Effect of precipitation on hot formability of high nitrogen steels

Materials Science and Engineering A 528 (2010) 519–525

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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

Effect of precipitation on hot formability of high nitrogen steels E. Erisir a,b,∗ , U. Prahl a , W. Bleck a a b

Department of Ferrous Metallurgy, RWTH Aachen University, Germany Department of Metallurgical and Materials Engineering, Kocaeli University, Turkey

a r t i c l e

i n f o

Article history: Received 26 July 2010 Received in revised form 10 September 2010 Accepted 15 September 2010

Keywords: High nitrogen steels Precipitation Forging Crack initiation Hot formability

a b s t r a c t This paper examines the effect of precipitation on hot formability of high nitrogen steels (HNSs) through modeling of the precipitation and experimental simulation of the hot deformation. The experiments were carried out on three grades 1.4501, 1.4882, 1.4452 of HNSs and austenitic grade 1.4301 steel. In the course of modeling, precipitation characteristics and phase transformations were investigated by calculation of phase diagrams and precipitation analysis using SEM/EDX. High temperature properties were determined at deformation temperatures between 900 ◦ C and 1300 ◦ C using hot tensile tests in order to investigate the hot formability. The effect of deformation temperature on crack sensibility was determined using compression test. The effects of carbide/nitride precipitation on hot formability in HNSs were discussed. © 2010 Elsevier B.V. All rights reserved.

1. Introduction High nitrogen steels (HNSs) exhibit higher yield and tensile strength as well as improved high temperature strength, compare to conventional austenitic steels. Nitrogen is a strong austenitestabilizing element and improves the corrosion resistance, pitting resistance and stress corrosion cracking resistance [1,2]. In addition, nitrogen alloying can be an alternative to the high cost of nickel alloying [3]. However, HNSs still present several limitations. Edge and surface cracks as well as forging cracks can appear during hot deformation. These types of cracks can be generated by secondary phases, segregations, dynamical recovery, strain induced precipitations, and technical forming parameters [4–8]. It has been reported that the surface and edge cracks of blocks are caused by the precipitation of carbide/nitride and intermetallic phases [4]. The precipitation of Cr2 N has a detrimental effect on the formation of these cracks. Intermetallic sigma phase can also form cracks, especially in duplex steels. The purpose of the present study is to reveal the hot formability in HNSs with particular emphasis on precipitation. Phase diagrams were calculated with Thermo-Calc software in order to determine precipitation characteristics and phase transformations. The investigation of precipitations after homogenization annealing helps to draw conclusions from appearing phases and dissolving tempera-

∗ Corresponding author at: Department of Metallurgical and Materials Engineering, Kocaeli University, Turkey. Fax: +90 2623033003. E-mail address: [email protected] (E. Erisir). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.09.045

tures. Hot formability of investigated steels was determined via hot tensile and compression tests at elevated temperatures between 900 ◦ C and 1300 ◦ C. 2. Experimental procedure 2.1. Materials The experimental samples were produced from pressurized electroslag remelting (PESR)-blocks of industrial heats. DIN EN 10088-1 grades 1.4501, 1.4882, and 1.4452 of high nitrogen steel, with nitrogen amount up to 0.800 wt.%, were investigated [9]. In addition, austenitic grade 1.4301 stainless steel was chosen as reference material. Grades 1.4301 and 1.4501 are Ø350 mm hot roll formed bars, while grades 1.4882 and 1.4452 are 400 mm square ingots. Their chemical compositions are listed in Table 1. Super duplex grade 1.4501 HNS is characterized by its high corrosion resistance and ferritic–austenitic microstructure. Austenitic grade 1.4882 HNS is commonly used as valve material and highly alloyed with the carbide/nitride stabilizer niobium. Austenitic grade 1.4452 HNS is nickel-free and used in medical technology. 2.2. Investigations on precipitation The calculations were made with the Thermo-Calc software using the TCS Steel Database TCFE5 [10,11]. The aim of the calculations was the prediction of phase equilibrium for various temperatures. All elements present in the examined steels were considered. The homogenization annealing was carried out in a Bähr DIL805 plastodilatometer in order to investigate the precipitation and phase transformations. The standard specimen size of 5 mm diameter and 10 mm length was used for the dilatometer experiments. The samples were annealed at 1000 ◦ C, 1100 ◦ C and 1200 ◦ C for 15 min, followed by cooling with cooling rates 200 K/s and 0.003 K/s. The samples were metallographically prepared. After grinding, they were polished using diamond pastes with particle sizes of 6, 3 and 1 ␮m. Different etchants like Nital, V2A and Beraha were applied. Scanning electron microscope (SEM) attached EDX was used to observe the morphology of secondary phase particles and to analyze their

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Table 1 Chemical composition (wt.%) of the examined steels. Steel grade

Alloy

C

Mn

Cr

Mo

Ni

W

Cu

Nb

N

1.4301 1.4501 1.4882 1.4452

X5CrNi18–10 X2CrNiMoCuWN25–7–4 X50CrMnNiNbN21–9 X13CrMnMoN18–14–3

0.023 0.022 0.480 0.070

1.86 0.62 8.45 12.37

18.2 25.6 20.6 16.1

0.29 3.65 0.21 3.26

8.35 6.15 3.60 0.12

– 0.55 0.83 –

0.21 0.90 0.15 0.04

– 0.009 2.000 0.008

0.085 0.256 0.480 0.800

chemical compositions. In addition, metallographic observations were made with an optical microscope (OM). Quantitative metallography was carried out with imageanalyzer software using area analysis. The fractions of phases were determined at magnifications of 200 and 500, depending of the size of phase. 2.3. Hot formability In order to investigate the hot formability, the hot tensile and compression tests were performed using the Gleeble hot deformation simulator at deformation temperatures between 900 ◦ C and 1300 ◦ C with a strain rate of 1/s. The cylindrical samples for tensile and compression tests are shown in Fig. 1. The samples were pre-annealed at 1150 ◦ C for 120 min under a gas mixture of 5 vol.% N2 and 95 vol.% CO2 and cooled in a furnace. After the pre-annealing, samples were tested using the hot deformation simulator. Through the tensile test, the maximum load and the reduction of area were measured. The samples were heated to 1250 ◦ C with a rate of 5 K/s and soaked for 15 min for homogenization annealing. A typical deformation temperature range between 900 ◦ C and 1300 ◦ C was chosen. The deformation rate of 1/s was applied during the tests before the samples cooled down to room temperature with a cooling rate of 50 K/s. Argon was used as protective and cooling gas. The temperature was controlled with a tolerance of ±5 K using a thermocouple. The reduction of area was plotted as a function of test temperature in order to determine the hot deformability. The fracture surfaces of tensile test samples were also investigated by SEM. The hot compression tests were performed using the hot deformation simulator at different deformation degrees up to crack formation on the collar samples in Fig. 1b. The aim was to determine a forging window. The collar compression samples were chosen because of their localized crack susceptibility to reduce experimental effort for identification of cracking locