stress corrosion cracking resistance of a duplex stainless steel LDX 2101® in the presence of chlorides and thiosulphate. La Metallurgia Italiana, 2013; 6-105:37-45. Zanotto F, Grassi V, Balbo A, Monticelli C, Zucchi F. Stress corrosion cracking of LDX 2101® duplex stainless steel in chloride solutions in the presence of thiosulphate. Corros Sci, 2014; 80:205–212. Zanotto F, Grassi V, Merlin M, Balbo A, Zucchi F. Effect of brief heat treatments performer between 650 and 850 °C on corrosion behaviour of a lean duplex stainless steel. Corros Sci, 2015; 94:38-47. Zanotto F, Grassi V, Zucchi F, Merlin M, Balbo A. Effect of sensitization on stress corrosion cracking behaviour of LDX 2101 stainless steel. La Metallurgia Italiana, 2016;12-108:23-33. Zanotto F, Grassi V, Balbo A, Monticelli C, Melandri C, Zucchi F. Effect of brief thermal aging on stress corrosion cracking susceptibility of LDSS 2101 in the presence of chloride and thiosulphate ions. Corros Sci, 2018; 130:22–30. Zanotto F, Grassi V, Balbo A, Monticelli C, Zucchi F. Influence of thermal aging on SCC susceptibility of DSS 2304 in the presence of chlorides and thiosulphates. La Metallurgia Italiana, 2018; 3-110:40-44. Zanotto F, Grassi V, Balbo A, Monticelli C, Zucchi F. Stress-Corrosion Cracking Behaviour of Lean-Duplex Stainless Steels in Chloride/Thiosulphate Environments. Metals, 2018; 237(8):1-16. Zanotto F, Grassi V, Balbo A, Monticelli C, Zucchi F. Resistance of Thermally Aged DSS 2304 against Localized Corrosion Attack. Metals, 2018; 1022(8):1-18. Lo KH, Shek CH, Lai JKL. Recent developments in stainless steels. Mater Sci Eng Reports, 2009; 65:39–104. Garfias-Mesias LF, Sykes JM, Tuck CDS. The effect of phase compositions on the pitting corrosion of 25 Cr duplex stainless steel in chloride solutions. Corros Sci, 1996; 38:1319–1330. He L, Guo Y-J, Wu X-Y, Jiang Y-M, Li J. Effect of Solution Annealing Temperature on Pitting Behavior of Duplex Stainless Steel 2204 in Chloride Solutions. J Iron Steel Res Int, 2016; 23:357–363. International Organization for Standardization. UNI EN ISO 17864:2005: Corrosion of Metals and Alloys—Determination of the Critical Pitting Temperature under Potientiostatic Control; International Organization for Standardization: Geneva, Switzerland, 2005. Ruel F, Saedlou S, Mendibide C, Manchet SL, Wolski K. Effect of the substitution of Ni by N and Mn in Lean Duplex Stainless steels on Stress corrosion Cracking assisted by H2S. In NACE–International Corrosion Conference Series; National Association of Corrosion Engineers: Houston, TX, USA, 2017, p. 3376–3387. An LC, Cao J, Wu LC, Mao HH, Yang YT. Effects of Mo and Mn on Pitting Behavior of Duplex Stainless Steel. J Iron Steel Res Int, 2016; 23:1333–1341. Chen L, Tan H, Wang Z, Li J, Jiang Y. Influence of cooling rate on microstructure evolution and pitting corrosion resistance in the simulated heat-affected zone of 2304 duplex stainless steels. Corros Sci, 2012; 58:168–174.
La Metallurgia Italiana - n. 3 2019
metallurgia
co
per non metallurgisti
rs
o
15-16-22-23-29-30 ottobre 2019, Milano (c/o Centro Congressi Fast)
L’Associazione Italiana di Metallurgia propone la settima edizione del Corso Metallurgia per non Metallurgisti. L’iniziativa di formazione si rivolge a chi lavora con i materiali metallici e sovente non ha potuto ricevere in tempo opportuno le basi metallurgiche necessarie. Il Corso si sviluppa in sei giornate, suddivise su tre settimane consecutive, per garantire continuità e, nel contempo, limitare assenze prolungate dalle aziende. I docenti, scelti tra esperti noti in Italia e all’estero, di estrazione sia accademica che industriale, assicurano un mix di competenze difficilmente raggiungibili in altra sede. Il Corso è dedicato a persone non coinvolte in attività che richiedono una preparazione avanzata nei vari settori metallurgici, ma alle quali è sufficiente la conoscenza, l’interpretazione e la spiegazione della metallurgia nei più svariati settori tecnologici, applicativi, selettivi, ispettivi e di collaudo. “Metallurgia per non metallurgisti”, con un linguaggio per “non addetti ai lavori”, è rivolto a manager, personale di vendita e di acquisto, progettisti, ispettori e in generale a tutti coloro che devono lavorare e utilizzare materiali metallici, senza avere necessariamente una preparazione universitaria. Il programma fornisce sia pure succintamente, e lezione per lezione, le principali indicazioni riguardanti i diversi argomenti trattati nel Corso. Le lezioni, si susseguiranno tra loro in modo da fornire al partecipante una conoscenza panoramica a trecentosessanta gradi dell’ampio settore metallurgico.
Il programma completo è disponibile su www.aimnet.it Evento patrocinato da
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co CMYK
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0 68 137
280 CV
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424 C
bianco 100%
metallurgia fuori forno
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28-29 marzo, 4-5-10-11 aprile 2019 MILANO, CREMONA, NARNI SCALO, TERNI, MARONE, SAN ZENO NAVIGLIO
L’iniziativa si sviluppa come di consueto alternando alla didattica in aula le visite agli impianti e alle aziende del comparto che si sono rese disponibili ad ospitare il Corso. I partecipanti possono così fruire di informazioni di carattere teorico e pratico, che consentono di migliorare la qualità dei prodotti senza compromettere l’efficienza produttiva e gli impianti. In un’ottica di consolidamento delle realtà siderurgiche, risultano determinanti fattori quali il contenimento dei costi produttivi, la produttività, la qualità dei prodotti e la rispondenza puntuale alle specifiche dei clienti. La conoscenza e la corretta applicazione dei processi di “metallurgia secondaria” riveste un ruolo fondamentale per il raggiungimento di questi obiettivi da parte delle imprese siderurgiche. In questo Corso, saranno approfonditi temi fondamentali per la buona riuscita del processo di “metallurgia secondaria”. Il programma completo è disponibile su www.aimnet.it Organizzato con collaborazione di
an
te
Le manifestazioni AIM AIM meetings and events METALLURGIA FUORI FORNO Corso itinerante 28-29 marzo, 4-5-10-11 aprile
LEGHE DI ALLUMINIO Corso di base Bologna, 27 giugno
METALLOGRAFIA Corso modulare Milano, 9-10-22-23 maggio, 12-13 giugno Vicenza, 17-18-19 settembre
XIII GIORNATE NAZIONALI SULLA CORROSIONE E PROTEZIONE Convegno Palermo, 3-4-5 luglio
METALLI A SCALA MICRO NANOMETRICA: TECNICHE DI INDAGINE Giornata di Studio Roma, 17 maggio
DEFORMAZIONE DEI PRESSOCOLATI: CAUSE E RIMEDI Giornata di Studio Torino, 18 settembre
BULLONERIA E TRATTAMENTI TERMICI Giornata di Studio c/o Vimifasterners - Reggiolo (RE), 21 maggio LEGHE DI NICHEL E SUPER LEGHE Giornata di Studio Milano, 28 maggio PROVE MECCANICHE Corso Milano, Monza, Crema 29-30 maggio, 5-6 giugno
“ADDITIVE METALLURGY”. MATERIALI METALLICI E FABBRICAZIONE ADDITIVA Corso Milano, 18-19 settembre MASTER PROGETTAZIONE STAMPI Corso modulare Brescia, 25-26 settembre, 9-10-23-24 ottobre, 6-7-20-21 novembre, 4-5 dicembre FORGIATORI Corso itinerante 2-3-9-10 ottobre
ECHT 2019 - HEAT TREATMENT & SURFACE ENGINEERING FOR AUTOMOTIVE Convegno Internazionale Bardolino (VR), 5-6-7 giugno
METALLURGIA PER NON METALLURGISTI Corso Milano, 15-16-22-23-29-30 ottobre
POLVERI E PROCESSI PER ALTE PRESTAZIONI Giornata di Studio Milano, 6 giugno
IGIENE DELLE LEGHE DI ALLUMINIO Corso 23-30 ottobre
DIFETTOSITA’ IN COLATA CONTINUA E LINGOTTI Giornata di Studio c/o TenarisDalmine - Dalmine (BG), 11 giugno
METALLURGIA SICURA Corso itinerante 30 ottobre, 6-13 novembre
CORROSIONE NEGLI IMPIANTI OIL&GAS Corso di base Milano, 12 giugno
PROVE NON DISTRUTTIVE Corso Milano, 20-21 novembre
RIVESTIMENTI - II MODULO Rivestimenti spessi - Placcatura e Termospruzzatura Corso Milano, 19-20 giugno
METALS FOR ROAD MOBILITY International Meeting Bergamo, 21-22 novembre
MICROSCOPIA ELETTRONICA IN SCANSIONE SEM PER METALLURGISTI II ed. Corso Lecco, 26-27-28 giugno
MATERIALI INNOVATIVI PER L’ADDITIVE MANUFACTURING Giornata di Studio 4 dicembre
Per ulteriori informazioni rivolgersi alla Segreteria AIM e-mail: info@aimnet.it oppure visitare il sito internet www.aimnet.it
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La Metallurgia Italiana - n. 3 2019
AttualitĂ industriale Technology to Control Inclusions in Stainless Steels edited by: S. Lee, E. Jeong, J.-W. Ki, K. Kim, J. Choi, S.-Y. Kim In order to improve surface quality in stainless steels, many methods have been developed. These methods not only remove nonmetallic inclusions but also modifies the inclusion composition. The following results were obtained from this study. (1) The origin of harmful inclusions is the AOD slag suspended above the molten steel, which was not removed and subsequently enriched in Al2O3 and MgO during casting. MgO and Al2O3 inclusions could be reduced by controlling the temperature of the molten steel and the slag basicity during the AOD (Argon Oxygen Decarburization) process. (2) The effects of bottom bubbling on the removal rate of inclusions during LT (Ladle Treatment) were also analyzed, which is generally known that inert gas soft stirring leads to greater removal of inclusions. In determining and controlling the beneficial level of bottom bubbling with an inert gas, a control system to accurately measure argon stirring energy irrelevant to the stirring condition was developed and applied to the production of austenitic stainless steels. As a result, in order to increase the removal rate of inclusions, it is essential to supply sufficient stirring power to the molten steel during the soft bubbling stage. (3) It has also been found that the impact pad during pouring in the tundish, which is designed to confine the turbulence and slow down the flow of molten steel, can enhance the opportunity for inclusions to float and attain effective removal of ladle fillers and slag entrapped in the steel melt during the mixing period between heats.
KEYWORDS: INCLUSION REMOVAL - INCLUSION MODIFICATION - LADLE TREATMENT - STIRRING - TUNDISH - IMPACT PAD
Sangbeom Lee, Eunju Jeong, Seong-Yeon Kim POSCO Technical Research Laboratories
Jun-Wan Ki, Kisu Kim, Jayong Choi Department of Stainless Steelmaking, POSCO 6261, Donghaean-ro, Nam-gu, Pohang-si, Gyeongsangbuk-do, Republic of Korea
INTRODUCTION During steelmaking, inclusions are generated by various causes. Typical inclusions are a result of deoxidation, which are called indigenous inclusions. In addition, inclusions can be generated by slag entrapment, reoxidation, impurities in ferro-alloys and refractories, which are termed exogenous inclusions.
La Metallurgia Italiana - n. 3 2019
Inclusions can cause clogging of the submerged entry nozzle during casting and generate many defects in the stainless steel product, as shown in Figure 1. For example, AISI 304 stainless steels can have fine linear defects called slivers observed after cold rolling when inclusion control is insufficient. Sliver defects occur as lines along the steel strip surface parallel to the rolling direction.
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Industry news
Fig. 1 – Nozzle clogging by Inclusions (Left) and sliver defects on the surface of as-rolled products (Right) Figure 2 shows the stainless steel manufacturing process at the Pohang works of POSCO. The raw material and scrap are melted in an electric arc furnace and the molten metal from the furnace is decarburized, deoxidized, and desulfurized in the AOD (Argon Oxygen Decarburization) furnace. After the AOD, the steel composition is finally adjusted and the temperature is controlled to the level suitable for casting and inclusions removed by bubbling into the ladle using a porous plug during the LT (Ladle Treatment) process, which is then cast by a continuous caster. Generally, in stainless steels, inclusions in steels are formed by slag infiltration and mixing during tapping after the AOD. Initially, CaO-SiO2 is produced by the combination of the deoxidation product (SiO2) and CaO. During this initial formation, the Al present in the de-oxidizer is mixed to form CaO-SiO2-Al2O3. This forms a liquid inclusion at steelmaking temperature that are usually spherical in shape. The inclusion characteristics are changed depending on the slag basicity and the initial composition of the inclusions is also changed by the reoxidation and subsequent reactions with the slag/ metal/refractory [1]. Exogenous inclusions can form from other sources including
impurities from raw materials and fluxes, broken refractory brickworks, ladle fillers, and ceramic lining particles. They may act as sites for heterogeneous nucleation of alumina surrounding the central particle. The inclusions can be removed by floating in the LT and the tundish. The inclusions in molten steel are floated during the LT process by bubbling and can be separated naturally by the upward flow of molten steel developed within the tundish. On the other hand, efforts should be made to prevent the entrainment of impurities. The inclusions originating from ferro-alloys, additives, and refractories should be suppressed by strict material management. In addition, the entrapment of tundish fluxes and ladle fillers should also be suppressed. Steel cleanliness can be controlled by a wide range of operating practices in the steelmaking process. These include the basicity of AOD slags, tapping temperature, the time and extent of stirring during the LT, tundish practices, and the absorption capacity of the various metallurgical fluxes [2]. In the present work, the methods for evaluating steel cleanliness are provided and the results of commercial operating practices to improve steel cleanliness at the LT and tundish are described.
Fig. 2 – Stainless steel manufacturing process at the Pohang works of POSCO and sources of Inclusions during typical operating conditions 46
La Metallurgia Italiana - n. 3 2019
Attualità industriale METHODS TO EVALUATE INCLUSIONS The steel cleanliness is evaluated by observing the surface of the as-cast slab and measuring the number of inclusions per unit area of a slab sample. For rapid analysis, an area of 50 mm x 50 mm is observed with an optical microscope at 50 times magnification. The sampling position is generally the surface of the slab. After cutting the desired area, the surface is milled to 1 mm depth and polished followed by inclusion analysis for the size and numbers using an automated analyzer, where the number of inclusions per unit area is given. For precise analysis, the Auto-SEM inclusion analysis method is used to observe the area of the slab sample similar to the
optical microscope. Only the inclusions having a size of 10 μm or more are analyzed and counted in the number of inclusions per unit area. The shape of the inclusions detected in AISI 304 stainless steel is shown in Figure 3. In the CaO-SiO2-Al2O3MgO quaternary system, MgO-Al2O3 spinel inclusions can form at high MgO and Al2O3 content. Al2O3, MgO, CaO-Al2O3 or SiO2 inclusions are classified as exogenous. The most important and direct indicator of cleanliness is the number of slivers detected from the coil after cold rolling. It is expressed as the number of slivers per unit area. It is regarded as a very important quality indicator since the coil cannot be sold as a regular product if the quality code is detected above the critical cleanliness threshold.
Fig. 3 – Shape and composition of Inclusions observed in AISI 304 stainless steel IMPROVEMENT IN THE AOD OPERATION When the number of inclusions in the slab increases, the number of slivers also increases linearly, as shown in Figure 4. In addition, depending on the composition of inclusions such
as higher MgO content, a steeper slope can be observed with the number of inclusions increasing by 1.7 times compared to inclusions containing lower MgO.
Fig. 4 – Relationship between the cleanliness in slab and the cleanliness in cold-rolled coil As a result of checking the composition of inclusions from good and bad quality samples at arrivals/departures of the LT process, the basicity and Al2O3 contents were similar, but the MgO content in the bad quality case was appreciably higher compared to the good quality case. This suggested a change La Metallurgia Italiana - n. 3 2019
in the AOD operations to reduce the MgO content in inclusions [2, 3]. It was expected that the MgO concentration in inclusions could be lowered, when the amount of AOD slag (Mg Source) is reduced by reducing the amount of oxygen blowing. In addition, it was also observed that shorter stir47
Industry news ring time after addition of the de-oxidizer and lower average residence times of the melt in the ladle results in lower MgO content. Furthermore, a reduction of the maximum temperature in the AOD operation decreased the amount of MgO dissolution into the melt. A higher molten steel temperature results in greater amounts of MgO elution from the refractory walls. This tendency appears even with high basicity.
LADLE OPERATIONS FOR CLEAN STEEL Ladle slags The content of Al2O3 in inclusions was higher than 25% in a low-quality product. Figure 5 shows the composition of inclusions containing low and high Al2O3, respectively.
Fig. 5 – Composition of inclusions containing low and high Al2O3 on the quaternary phase diagram In order to decrease the Al2O3 content in inclusions, the LT top slag test was conducted. The LT top slag test attempts to lower the inclusion basicity by mixing AOD slag and rice hulls instead of using only the AOD slag. The basicity of the modified top slag was between 1.0 and 1.2. As a result of the test, the slab cleanliness applying the modified top slag was improved by 10% compared to slabs utilizing only the AOD slag. As a result of the inclusion compositional analysis using AutoSEM, the inclusion basicity was 0.88 for the AOD slag and rice hull test, which was slightly different from 0.93 of the
conventional AOD slag only case. The content of Al2O3 in the inclusions was lower than that of the conventional ones and the content of MgO was lower than that of the conventional ones, as shown in Figure 6. It was confirmed that spinel inclusions are prevented from being formed, when the top slag is applied and most of the compositional range of the inclusions corresponded to low melting point phases. When the top slag was used, the effect of Al2O3 reduction in inclusions was observed, but the lowering of the inclusion basicity was found to be less than expected.
Fig. 6 – Composition of inclusions when the modified top slag is applied
48
La Metallurgia Italiana - n. 3 2019
Attualità industriale Stirring in the ladle The graph below shows the rate of floatation of the inclusions
during steady state without bottom bubbling cal-culated using the creeping flow approximation of Stoke’s Law. [1]
where u is the velocity (m/s), g is the gravitational acceleration (m/s2), r is the radius of the spherical particle, ρp is the mass density of the particles (kg/m3), ρf is the mass density of the fluid (kg/m3) and μ is the dynamic viscosity (kg/m*s). It is well-known that as the size of inclusions increases the time of floatation decreases drastically. Inclusions of 10 µm travel
30 cm in 30 minutes, but inclusions of 20 µm can travel about 1.3 m for the same time. An inclusion of less than 20 µm in the molten steel at a depth of 1.8 m requires a lifting time of 40 minutes or more, as shown in Figure 7. Therefore, it can be understood that sufficient bottom bubbling is necessary for inclusion flotation and separation from the molten steel.
Fig. 7 – Floating time from the depth of 1.8m as a function of inclusion diameter In addition, it can be deduced that a difference in the number of inclusions is present according to the depth of inclusion suspended in the molten steel [4]. To confirm this expectation, the number of inclusions with respect to the depth of
steel melt was analyzed. As shown in Figure 8, samples were taken at 0.24 m and 0.84 m from the top line of an AISI 304 stainless steel melt.
Fig. 8 – Schematic of the sampling depth in the ladle
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49
Industry news According to the test samples taken at different depths, there were insignificant differences in the composition of the molten steel as a function of depth or time. When molten steel agitation is not provided, which includes before and after the LT operation, the temperature at the upper region of the steel is about 8-9 degrees lower than the lower region. After mixing is initiated with bottom bubbling, the temperature difference decrease to less than 1 degree. According to Figure 9, the number of inclusions is larger than at the upper (240 mm) inclusions at the time of LT arrival and the number of inclusions decreases with bubbling. During the stationary state (No bubbling) after 30 minutes of bottom bubbling, the inner inclusions decreased, but the upper inclusions increased, resulting in an inversion of the inclusions
in the upper and the inside. The inclusion basicity decreases during ladle treatment, which is presumed to be due to Si reoxidation. The compositions of the upper and inner inclusions were comparable. As the LT operation progress, an increase in the number of re-entrained inclusions and the inclusion composition deviation was observed. Therefore, it can be seen that the bubbling time to float inclusions can be optimized for increased cleanliness of the steel. An increase in the reoxidation seems to have no adverse effects, which lowers the inclusion basicity. It is not desirable to increase the AOD temperature to increase the bubbling time since it increases the Al2O3 and MgO content in the inclusions.
Fig. 9 – Cleanliness index of melt sample taken at various depths in the ladle The stirring of molten steel is controlled by the amount of injected Ar gas. However, it is difficult to accurately control the flow rate, when problems such as clogging of the porous plug located at the bottom of the ladle or leaking of the Ar gas piping occur. In order to determine the proper flow rate, the plume eye on the surface of the steel melt was monitored by a camera. However, since the judgment of the open area relies on visual inspections by the operator, there can be some variation in the actual effective stirring force. To solve this issue, it was necessary to develop a continuous monitoring system with a flow control system. Methods for continuously measuring the effective agitation force include utilizing an accelerometer sensor combined with a thermal imaging camera, as shown in Figure 10. 50
The accelerometer sensor measures the agitation in the melt by detecting and quantifying the vibrations of the melt. When the molten steel is stirred by the Ar gas, the resulting vibration of the melt is assumed to be proportional to the stirring power of the molten steel [5]. The modified thermal imaging camera can clearly distinguish the molten steel and the slag compared to a con-ventional imaging camera to quantify the open area of the hot melt and to control the flow rate. By changing the Ar gas flowrate at real time using a combination of the accelerometer and thermal imaging camera, it was possible to dynamically adjust the flow rate and suppress the generation of an open area on the surface due to an excessive stirring intensity. La Metallurgia Italiana - n. 3 2019
Attualità industriale
Fig. 10 – Schematic diagram of Ar bubbling control system MODIFICATION OF TUNDISH OPERATIONS Beyond the LT operation and inclusion floatation through actively controlled bubbling, the slag cleanliness showed the number of inclusions to increase with increasing cast sequence in the continuous caster. It is speculated that the cause of greater sliver defects is attributed to the tundish since the cleanliness of each steel heat taken at the ladle after gentle bubbling is similar with every heat, but the cleanliness of the slab is worse with consecutive
cast heats. Using the tundish flux composition of each heat analyzed by XRF, the melting point and the viscosity were calculated as shown in Figure 11. It can be seen that the flux composition deteriorates toward the end of the cast sequence. Therefore, it is presumed that not only does the inclusions in the molten steel continuously absorb into the tundish flux, but portions of the ladle slag carryover and the filler sand also influence the change in tundish flux composition [6, 7]
Fig. 11 – Change of melting point and viscosity of tundish flux with casting sequence
In other word, the entrapment of inclusions can occur during the ladle exchange operation from the contaminated tundish flux, as depicted in Figure 12. Figure 13 shows the flux near the pouring zone of the tundish. The amount of flux increases with the cast sequence. In order to ascertain the direct cause,
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SrO tracers were added to the slag in the ladle or the flux near the pouring zone [1, 2, 8]. The SrO content was observed in the inclusions of the slab in each case, from which it was confirmed that both the ladle slag and tundish flux were entrained.
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Fig. 12 – Sources of inclusions; 1 ladle slag carryover, 2 filler sands and 3 tundish flux near pouring zone During ladle tapping, infiltration of the filler sand and the tundish flux into the molten metal occurs. It is well known that the ladle filler can be a source of Al2O3 and MgO in inclusions because the ladle filler is composed of Cr2O3, Fe2O3,
SiO2, Al2O3 and MgO. Reoxidation of molten steel (MnO, SiO2 inclusions observed) and refractory loss (Al2O3 inclusions observed) also occur.
Fig. 13 – Slag near pouring zone in the tundish after consecutive casting sequences From various past publications [9, 10, 11], the efficacy of installing an impact pad in the tundish has been suggested to be an excellent method to improve the floatation of inclusions. According to Fan et al., a weir-dam in the tundish increased the inclusion removal rate by about 30% compared to a standard tundish without a weir-dam configuration. When an impact pad was applied, the inclusion removal rate was increased by more than 10% compared to the weir-dam tundish configuration. In the case of a tundish with an impact pad, the residence time of the melt in the tundish was increased by 10% compared to the weir-dam tundish. These results were similar to the simulations carried out in this work. It was 52
also confirmed that the case of the impact pad without the weir was more effective for the removal of inclusions than the combination of the impact pad and weir from the analysis of the flow of molten steel in the tundish. As can be seen from the results of the flow analysis in Figure 14, the rising flow of the weir-dam tundish is weak and the tundish flux is mixed with the high-speed flow of the hot melt. Further, when the flow of inclusion floatation is weak, the inclusions are moved to the outlet directly with a high velocity through the discharge hole at the bottom of the tundish. In the case of the weir-impact pad tundish, the upward flow was blocked by the weir in the bottom of the dam and flux La Metallurgia Italiana - n. 3 2019
Attualità industriale contamination is possible due to the return flow by the weir. In the case of no weir-impact pad tundish, strong upward flow is generated by the impact pad and maintained. After passing through the ejection hole of the impact pad, the inclusions were separated by floatation. In the case of the impact pad, the impact pad was installed at the bottom of the pouring zone of the tundish as shown
in Figure 15. A weir was not installed. In order to maximize and clearly ascertain the improvement effect, 5 tons of residual molten steel was left in the ladle to prevent exogenous inclusion sources of ladle slags. In addition, Ar and LNG gas sealing of the tundish was applied continuously during the casting to prevent reoxidation of the melt.
Fig. 14 – Computational flow analysis for weir-dam, impact pad + weir and impact pad without a weir, respectively The results of the cleanliness analysis showed a 30% improvement on average and the inclusion basicity decreased to 0.83. The phenomenon of a gradual increase in the number of inclusions with successive cast sequences was reduced, as shown in Figure 16. Figure 17 shows most inclusions in the slab are located in the low melting point region (<1400 °C).
The cleanliness of slab was greatly improved by about 50%. This was quantified to be the direct result of improving the flow of molten steel by installing weir-free impact pads and reducing exogenous inclusions from the tundish flux and preventing reoxidation by Ar and LNG sealing in the tundish.
Fig. 15 – Concept of weir-free impact pad and Ar+LNG gas sealing test and installation of impact pad in the tundish
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Fig. 16 â&#x20AC;&#x201C; Cleanliness index change with casting sequence after installation of the impact pad
Fig. 17 â&#x20AC;&#x201C; Composition of inclusions after implementation of the impact pad at the tundish
SUMMARY Many methods have been developed in order to improve the surface quality of stainless steels. These methods not only remove non-metallic inclusions, but also modify the inclusion composition. The following results were obtained from this study. The origin of the harmful inclusions is the AOD slag suspended in the molten steel, which was not removed and enriched in Al2O3 and MgO during the casting process. MgO and Al2O3 content in the inclusions were reduced by controlling the temperature of the molten steel and the slag basicity in the AOD process. It is generally known that a soft stirring with an inert
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gas leads to an efficient removal of inclusions. In determining and controlling the level of bottom bubbling with an inert gas, a control system to accurately measure argon stirring energy regardless of the stirring condition was developed and applied. It has also been found that the impact pad in the pouring side of the tundish, which is designed to confine the turbulence in the pouring zone and slow down the flow of molten steel, can enhance the opportunity for the inclusions to float and ensure effective removal of ladle filler and ladle slags entrapped in the steel melt during mixing between heats.
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AttualitĂ industriale REFERENCES
[1]
Kim JW, Kim SK, Kim DS, Lee YD, Yang PK. Formation mechanism of Ca-Si-Al-Mg-Ti-O inclusions in type 304 stainless steel. ISIJ International. 1996;36(Supplement):S140-3 [2] Hojo M, Nakao R, Umezaki T, Kawai H, Tanaka S, Fukumoto S. Oxide inclusion control in ladle and tundish for producing clean stainless steel. ISIJ International. 1996;36(Supplement):S128-31 [3] Steneholm K, Andersson NAI, Tilliander A, Jonsson PG. Ther role of process control on the steel cleanliness. Ironmaking & Steelmaking. 2018;45(2):114-24 [4] Pitts-Baggett A, Turner P, Hallmark J, William B, Nastac L. Method for simultaneous sampling of steel at multiple depths in the refining ladle. Iron & Steel Technology. 2017 Jan:98-106 [5] Lee S, Byun SM, Park J, Suk MO, Jones JA, Walker DI. Effect of bottom bubbling on decarburization rate in VOD process. at the SCENMET III Conference, Lulea, Sweden, 8-11 June 2008;Vol.1:131-40 [6] Bessho N, Yamasaki H, Fujii T, Nozaki T, Hiwasa S. Removal of inclusions from molten steel in continuous casting tundish. ISIJ International. 1992;32(1):157-63 [7] Tanaka H, Nishihara R, Kitagawa I, Tsujino R. Quantitative analysis of contamination of molten steel in tundish. ISIJ International. 1993;33(12):1238-43 [8] Fuhr F, Torga G, Medina F, Cicutti C. Application of slag tracer to investigate source of non-metallic inclusions. Ironmaking & Steelmaking. 2007;34(6):463-70 [9] Zhang L, Thomas BG. State of the art in evaluation and control of steel cleanliness. ISIJ International. 2003;43(3):271-91 [10] Fan CM, Shie RJ, Hwang WS. Studies by mathematical and physical modelling of fluid flow and inclusion removal phenomena in slab tundish for casting stainless steel using various flow control device designs. Ironmaking & Steelmaking. 2003;30(5):341-7 [11] Badr K, Tomas M, Kirschen M, McIlveney G. Refractory solutions to improve steel cleanliness. RHI Bulletin. 2011;1:43-50
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Industry news Control of Large-sized Inclusions and Macro Segregation in Steel for High-Speed Train Wheels and Axels edited by: X. Wang, M. Jiang, K. Wang, W. Sun Investigation was made on control of large-sized non-metallic inclusions and macro segregation in high speed train wheel and axle steels produced by EAF-LF-RH-CC steelmaking route. It was found that, in the continuously cast round blooms containing less than 0.0006 mass% T.O, the inclusions larger than 100μm were all cluster typed and, from edge to center of the bloom, negative and positive segregation alternately occurred like the letter “W”. In LF refining, inclusions changed from Al2O3 firstly to MgO-Al2O3, then to CaO-MgO-Al2O3 and finally to CaO-Al2O3 system. As solid inclusions in liquid steel are easier to aggregate to larger ones, measures were taken to (1) delay the change of the solid inclusions to liquid ones of CaO-Al2O3 system by shorten the LF refining time and, (2) eliminate most solid inclusions in RH degassing by extending the RH refining time. The remnant inclusions after RH were mostly small-sized liquid inclusions which were difficult to aggregate to larger inclusions in subsequent continuous casting. The “W” typed macro segregation was proved mainly due to actions of M-EMS and S-EMS. By moving position of S-EMS far from the mold and adopting the EMS control pattern of “mild M-EMS / weak S-EMS / mild F-EMS”, problem of “W” typed macro segregation has been solved. Carbon content variation range of wheel and axle steel blooms of Φ380mm, Φ450mm and Φ600mm has been decreased to less than 0.03% and the segregation index of carbon (C/C0) has been lowered to less than 1.0±0.03. KEYWORDS: HIGH-SPEED TRAIN - WHEEL - AXLE - NON METALLIC INCLUSIONS - MACRO SEGREGATION
Xinhua Wang Shougang Research Institute of Technology, Shougang Group Co., Beijing, China; School of Metallurgical and Ecological Engineering, University of Science and Technology Beijing, Beijing, China
Min Jiang School of Metallurgical and Ecological Engineering, University of Science and Technology Beijing, Beijing, China
Kaizhong Wang, Wei Sun Technical Center, Ma Steel Co. Ltd., Maanshan, China
INTRODUCTION Wheels and axles are important parts of high-speed trains. Compared to common speed trains, the stresses of friction, impact, torque, etc. which the wheels and axles of high-speed train subjected to are much severer. As metallurgical defects in steels, such as the inner cracks, large-sized non-metallic inclusions, abnormal structures formed due to macro segregation, etc. can be origins of the fatigue fractures, their existence in steels for the wheels and axles of high-speed train must be very strictly controlled. 56
High-speed railway transportation has been developed very rapidly in the past 15 years in China. However, for a quite long time after the high-speed railway began to run in the country, the steels for high-speed train wheels and axles were mostly imported from abroad with high prices. In order to change this situation, industrial investigation was jointly carried out by Ma Steel Co. Ltd. and University of Science and Technology Beijing to develop the steelmaking technology for high-speed train wheel and axle steels. In this article, result of the investigation on control of large-sized non-metallic inLa Metallurgia Italiana - n. 3 2019
Attualità industriale clusions and macro segregation in continuously cast round blooms are described. EXPERIMENTAL Industrial experiments were carried out at Ma Steel Co. Ltd., in which steels for high-speed train wheels and axles were produced through the process of ‘‘Electric arc furnace (EAF) →Ladle furnace (LF) → RH degassing → Continuous casting’’. In EAF steelmaking, hot metals were charged together with scraps at a ratio around 55 mass%. During tapping (EBT), aluminum and part of refining fluxes were added into the ladle. In the subsequent LF refining, aluminum was added both into liquid steel and slag. [Al] content of liquid steel at end of LF refining was controlled between 0.035-0.05 mass% and, the slag basicity (mass% CaO)/ (mass% SiO2) and FetO content were controlled between 4.5-5.5 and ≤0.8 mass%, respectively. In RH degassing, the pressure was lowered to 67-90Pa and the vacuum refining time was between 2035min. After RH degassing, the molten steel was continuously cast to round blooms by a five-strand caster of curved type and with M-EMS, S-EMS and F-EMS devices. Chemical compositions the steel and slag samples taken in the experiments were made with conventional analytical methods. Non-metallic inclusions in steel samples were analyzed with a scanning electron microscope of ASPEX made of PEI Co., which can make automatic analysis of the inclusions by
combining SEM, EDS and a digital scan generator under computer control, obtaining size, morphology, location and chemical composition of each detected inclusion in the analyzed area [1]. In measurement of macro segregation, samples for chemical composition analysis were taken by drilling holes of 6mm in diameter on the cross-section specimen cut from the cast bloom with an interval of 10mm from loose side to fix side. Index of C/C0 is used to indicate magnitude of the macro segregation, where C is content of carbon of each individual sample and C0 is the average content of carbon of all samples. RUSULTS AND DISCUSSION Control of large sized non-metallic inclusions For decreasing the population and size of the non-metallic inclusions, total oxygen contents (T.O) of the steels for high speed train wheels and axles are controlled less than 0.0006 mass% in Ma Steel Co. However, large inclusions of several hundred micro meters can still be detected by ultrasonic test on hot rolled steel bars. Kawakami et al [2] found similar phenomenon that, although the inclusions in ultra-low T.O bearing steels were all smaller than 10μm by conventional inspection methods such as optical microscopy detection, inclusions as large as of several ten or hundred micro meters could still be found when ultrasonic test method was applied on bigger volume specimens or on fracture surfaces of specimens in fatigue test.
Fig. 1 – Illustration of the specimens taken from the continuous cast bloom
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Industry news For investigating the mechanism of the large-sized inclusions formed in ultra-low T.O steels, 10 large-sized specimens in the present study were taken from the continuously cast bloom of high-speed train axle steel containing 0.00059 mass% T.O, as shown in Fig.1. With the automatic detecting device of ASPEX, totally 33170 mm2 area of the polished surface of the specimens was scanned. For saving on analysis time, only the inclusions larger than 30μm were analyzed in the measurement.
Totally, 56 inclusions of oxides larger than 30μm were detected in the measurement. The detected inclusions were mainly of two categories: (1) singular spherical or near-spherical shaped inclusions of CaO-Al2O3 system and (2) cluster typed inclusions of CaO-MgO-Al2O3 system. As indicated in Fig.2 which shows sizes and locations of the detected inclusions, the inclusions larger than 100μm were all of cluster typed ones. While, the singular spherical or near spherical shaped inclusions were relatively smaller (all less than 40μm).
Fig. 2 – Sizes and locations of the inclusions detected in cast bloom specimens For investigating the mechanism of the large-sized inclusions formed in ultra-low T.O steels, 10 large-sized specimens in the present study were taken from the continuously cast bloom of high-speed train axle steel containing 0.00059 mass% T.O, as shown in Fig.1. With the automatic detecting device of ASPEX, totally 33170 mm2 area of the polished surface of the specimens was scanned. For saving on analysis time, only the inclusions larger than 30μm were analyzed in the measurement.
Totally, 56 inclusions of oxides larger than 30μm were detected in the measurement. The detected inclusions were mainly of two categories: (1) singular spherical or near-spherical shaped inclusions of CaO-Al2O3 system and (2) cluster typed inclusions of CaO-MgO-Al2O3 system. As indicated in Fig.2 which shows sizes and locations of the detected inclusions, the inclusions larger than 100μm were all of cluster typed ones. While, the singular spherical or near spherical shaped inclusions were relatively smaller (all less than 40μm).
Fig. 3 – Chemical composition distribution of small branches of the cluster typed inclusion (a), (b), (c) and (d), respectively 58
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Attualità industriale In Fig.3, the dark color denoted zone is the region calculated by FactSage, within which melting temperatures of the inclusions are lower than 1600˚C. It is seen that the small particles of the clusters were mostly solid inclusions whose melting temperatures are higher than 1600˚C. As solid inclusions usually have larger surface tensions than liquid ones to liquid steel [3], they are easier to aggregate to larger inclusions. Fig.4 shows SEM photographs of four typical singular spherical or near spherical shaped inclusions. It is seen that the
singular inclusions are either of single phase with light gray color like inclusion (a), or of two different phases with light gray or dark gray color, respectively, such as inclusion (b), (c) and (d). EDS analysis was made on different phases of the singular inclusions and the result is also given in Fig.3. It can be seen that the light gray phases of the inclusions are of liquid with melting temperatures less than 1600˚C, while the dark gray phases are of solid in molten steel.
Fig. 4 – Chemical composition distribution of light and dark gray phases of inclusion (a), (b), (c) and (d), respectively As can be seen in Fig.4, the dark gray colored parts of the singular inclusions are surrounded by light gray parts, indicating that outer surfaces of the inclusions are of liquid. Yin, et al [4] found that liquid inclusions of CaO-Al2O3 system were difficult to aggregate to larger ones owing to their lower surface tensions to liquid steel. It might be the reason why the singular spherical or near spherical shaped inclusions detected in the cast bloom specimens in present study were all smaller than 40μm. It was found in previous studies of the authors’ group [5-9] that, in secondary refining of the low oxygen special steels, the non-metallic inclusions changed firstly from Al2O3 inclusions which were products of Al deoxidation to MgO-Al2O3 system inclusions, then from MgO-Al2O3 inclusions to CaOMgO-Al2O3 and finally to liquid inclusions of CaO-Al2O3system. The RH degassing had much stronger ability to lower the T.O content and eliminate the inclusions, particularly the solid inclusions [4,9]. Based on the previous study results and the above introduced analysis result in present study that all the inclusions larger than 100μm detected in cast bloom specimens are clusters constituted by high melting temperature small Al2O3-MgO-CaO particles, the following two countermeasures are made for control of large-sized inclusions in hi-
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gh-speed train wheel and axle steels: (1) Shorten as much as possible the LF refining time to delay the change of the solid inclusions of MgO-Al2O3 and CaO-MgO-Al2O3 system to liquid inclusions of CaO-Al2O3 system. (2) Prolong the RH degassing time to more than 33 minutes to fully develop RH’s ability to eliminate as more as possible the solid inclusions. After RH degassing, the residual inclusions in molten steel could be mostly small-sized liquid ones which were difficult to aggregate to large inclusions in subsequent continuous casting. Fig.5 and Fig.6 show the Chemical composition distributions in CaO-MgO-Al2O3 system of inclusions at the time before LF refining, at 26min of LF, at end of LF refining and at end of RH degassing, respectively. It can be seen that, owing to decreasing LF refining time, the inclusions after LF refining were still mostly of solid ones of CaO-MgO-Al2O3 system, that are easier to be eliminated in RH degassing than the liquid inclusions of CaO-Al2O3 system. While, by extending RH degassing time to 33min, the inclusions after RH were almost all of small-sized (≤5μm) liquid ones which are difficult to aggregate to large inclusions in subsequent continuous casting.
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Fig. 5 – Chemical composition distribution of the inclusions at the time before LF, at 26min of LF and end of LF refining
Fig. 6 – Chemical composition distribution of the inclusions after 33min of RH refining Good result has been obtained in Ma Steel Co. after the above mentioned technique is applied. In ultrasonic test of the hot rolled bars, the reject ratio due to large inclusions has been lowered from 5.4% to 0.6%.
Fig. 7 – Variation of segregation index of carbon with distance to the center of the round bloom 60
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Attualità industriale Control of macro segregation in continuous cast round blooms Fig.7 shows the segregation indexes of carbon measured in the cast round bloom of Φ380mm of high-speed train wheel steel containing about 0.53 mass% carbon. For better describing feature of the segregation, the photograph showing the macro structure of the bloom specimen is also given in the figure. As Fig.7 shows, macro segregations formed mainly at four locations from edge to center of the bloom: (1) negative segregation at the location near surface, (2) positive segregation at the interface between columnar crystal zone and equiaxed crystal zone, (3) negative segregation at the location about 20-30mm to the center and (4) positive segregation at the center. This typed macro segregation can also be found in rectangular shaped continuously cast blooms of special steels. Since the C/C0 curve from edge to center like the letter “W” and the positive segregation peak at the interface between the
columnar crystal zone and equiaxed crystal zone often locates at 1/4 of the diameter of the round bloom or 1/4 of the thickness of the rectangular bloom, it is also called “W” typed segregation or “1/4 segregation”. Some researchers [10-13] also reported this typed segregation, but the mechanism is still not very clear. An experiment was made In the present study, in which power of electro-magnetic stirring (EMS) was adjusted. As can be seen in Fig.8, magnitude of the “W” typed macro segregation could be significantly decreased when the electricity current of S-EMS was lowered from 450A to 150A. It is also known that, in continuous casting of slabs, most casters are not equipped with S-EMS and F-EMS. There do also exist the negative segregation at the location 20-30mm to the center and the positive segregation at the center, but not exist the negative segregation at the location near surface and the positive segregation at the interface between the columnar crystal zone and equiaxed crystal zone.
Fig. 8 – Influence of S-EMS power on segregation index of carbon of Φ380mm round bloom of high-speed train wheel steel Base on the above analysis, It is considered that the electromagnetic stirring might be the main reason for the “W” typed segregation in round and rectangular bloom casting of special steels and the following mechanism of the segregation is proposed:
(1) In continuous casting mold, solidified shell of the bloom forms at the meniscus and then is withdrawn into the MEMS and S-EMS action zones (Zone A and Zone B in Fig.9). Owing to the flow of the liquid steel aroused by the EMS at the solidification front, contents of solutes in the solidified shell formed in the EMS action region decrease, i.e. negative segregation in the shell near the surface take place. In adLa Metallurgia Italiana - n. 3 2019
dition, as the magnetic force in the EMS action region from top to bottom increases firstly, reaches maximum at about middle height of the region and then gradually decreases till zero, the negative segregation aroused by the EMS behaves in quite similar way, i.e. increases firstly, reaches maximum at about middle height of the region, then decreases and finally disappeared. (2) When the bloom moves down into Zone C (shown in Fig.9) where no EMS acts, contents of solutes in molten steel at the solidification front increase gradually with the bloom going down. Correspondingly, contents of solutes in the solid phase formed also increase gradually. In this region, the segregation 61
Industry news turns from negative to positive segregation and, with the solidification process going on, the magnitude of the positive segregation keeps increasing. (3) When the columnar crystal zone advances inwards to a certain extent with the solidification process going on, equiaxed crystals are formed in molten steel some distances away from the solidification front owing to the so called “chemical composition supercooling” or “crystal rain” down from the molten steel. The newly formed equiaxed grains are purer
than the already formed columnar crystals. This is the reason why a positive segregation peak appears at the boundary between the columnar crystal zone and equiaxed crystal zone and there exists a negative segregation zone near the center of the bloom. (4) At final solidification stage, a slit is formed at center of the round bloom owing to the shrinkage of the solidified shell. Molten steel which contains more impurities flows into the slit, forming the centerline positive segregation.
Fig. 9 – Illustration of M-EMS affected zone, S-EMS affected zone, columnar crystal zone and equiaxed crystal zone, respectively Experiment was made to verify the mechanism, in which SEMS devices were moved far away from the mold to the position of 9-12m from the meniscus and various control patterns of M-EMS, S-EMS and F-EMS were tested. As can be seen in Table 1, by changing the EMS pattern from previously used
“strong M-EMS / strong S-EMS / strong F-EMS” to “mild MEMS / week S-EMS / mild F-EMS”, the magnitude of macro segregation as well as the inspection ratings of porosity, central cavity and central crack were remarkably lowered.
Tab. 1 – Result of the experiment of optimizing EMS parameters Power of EMS M-EMS
S-EMS
F-EMS
strong
strong
strong
1.0±0.08
1.0
0
0
mild
weak
mild
1.0±0.03
1.0
0
0
mild
weak
weak
1.0±0.04
2.0
2.0
0.5
Fig.10 and Fig.11 show the C/C0 curve and photographs of cross-section specimens of the cast round bloom of Φ600mm for high-speed train axles, respectively. By applying the developed technology in Ma Steel Co., problem of the “W” typed macro segregation has been solved. Now, the carbon content
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Defect rating
Segregation index of carbon (C/C0)
Porosity
Central cavity
Central crack
variation range of wheel and axle steel blooms of Φ380mm, Φ450mm and Φ600mm has been decreased to less than 0.03%, the segregation degree of carbon (C/C0) lowered to less than 1.0±0.03 and the central porosity ratings were mostly less than 1.0.
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Attualità industriale
Fig. 10 – Variation of segregation index of carbon with distance to the center of Φ600mm round bloom of high-speed train axle steel
Fig. 11 – Photographs of horizontal and vertical cross-section specimens of round blooms of high-speed train axle steel CONCLUSIONS Investigation was made on control of large-sized inclusions and macro segregation in high speed train wheel and axle steels produced by EAF-LF-RH-CC steelmaking route. It was found that, in the continuously cast round blooms which contain less than 0.0006 mass% T.O, the inclusions larger than 100μm were all cluster typed CaO-MgO-Al2O3 system inclusions and, from edge to center of the cast round bloom, negative and positive segregation alternately occurred like the letter “W”. In LF refining, inclusions changed from Al2O3 firstly to MgOAl2O3, then to CaO-MgO-Al2O3 and finally to CaO-Al2O3 system. As solid inclusions in liquid steel are easier to aggregate to larger ones, measures were taken to (1) delay the change of the solid inclusions to liquid ones of CaO-Al2O3
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system by shorten the LF refining time and, (2) eliminate most solid inclusions in RH degassing by extending the RH refining time to more than 33min. The remnant inclusions after RH were mostly small-sized liquid inclusions which were difficult to aggregate to larger inclusions in the subsequent continuous casting. The “W” typed macro segregation was proved mainly due to actions of M-EMS and S-EMS. By moving position of SEMS far from the mold and adopting the EMS control pattern of “mild M-EMS / weak S-EMS / mild F-EMS”, the problem of “W” typed macro segregation has been solved. Carbon content variation range of wheel and axle steel blooms of Φ380mm, Φ450mm and Φ600mm has been decreased to less than 0.03% and the segregation degree of carbon (C/C0) has been lowered to less than 1.0±0.03.
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Industry news REFERENCES
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Wang X, Li X, et al. Control of Stringer Shaped Non-Metallic Inclusions of CaO-Al2O3 system in API X80 Linepipe Steel Plates. steel research international. 2014; 85: 155-163. Kawakami K, Clean Baring Steel. The 182-183th Nishiyama Technical Memorial Seminar. Tokyo-Osaka, Japan. 2004; 151-179 Yang G, Wang X, et al. Transient Inclusion Evolution during RH Degassing. steel research international. 2014; 85: 26-34 Yin H, Shibata H, et al. Characteristics of Agglomeration of Various Inclusion Particles of Molten Steel Surface. ISIJ International. 1997; 37: 946-955 Wang X, Wang W, et al. Composition Change of the Non-metallic Inclusions during the Secondary Refining of Low Oxygen Alloyed Structural Steels. Asia Steel 2009 International Conference. Busan, Korea. 2009 Jiang M, Wang X, et al. Control of Non-Metallic Inclusions by Slag-Metal Reactions for High Strength Alloying Steels. steel research international, 2010; 81: 759-765. Jiang M, Wang X, et al. Laboratory Study on Evolution Mechanisms of Non-Metallic Inclusions in High Strength Alloyed Steel Refined by High Basicity Slag. ISIJ international, 2010; 50: 95-104. Wang X, Jiang M, et al. Study on Formation of Non-metallic Inclusions with Lower Melting Temperatures in Extra Low Oxygen Special Steels. Science China-Technological Sciences. 2012; 55: 1863-1872. Li J, Jiang M, et al., Investigation on Nonmetallic Inclusions in Ultra-Low-Oxygen Special Steels. Metallurgical and Materials Transaction B, 2016; 47: 2386-2399 Choudhary S K and Ghosh A. Morphology and Macrosegregation in Continuously Cast Steel Billets. ISIJ International. 1994; 34:338-345 Wu X, Wang F, et al. Effect of M-EMS on Composition Segregation of Midiem Carbon Steel. Steelmaking. 2005; 21(5): 21-23 Xu W, Cao Z, et al. Macro-segregation in cast bloom of gear steel and its formation mechanism. Steelmaking. 2011; 27(5): 63-67 Jing C, Wang X, et al. Study on Formation Mechanism and Countermeasure for control of 1/2R Segregation in Round Bloom of High Carbon Steel. Foundry Technology. 2011; 32(7):989-993
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difettosita’ in colata continua e lingotti
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11 giugno 2019, Dalmine (c/o TenarisDalmine)
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La qualità della struttura metallurgica di billette, blumi, bramme e lingotti è un aspetto fondamentale per garantire la qualità del prodotto finale ed evitare spiacevoli, e costosi, inconvenienti durante le lavorazioni successive o cedimenti anomali in esercizio. Poiché, al giorno d’oggi, molti produttori d’acciaio riorientano la propria attività verso la produzione di nuovi acciai speciali, l’attenzione alla qualità e alla riduzione dei difetti deve essere ancora maggiore. Inoltre, nella produzione delle billette l’uso di apparecchiature elettromagnetiche deve essere fatto con grande attenzione per evitare l’insorgere di effetti negativi, che potrebbero rendere controproducente l’applicazione di tali tecnologie. Anche la produzione di lingotti di grosse dimensioni necessita di estrema cura nella prevenzione della formazione di difetti e l’utilizzo di strumenti come la simulazione fluido-dinamica può essere di grande aiuto. La Giornata di Studio intende affrontare i seguenti temi: · il ruolo degli elementi chimici inquinanti nella formazione dei difetti · il ruolo della lubrificazione nella qualità superficiale dei semilavorati colati in continuo · difettosità prodotte dalla gestione termica della macchina di colata continua · utilizzo degli agitatori elettromagnetici per il miglioramento della qualità interna · gestione dei lingotti per la minimizzazione della difettosità dei forgiati · tecniche di simulazione per limitare i difetti nel prodotto finito
Il programma completo è disponibile su www.aimnet.it
metallografia 1° Modulo Giornate propedeutiche: Milano, 9-10 maggio 2019 2° Modulo Materiali ferrosi e saldatura: Milano, 22-23 maggio 2019 3° Modulo TT superficiali, sinterizzati e leghe di rame: Milano, 12-13 giugno 2019
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4° Modulo Leghe leggere: Vicenza, 17-18-19 settembre 2019 Visita tecnica c/o CNR ICMATE: Milano, 4 ottobre 2019
La metallografia trova applicazione primaria nella metallurgia sia nella fase di produzione dei metalli e leghe e sia nei processi come trattamenti termici, processi galvanici, fonderia, etc… e soprattutto nello “studio anomalie” meglio espresso da “failure analysis”. Il Corso è strutturato in un primo modulo propedeutico e tre successivi moduli dedicati rispettivamente a materiali ferrosi e saldatura, TT superficiali con e senza materiale d’apporto, sinterizzati e leghe di rame e infine a settembre un modulo di tre giorni relativo alle leghe leggere (alluminio, magnesio, titanio). Quest’ultimo modulo organizzato con il Centro di Studio Metalli leggeri, prevede delle esercitazioni pratiche presso i laboratori dell’Università di Padova, sede di Vicenza. Le lezioni tenute da docenti universitari e da tecnici di provata esperienza presenteranno un “taglio” applicativo, con riferimenti alla failure analysis, relativa sia ai processi di fabbricazione sia gli impieghi e soprattutto facendo riferimenti alle norme UNI – EN – ISO – ASTM (grano austenitico, decarburazione, precipitati, inclusioni non metalliche, bandosità, etc.). Ampio spazio sarà dato alla collaborazione tra docenti e partecipanti, allo scopo di beneficiare dei momenti di discussione e confronto al termine delle singole lezioni. All’atto dell’iscrizione si consiglia di scegliere il modulo o i moduli di interesse, non trascurando quello propedeutico, volto a facilitare la comprensione dei moduli successivi. A chiusura del Corso, verrà organizzata la consueta e interessante visita al Laboratorio CNR ICMATE di Milano, alla quale potranno accedere coloro che abbiano espresso interesse a parteciparVi in fase di iscrizione. L’ attestato di partecipazione rilasciato dall’AIM a fine corso, per chi abbia partecipato ad almeno due moduli, rientra tra la documentazione da presentare per la domanda di certificazione come esperto di 2° Livello di Controlli Metallografici.
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Experts’ corner Intervista al Dr. Roberto Moreschi Roberto Moreschi si è laureato in Scienze Biologiche nei primi anni Novanta. Direttore della Divisione R&D del Gruppo Unicalce, si occupa di ricerca applicata al miglioramento dei processi produttivi e allo sviluppo di prodotti a base calce nonché di industrializzare e fornire soluzioni tecniche a favore degli utilizzatori da ormai più di 25 anni. Nel corso della propria carriera ha altresì rivestito specifici ruoli di responsabile di prodotto per svariate applicazioni della calce; all’interno del Gruppo Unicalce ricopre anche il ruolo di Direttore dell’Ufficio Qualità. È membro attivo dell’Innovation Task Force e collabora in vari Gruppi Tecnici dell’Associazione Europea dei Produttori di Calce (European Lime Association – EuLA). In questa intervista il Dr. Moreschi presenta il Gruppo Unicalce e ne illustra la mission da cui, in particolare, nell’ottica di porsi come partner tecnologico per lo sviluppo e la generazione di valore, nascono il concetto di prodotti di filiera e soluzioni applicative della calce specificatamente concepite per l’industria dell’acciaio. Ci può presentare il gruppo Unicalce? Il Gruppo Unicalce è il primo produttore di calce aerea in Italia con 11 stabilimenti dislocati sul territorio nazionale e una capacità produttiva di circa 1,6 milioni di tonnellate all’anno. Il Gruppo Unicalce fornisce prodotti e soluzioni in vari settori applicativi afferenti a diversi mercati quali, solo per citarne alcuni, l’industria siderurgica, l’industria del vetro, l’industria chimica, l’agricoltura, le infrastrutture e l’edilizia, le applicazioni tecnologiche di salvaguardia ambientale e tutela del territorio, ecc. In un tale contesto multidisciplinare, oltre a essere impegnato a mantenere e consolidare la business leadership garantendo il soddisfacimento delle richieste di una così grande diversità di segmenti di mercato, il Gruppo Unicalce, in conformità con la propria missione di migliorare l’efficienza dei processi produttivi, propri e della clientela cui si rivolge, è da sempre attento alle tematiche della sostenibilità che rispondano responsabilmente alle esigenze della società moderna.
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Come riuscite a coniugare la sostenibilità con la produzione della calce? La produzione della calce, che come detto è una materia prima o un intermedio indispensabile in svariate applicazioni e impieghi manifatturieri, è di per sé un processo che presenta intrinsecamente un certo impatto sull’ambiente, tanto in termini di consumo di risorse quanto in termini energetici. Grazie all'elevata qualità dei propri giacimenti minerari, a impianti di produzione moderni e a un’attenta pianificazione e progettualità, il Gruppo Unicalce rappresenta oggi una solida realtà orientata a soddisfare la crescente domanda di prodotti a base calce innovativi e di qualità senza perdere di vista la sostenibilità ambientale. Lo sfruttamento responsabile delle risorse naturali avviene attraverso l’accesso a giacimenti di calcare di elevata purezza e la massimizzazione delle frazioni idonee alla produzione di calce che consente di estendere la vita utile dei giacimenti minerari medesimi, così come l’adozione di forni industriali tecnologicamente avanzati che si traduce in più ridotti livelli di
emissioni di inquinanti atmosferici e gas a effetto serra. Ci racconti del connubio tra calce e acciaio. L’indispensabilità dei prodotti a base calce nel comparto siderurgico discende sia dalla grande efficacia e versatilità di utilizzazione, quanto nel soddisfare la necessità di competitività richiesta nella produzione dell’acciaio. Pur essendo tecnicamente possibile l’adozione di diversi altri composti alcalini, i prodotti dell’industria della calce rappresentano oggi il migliore compromesso costi-benefici. Oltre a ciò è da sottolineare l’ampia disponibilità geografica dei prodotti a base calce: basti pensare, per esempio, che in Italia gli stabilimenti di produzione della calce sono sorti, non a caso, nelle zone di prossimità dei maggiori poli siderurgici. L’utilizzo della calce nella produzione dell’acciaio ha radici molto antiche anche se la descrizione tecnica dettagliata dell’impiego della calce in siderurgia, contenuta nell’opera “De re metallica” dello scienziato tedesco Georg Bauer La Metallurgia Italiana - n. 3 2019
Scenari (1494-1555) meglio noto col nome latino di Georgius Agricola, si ha solamente a partire dal XVI secolo. Ancora oggi, come in passato, la sinergia tra industria metallurgica e produzione di calce è estremamente stretta. A livello mondiale, nell’anno 2017, la stima del consumo di calce per i processi siderurgici è stata dell’ordine di 140-160 milioni di tonnellate, che ha rappresentato circa il 40-45% della produzione totale di calce. La sinergia tra acciaio e calce è così importante che, in alcuni casi, una parte di questa viene prodotta proprio dall’industria siderurgica stessa che, specialmente in stabilimenti “a ciclo integrale”, può avere integrata nei siti produttivi di acciaio la produzione della calce. Nel 2017 nei 28 paesi membri dell’Unione Europea (EU-28) a fronte di una produzione di calce circa pari a 20 milioni di tonnellate il comparto siderurgico ne ha consumate 8 milioni di tonnellate: risulta chiaro quindi che l’utilizzo in siderurgia rappresenta il primo settore di applicazione per l’industria della calce. L’uso della calce è imprescindibile nella siderurgia moderna in quanto assolve a molteplici ruoli fondamentali nelle differenti fasi del processo metallurgico, tanto per quanto concerne la fusione (tipicamente nei processi di metallurgia primaria cosiddetti “in forno”) che l’affinazione (processi di metallurgia secondaria anche definiti “fuori forno”), ma principalmente concorre alla formazione di una scoria basica, caratterizzata da una chimica equilibrata e una corretta viscosità, in grado di neutralizzare elementi acidi e che faciliti la rimozione del fosforo e dello zolfo. La formazione di una buona scoria deve altresì assicurare una adeguata copertura dell’arco elettrico. Pertanto, sebbene il costo unitario della calce abbia un impatto relativamente modesto se calcolato sul costo per tonLa Metallurgia Italiana - n. 3 2019
nellata di acciaio prodotto, la qualità e la quantità di calce utilizzata hanno un effetto diretto non solo sulla qualità della scoria ma più in generale sui risultati metallurgici in termini di qualità e di produttività, nonché sulla durata dei refrattari e sul costo totale della produzione. Cosa rappresenta per Unicalce la ricerca in ambito siderurgico? Dal punto di vista della ricerca Unicalce è sempre stata un’azienda orientata all’innovazione. Ne sono esempio le sfide affrontate per mettere a punto sistemi innovativi di coltivazione dei propri giacimenti o l’adozione di tecnologie emergenti per i processi di calcinazione del calcare. Per esempio, Unicalce è stata la prima in Italia ad adottare, allo stadio prototipale, forni tecnologicamente sviluppati per produrre calce da calcare fine che hanno permesso un migliore sfruttamento dei giacimenti garantendo l’ottenimento di prodotti aventi un livello qualitativo più controllato: questa tecnologia è ora ben diffusa a livello mondiale. Nel corso dell’ultimo decennio, Unicalce ha inoltre sviluppato e brevettato tecnologie innovative per la depurazione degli effluenti gassosi derivanti da processi di combustione, oggi largamente applicate in Italia e sempre più anche all’estero. Guardando più indietro su quanto fatto per l’ambito siderurgico, negli anni ’90 siamo stati i primi a studiare e proporre prodotti, che possiamo definire “customizzati”, calibrati con specifiche quantità di magnesio, atti a migliorare le prestazioni di fusione permettendo di ottenere un corretto bilanciamento di massa e la piena saturazione della scoria medesima che sono altresì necessari alla salvaguardia dei rivestimenti refrattari. Abbiamo inoltre sviluppato prodotti a bassissimo contenuto di carbonio re-
siduo (Extra Low Carbon) utilizzati per garantire livelli molto elevati di desolforazione nei trattamenti di metallurgia secondaria. Negli ultimi 15 anni in Unicalce si è investito molto sulla ricerca, cambiando anche il modo di farla: si è infatti passati da una ricerca fatta principalmente internamente e attraverso partners istituzionali come Università e Istituti di Ricerca, ad una maggiore condivisione di progetti con gli utilizzatori dei nostri prodotti. La sinergia con gli utilizzatori ha dimostrato di essere una strada vincente per l’ottenimento di risultati importanti e concreti. Seguendo questa filosofia, negli ultimi anni sono nati importanti progetti di ricerca che coinvolgono come partners alcuni dei nostri principali clienti. Il confronto con gli utilizzatori, oltre ad aver generato una grossa accelerazione alla crescita di idee e soluzioni innovative (cross-fertilization) ci ha fatto comprendere anche la necessità di allargare le nostre competenze ad ambiti applicativi che prima erano unicamente deputati all’utilizzatore finale e ci ha permesso di aumentare la comprensione e la consapevolezza su nuovi aspetti collegati all’utilizzo dei prodotti a base calce. Il percorso di crescita intrapreso negli ultimi anni ci ha portato oggi ad avere un centro di ricerca altamente qualificato in parte costituito da risorse interne e in parte costituito da un laboratorio congiunto appositamente realizzato in partenariato con una primaria Università Italiana dove operano ricercatori universitari dedicati alle tematiche della calce. Come immagina l’evoluzione dei prodotti a base calce per i futuri utilizzi siderurgici? Faccio una premessa. Guardando “con un occhio critico” all’utilizzo della calce nei processi siderurgici notiamo la possi67
Experts’ corner bilità di apportare innovazione non solo nei nostri prodotti ma anche nei sistemi di dosaggio della calce: gli impianti siderurgici sono spesso legati a concetti tradizionali di adduzione dei cosiddetti agenti scorificanti, specialmente durante la fase fusoria dove il consueto caricamento in cesta di un prodotto in zolle o in tramoggia di un prodotto granulare pone dei limiti nell’efficientamento del processo. Riteniamo che il futuro, data l’oggettiva necessità di armonizzare gli aspetti ambientali e della sicurezza con quelli di una maggiore efficienza di processo, porterà allo sviluppo e alla diffusione dei sistemi pneumatici di iniezione di prodotti a base calce in forma polverulenta. In tale maniera, grazie all’utilizzo di prodotti a maggiore grado di finezza, caratterizzati da una più alta superficie specifica e reattività, si potrebbe gestire l’apporto del flussante con modularità, maggiore controllo e precisione nel rispetto dei tempi metallurgici, con conseguenti migliori rese in forno. Nei processi di metallurgia primaria ciò si traduce nella massima omogeneità di dosaggio del prodotto direttamente all’interfaccia scoria/acciaio e ottimizzazione delle caratteristiche di schiumeggiamento della scoria, migliore sfruttamento delle caratteristiche del prodotto con ridotti consumi specifici e incremento della velocità nel processo. L’additivazione controllata di prodotti a base calce è altrettanto importante nei processi di metallurgia secondaria al fine di assicurare una rapida formazione di scoria e massimizzare l’interazione scoria/acciaio ottimizzando la viscosità e la permeabilità per una migliore rimozione tanto dello zolfo e di eventuali inclusioni, quanto delle condizioni di degasaggio. Per quanto concerne l’evoluzione dei nostri prodotti, penso all’importanza di comprendere ancora meglio gli effetti combinati che la presenza di alcune
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impurezze, peraltro caratteristiche dei calcari di partenza, possono rivestire nel processo siderurgico: mi riferisco per esempio al contenuto di Fe2O3 e Al2O3 che potenzialmente possono avere effetti favorevoli sulla dissoluzione della calce nella scoria che, seppure generalmente disponibili solo in piccole quantità, potrebbero essere meglio bilanciati in alcuni particolari prodotti. Oltre alle caratteristiche chimiche stiamo lavorando per ottimizzare anche le proprietà fisiche quali la distribuzione granulometrica, la superficie specifica (BET) e la porosità (BJH) nonché la conformazione delle particelle (siano esse polveri o granuli) al fine di migliorare ulteriormente la reattività; a questo scopo siamo convinti che possa risultare assai interessante in alcuni casi l’utilizzo di specifici additivi nonché il ricorso a opportuni trattamenti superficiali delle particelle. Un ulteriore aspetto di grande rilevanza è legato alle fasi di movimentazione e dosaggio dei prodotti per i quali, soprattutto nel caso di utilizzo di prodotti a più elevato grado di finezza, sono sempre più richieste buone caratteristiche di scorrevolezza e fluidità che si traducono nello studio approfondito della reologia e nell’eventuale massimizzazione delle proprietà di flusso. In ultimo, non solo il miglioramento continuo delle caratteristiche dei prodotti ma anche lo studio della sinergia della calce con altri agenti flussanti potrebbe aprire nuove prospettive non ancora completamente studiate in termini di efficienza. Come possiamo chiudere il cerchio della sostenibilità? In un’ottica di economia circolare, negli ultimi anni, la nostra attività di ricerca si sta sempre più orientando al miglioramento della sostenibilità dei processi
produttivi sia interni, attraverso l’implementazione tecnologica improntata al miglioramento continuo delle produzioni e allo sviluppo di prodotti sempre più performanti, che degli utilizzatori dei nostri prodotti, collaborando con loro ad ottimizzare i consumi di calce e conseguentemente a ridurre gli scarti generati. In un contesto di repentino mutamento dei paradigmi operativi ed organizzativi in cui sarà sempre più necessario coniugare ambiente e sviluppo attraverso l’attuazione di un modello circolare, stiamo guardando con grande interesse alla possibile valorizzazione di alcuni residui o sottoprodotti generati dal processo metallurgico in quanto riteniamo che le nostre competenze possano coniugarsi bene con tali materiali che spesso hanno, per propria genesi, assoluta affinità coi geomateriali naturali nonché una stretta parentela con la calce. In conclusione, come vede il futuro? Sono convinto che nel contesto di una sempre maggiore globalizzazione, l’industria siderurgica italiana sarà chiamata nei prossimi anni a sfide sempre più importanti. Per potere rispondere alle domande di mercati sempre più complessi ed esigenti, dovrà prevedibilmente continuare in un processo di trasformazione improntato sulla competitività e sostenibilità di processo. È una sfida difficile anche per noi, in qualità di produttori di calce e partner tecnologici dell’industria siderurgica, in quanto dovremo affrontarla insieme con un atteggiamento flessibile e dinamico sempre più improntato sulla connettività e con un continuo spirito di innovazione.
La Metallurgia Italiana - n. 3 2019
Scenari Challenges for Secondary Metallurgy to meet future needs edited by: M. Dorndorf - C. Schrade Dr. Markus Dorndorf started his career in 1997 with studies for Mechanical Engineering at the Technical University Berg-Akademie Freiberg, Germany. After a time period of 2 years with experiences in R&D management at the Fraunhofer Institute, he joined the steelmaking industry and gained his PhD. Between 2008 and 2014, Markus Dorndorf headed the department of R&D for Electric Steelmaking and Secondary Metallurgy for Siemens VAI and was responsible for the Product Portfolio Management. With his experience in product design, the receipt of the Siemens Environmental Award in 2011, the Steel Manufacturers Association Award in 2012 (USA) and as owner of more than 12 patent families he was nominated as Product Manager for Melt Shops for the Tenova S.p.A. / Tenova Metals Deutschland GmbH in April 2015. He is focusing on global trend evaluation, product portfolio development and strategic customer projects. Christian Schrade graduated in Mechanical Engineering. In 1993, he joined Voest Alpine Industrieanlagenbau as Sales and Technology Manager where he stayed until 2000. Between 2001 and 2009, he became Head of Business Segment Metallurgy and Vice President at Siemens VAI Metals Technology. Then, in 2010, he moved to TECHNOMETAL GmbH where he was appointed Managing Director. In July 2013, he joined Tenova Metals Deutschland as Managing Director.
Mr. Dorndorf, Mr. Schrade, you have been in steelmaking business for several decades, with extensive experiences in electrical steelmaking and secondary metallurgy. In the latter, where do you see main challenges and demands for secondary metallurgy in the future? To start with some basic considerations: several aspects in secondary metallurgy will become a bottleneck and push its further development, even if in the last years we have seen only incremental improvements instead of breakthrough technologies or game changing innovations in this technological field. In general, rapidly changing conditions on raw material quality and availability, difLa Metallurgia Italiana - n. 3 2019
ferent energy sources on changing price levels, stricter environmental regulations and in general difficulty in predicting market conditions are topics iron- and steel makers have to deal with. These terms are representing threats, but also opportunities; those who act now will also play an important role in the future. In particular, two main facts have to be considered mid-term. First, the final asymptotic target world steel demand in future will be covered by 80 - 85% with recycled steel scrap and only 15 â&#x20AC;&#x201C; 20% has to be substituted by iron ore considering worldwide economic growth and the related product lifecycle in emerging markets. Additionally, scrap qualities worsening due
to higher complexity in structures and components (compound materials); this, in turn, leads to higher efforts in liquid steel treatment (Secondary Metallurgy) to reach the requested degree of purity and chemical composition. Second, increasing demands in higher steel grades and qualities specified by purpose of use and mechanical properties (strength, ductility, toughness and weldability) lead to higher requirements for metallurgical processes. The fact, that the purpose of the electric arc furnace (EAF) is reduced to a highly efficient melting unit promotes the importance of secondary metallurgy processes in the future.
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Experts’ corner Where do you see the limits of liquid steel treatment and secondary metallurgy in the electric steel making route considering the trends described before? Frankly speaking, it’s all about process efficiency and related demand in energy and time. In well performing melt shops the necessary process time, meaning from incoming material (steel scrap), up to the point when the final product (coils, blooms, etc.) leaves the factory is only about 2 - 3h. In the future, when the “steel on demand” approach (producer – customer) and its inherent request for highest flexibility will become more and more important all processes (scrap treatment – melting – liquid steel treatment – casting and rolling) will have a certain period of time to perform. Thus, considering worsening scrap qualities and in parallel higher requests for steel grades and qualities, can be handled only in a holistic approach optimizing all different process steps and the interfaces in between the process units. For instance, a dedicated scrap treatment in combination with a continuous scrap charging and melting unit leads to higher yield (upstream sepa-
ration of non-ferrous metals as well as copper and aluminum), increased scrap density and finally enhanced preheating and melting efficiency. This in line with flexible atmospheric (chemical heating station, ladle furnace, ladle treatment station) and vacuum based processes (Vacuum Oxygen Decarburization, Vacuum Stream Degassing, Vacuum Circulation Processes (RH), Vacuum Degasser) provides highest process flexibility and utmost efficient utilization of charging materials in conjunction with lowest operational expenditures. Nevertheless, we have already designed and we are improving further our cutting-edge solutions in the field of secondary metallurgy to meet future needs. Especially with latest measurement technologies in conjunction with our long-term experience we improve our process control algorithms continuously to minimize impurity fractions [C], [P], [N], [H2], [S] and related inclusion sizes for special steel grades. Besides this, in our opinion, the steel plant of the future will be more than a specialized steel producing unit; it will be a recycling facility dealing with various input materials and produce va-
rious products next to steel, to create additional business with by-products and so to remain competitive in a longterm. You have recently realized a RH plant for Salzgitter Flachstahl GmbH in Germany; it is very difficult to realize a brown-field project in an existing steelmaking area. What are the results? The metallurgical aim was to improve the quality of the overall production comprising high carbon, carbon, microalloyed and HIC-resistant grades for plates as well as IF and low carbon grades for automotive parts. Referring to this Salzgitter Flachstahl GmbH awarded Tenova Metals Deutschland GmbH with the new installation of Duplex RH plant in the integrated steelmaking site at Salzgitter in 2015. After two years of design and erection, the plant has been commissioned in April 2017 with an additional annual vacuum capacity of 1.5 Mill. tons to produce IF grades with C – contents with < 20 ppm (automotive). The scope of supply and services contained RH core equipment, a ladle tre-
Fig. 1 – Plant view south-west Salzgitter Flachstahl GmbH 70
La Metallurgia Italiana - n. 3 2019
Scenari atment station with wire feeding, a ladle lifting / vessel exchange facility, a top lance system, a combination of mechanical and steam ejector based vacuum system (hybrid), alloying system as well as fume treatment and water treatment plant. Furthermore, the scope comprises 4 ladle cars for quick and easy ladle handling, two snorkel deskulling cars as well as one snorkel maintenance car for quick snorkel cleaning and highest plant efficiency and availability. It has been a challenge considering the complex task in an existing brownfield area, the limited space available and the tight time schedule to realize this project. Especially the seamless integration of the new designed vacuum process under operating conditions has been an important fact for the realization of this project. Thanks to our well experienced team, working together for so many years, we were able to write a success story together with our customer. One of the main features is the use of steam ejector technology for vacuum generation, which youâ&#x20AC;&#x2122;ve called hybrid. Do you see advantages compared to solutions with
La Metallurgia Italiana - n. 3 2019
mechanical pumps only? We prefer a hybrid solution consisting of liquid ring vacuum pumps and steam ejector technology to realize the best compromise of the two worlds: lowest investment costs and energetically highest efficiency of steam ejectors for vacuum below 150 mbar and easy to maintain LRVPs with high reliability for vacuum above 150 mbar. Especially for deep vacuum below 20 mbar for lowest carbon and hydrogen contents a high vacuum capability at high suction loads is needed to perform in a certain time period. Nevertheless, the main driver is the availability of steam from the CAPEX and OPEX point of view. Again, the example of Salzgitter Flachstahl: The Hybrid vacuum system is designed in an economical configuration: The cooling water is used twice for condensation and by this measure the cooling water flow rate can be reduced by 30 % â&#x20AC;&#x201C; compared to conventional systems. How do you rate changes in availability and quality of charging materials and the in parallel increasing demand for higher steel qualities/
grades? What will be the effects and efforts for secondary metallurgy and its processes in the future? There is an increasing demand in high quality product and many steelmakers have taken up the idea of a 100% vacuum production. This gives the advantage of a higher cleanliness in the steel products and, by applying the appropriate process route, widens the range of products which can be offered to the market. In any case the request by the steel processing companies for clean steel with lowest residual impurities is rising. Final statements: Tenova provides sophisticated technologies and features to meet future needs of our customers even in the field of secondary metallurgy. Considering present challenges for iron- and steel makers facing worsening scrap qualities and in parallel higher requests for steel grades and qualities Tenova with its long-term experience, know-how and technologies can support the transformation of making steel in future in a more sustainable, environmental-friendly and economic way.
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Aim news Calendario degli eventi internazionali International events calendar 2019 March 25-26, Sydney, Australia, Future of Mining Australia 2019 March 25-29, Khao Lak, Thailand, 11th International Conference on the Science of Hard Materials (ICSHM11) April 3-5, Auckland, New Zealand, PACRIM 2019: Mineral Systems of the Pacific Rim May 6-9, Houston, USA, Offshore Technology Conference (OTC) 2019 May 18-25, Perth, Australia, ALTA 2019 Nickel-Cobalt-Copper, Uranium-REE, Gold-PM, In Situ Recovery and Lithium Conference & Exhibition
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May 22-24, Brno, Czech Republic, 28th International Conference on Metallurgy and Materials (Metal 2019)
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May 27-29, Sydney, Australia, 9th International Conference on Sustainable Development in the Minerals Industry (SDIMI)
AIM, Via F. Turati 8
June 5-7, Bardolino, Garda Lake, Italy ECHT 2019 - heat treatment & surface engineering for automotive
Tel.: 02 76021132/76397770,
June 10-14, Nantes, France 14th World Conference on Titanium (Ti-2019)
e-mail: amm.aim@aimnet.it
June 13-15, Guangzhou, China 2019 China International Metal & Metallurgy Exhibition
AIM e per l’acquisto delle
20121 Milano fax: 02 76020551 www.aimnet.it
June 23-27, Portsmouth, USA NUMIFORM 2019: The 13th International Conference on Numerical Methods in Industrial Forming Processes June 25-29, Düsseldorf, Germany METEC & 4th ESTAD 2019 July 21-25, Indianapolis, USA 5th World Congress on Integrated Computational Materials Engineering (ICME 2019) July 22-24, Perth, Australia Iron Ore Conference 2019 July 22-24, Osaka, Japan, BIT’s 8th Annual Worls Congress of Advanced Materials (WCAM 2019) August 13-15, Toronto, Canada, 8th International Conference on Modeling and Simulation of Metallurgical Processes in Steelmaking (STEELSIM2019) September 1-5, Stockholm, Sweden, EUROMAT 2019 September 9-13, Cracow, Poland, EUROCORR 2018 September 30 - October 2, Graz, Austria, 10th European Stainless Steel Congress, Science and Market - 6th European Conference and Expo Duplex August, Xi’an, China, 10th Pacific Rim International Congress on Advanced Materials an Processing (PRICM10)
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La Metallurgia Italiana - n. 3 2019
Atti e notizie AIM - UNSIDER Norme pubblicate e progetti in inchiesta (aggiornamento 28 febbraio 2019) NORME UNSIDER PUBBLICATE DA UNI NEL MESE DI FEBBRAIO 2018 UNI EN ISO 6506-2:2019 Materiali metallici - Prova di durezza Brinell - Parte 2: Verifica e taratura delle macchine di prova UNI EN ISO 19277:2019 Industrie petrolifere, petrolchimiche e del gas naturale - Prove di qualificazione e criteri di accettazione per sistemi di rivestimento protettivo sotto isolamento EC 1-2019 UNI EN ISO 8434-1:2018 Connessioni di tubi metallici per oleoidraulica e pneumatica e per impieghi generali - Parte 1: Raccordi conici a 24° EC 1-2019 UNI EN ISO 6892-2:2018 Materiali metallici - Prova di trazione - Parte 2: Metodo di prova a temperatura elevata UNI EN ISO 4945:2019 Acciaio - Determinazione del contenuto di azoto - Metodo spettrofotometrico UNI EN 10058:2019 Piatti e larghi piatti di acciaio laminati a caldo per impieghi generali - Dimensioni e tolleranze sulla forma e sulle dimensioni UNI CEN/TR 10364:2019 Acciai e ghise - Determinazione delle sostanze elencate nelle direttive 2011/65 / UE (RoHS) e 2000/53 / CE (ELV) – Limitazioni
NORME UNSIDER RITIRATE DA UNI NEL MESE DI FEBBRAIO 2019 UNI EN ISO 6506-2:2015 Materiali metallici - Prova di durezza Brinell - Parte 2: Verifica e taratura delle macchine di prova UNI EN ISO 4945:2009 Acciaio - Determinazione del contenuto di azoto - Metodo spettrofotometrico UNI EN 10058:2004 Barre di acciaio piane laminate a caldo per impieghi generali - Dimensioni e tolleranze sulla forma e sulle dimensioni
PREN E ISO/DIS – MARZO 2019 NORME UNSIDER PUBBLICATE DA CEN E ISO NEL MESE DI FEBBRAIO 2019 EN ISO 683-3:2019 Heat-treatable steels, alloy steels and freecutting steels - Part 3: Case-hardening steels (ISO 683-3:2019)
prEN ISO 13680 Petroleum and natural gas industries Corrosion-resistant alloy seamless tubular products for use as casing, tubing, coupling stock and accessory material - Technical delivery conditions (ISO/DIS 13680:2019)
ISO 15630-1:2019 Steel for the reinforcement and prestressing of concrete -- Test methods -- Part 1: Reinforcing bars, rods and wire
prEN ISO 945-1 Microstructure of cast irons - Part 1: Graphite classification by visual analysis (ISO/FDIS 945-1:2019)
ISO 15630-2:2019 Steel for the reinforcement and prestressing of concrete -- Test methods -- Part 2: Welded fabric and lattice girders
prEN 877 Cast iron pipes systems for the evacuation of water from works - Characteristics and test methods
ISO 15630-3:2019 Steel for the reinforcement and prestressing of concrete -- Test methods -- Part 3: Prestressing steel
prEN ISO 439 Steel and cast irons - Determination of silicon content - Gravimetric method (ISO/DIS 439:2019)
ISO 10893-6:2019 Non-destructive testing of steel tubes -Part 6: Radiographic testing of the weld seam of welded steel tubes for the detection of imperfections
prEN 10217-7 Welded steel tubes for pressure purposes - Technical delivery conditions - Part 7: Stainless steel tubes
ISO 10893-7:2019 Non-destructive testing of steel tubes -Part 7: Digital radiographic testing of the weld seam of welded steel tubes for the detection of imperfections
PROGETTI UNSIDER MESSI ALLO STUDIO DAL CEN (STAGE 10.99) – MARZO 2019 prEN ISO 19901-2 rev Petroleum and natural gas industries - Specific requirements for offshore structures - Part 2: Seismic design procedures and criteria prEN ISO 8062-3 rev Steel and cast iron - Determination of copper content - Flame atomic absorption spectrometric method
PROGETTI UNSIDER IN INCHIESTA La Metallurgia Italiana - n. 3 2019
prEN – progetti di norma europei
ISO/DIS – PROGETTI DI NORMA INTERNAZIONALI ISO/DIS 20088-2 Determination of the resistance to cryogenic spill of insulation materials -- Part 2: Vapour release ISO/DIS 19901-10 Petroleum and natural gas industries -- Specific requirements for offshore structures -Part 10: Marine geophysical investigations ISO/DIS 13765-7 Refractory mortars -- Part 7: Determination of permanent change in dimensions on heating ISO/DIS 13680 Petroleum and natural gas industries -Corrosion-resistant alloy seamless tubular products for use as casing, tubing, coupling stock and accessory material -- Technical delivery conditions
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Aim news ISO/DIS 11971 Steel and iron castings -- Visual testing of surface quality ISO/DIS 9647 Steels -- Determination of vanadium content -- Flame atomic absorption spectrometric method (FAAS) ISO/DIS 439 Steel and cast irons -- Determination of silicon content -- Gravimetric method
PROGETTI UNSIDER AL VOTO FPREN E ISO/FDIS – MARZO 2019 FprEN – progetti di norma europei FprEN ISO 10418 Petroleum and natural gas industries - Offshore production installations - Process safety systems (ISO/FDIS 10418:2019) FprEN ISO 19904-1 Petroleum and natural gas industries - Floating offshore structures - Part 1: Ship-shaped, semi-submersible, spar and shallowdraught cylindrical structures (ISO/FDIS 19904-1:2019)
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FprCEN/TR 10367 Alloyed steels - Determination of chromium content - Inductively coupled plasma optical emission spectrometric method FprEN 10136 Steels and cast irons - Determination of nickel content - Flame atomic absorption spectrometric method (FAAS) FprEN 10177 Steels - Determination of calcium content - Flame atomic absorption spectrometric method (FAAS) FprEN 10181 Steels - Determination of lead content - Flame atomic absorption spectrometric method (FAAS)
ISO/FDIS – PROGETTI DI NORMA INTERNAZIONALI ISO/FDIS 19906 Petroleum and natural gas industries -- Arctic offshore structures ISO/FDIS 19904-1 Petroleum and natural gas industries -- Floating offshore structures -- Part 1: Ship-sha-
ped, semi-submersible, spar and shallowdraught cylindrical structures ISO/FDIS 10418 Petroleum and natural gas industries -- Offshore production installations -- Process safety systems ISO/FDIS 5000 Steel sheet, aluminium-silicon alloy-coated by the continuous hot-dip process, of commercial and drawing qualities ISO/FDIS 4701 Iron ores and direct reduced iron -- Determination of size distribution by sieving ISO/PRF 2597-2 Iron ores -- Determination of total iron content -- Part 2: Titrimetric methods after titanium(III) chloride reduction ISO/FDIS 945-1 Microstructure of cast irons -- Part 1: Graphite classification by visual analysis ISO/FDIS 945-4 Microstructure of cast irons -- Part 4: Test method for evaluating nodularity in spheroidal graphite cast irons
La Metallurgia Italiana - n. 3 2019
www.aimnet.it/echt2019.htm
Heat Treatment & Surface Engineering for Automotive
Organised by
REGISTRATION FEES Early bird registration fees (by May 8, 2019) SPEAKER (presenter) SESSION CHAIRPERSON
MEMBER (*)
NON MEMBER
€ 480
€ 580
COMMITTEE MEMBER PARTICIPANT (non-presenter) EXHIBITOR / SPONSOR
€ 690
€ 590
STUDENT **
Patronized by
€ 360 € 630
€ 530
assofluid
REGISTRATION FEES Standard registration fees (after May 8, 2019) SESSION CHAIRPERSON COMMITTEE MEMBER PARTICIPANT (non-presenter)
MEMBER (*)
NON MEMBER
€ 480
€ 580
COMITATO ITALIANO DEI COSTRUTTORI DI FORNI INDUSTRIALI th
7 international congress on science and technology of steelmaking the challenge of industry 4.0
€ 400 € 530
P roduttori Italiani e B ulloneria 26 Viteria convegno nazionale trattamenti termici o
Sponsored by
€ 630
HEAT TREATMENT
* AIM, ASSIOT, ASSOFLUID, CICOF and UPIVEB Member ** Students will have to provide valid proof of student status. The social event on June 6 is not included in the student registration fee.
(Includes only the social event on June 6)
CONFERENCE REGISTRATION FEES INCLUDE:
Admittance to technical sessions and to the exhibition Conference bag with electronic proceedings Social event on June 6 coffee breaks lunches
For non-members the fee includes AIM Membership for the second half of 2019 and for the year 2020.
Accredited TM
GROUP
Additional ticket for Social event for accompanying persons: € 122 (22% VAT included)
• • • • •
convegno nazionale Unione termici trat tamenti
€ 740
€ 640
STUDENT ** EXHIBITOR / SPONSOR
FEDERATA
EXHIBITION & SPONSORSHIP
registration informaation
5-7 June 2019
Heat Treating
The Conference will feature a table-top exhibition that will represent many areas of industry with latest equipment, facilities and instruments, products and services in the field of heat treatment and surface engineering. Companies will be able to reinforce their participation and enhance their corporate identification by taking advantage of benefits offered to them as Contributing Sponsors of the Conference. Companies interested in taking part in the table-top exhibition or sponsoring the Conference may contact the Organising Secretariat.
e c
10
c c
1st announcement call for papers
10th european conference on continuous casting 2020
10
Bari . Italy 17-19 June 2020
e c
www.aimnet.it/eccc2020
c c Organised by
The 10th European Continuous Casting Conference - ECCC 2020 - will be organised by AIM, the Italian Association for Metallurgy, in Bari (Italy) on 17-19 June 2020 with focus on the status and future developments in the casting of steel. The ECCC is a unique forum for the European continuous casting community to exchange views on the status and the future development of the continuous casting process. The Conference program is abreast of the latest developments in control and automation, advanced continuous casting technologies, application of electromagnetic technologies and mechanical devices to improve the core microstructure, the lubrication issues for improving the surface qualities. Steel metallurgical issues will be addressed as well as their physical and numerical simulation. The exchange of experience in operational practice, maintenance and first results from the recently commissioned plants will integrate the program. The Conference aims at promoting the dialogue among the delegates with industrial and academic background and among the participants in former Conferences and new members of the continuous casting community.
Topics • Trends of innovation in casting technologies • New developments and advanced technologies for the casting of slabs, blooms and billets • Ladle and tundish recent metallurgical solutions for steel cleanness • Flow control, refractories and clogging • Mold lubrication and heat transfer • Product quality control: Surface quality and internal soundness • Numerical simulation and modelling (solidification, metallurgy, fluid flow, validation) • Safety and environmental aspects • Continuous casting technologies and circular economy • Operational practice and maintenance • Measurement, automation and process control • Post-processing of semi-finished products (Scarfing, machining and heat treatment) • Modernization and new implementations • Industry 4.0, machine learning and digitalisation
Call for Papers - Abstract Submission Prospective authors wishing to present papers are invited to submit, by 31 October 2019, a tentative title and an abstract of about 300 words (in English), specifying a maximum of two topics for each proposal, to the Organising Secretariat (aim@aimnet.it). The abstract should provide sufficient information for a fair assessment and include the title of the paper, the author’s names and contact details (address, telephone number and e-mail address). The name of the presenting author should be underlined. A poster session might be organized as well. There are two ways to submit papers: • fill in the form on the Conference website at: www.aimnet.it/eccc2020 • send the requested information by e-mail to: aim@aimnet.it.
Contacts ECCC 2020 Organising Secretariat AIM - Associazione Italiana di Metallurgia Via Filippo Turati 8, 20121 Milan - Italy Tel. +39 02 76021132 / +39 02 76397770 aim@aimnet.it - www.aimnet.it/eccc2020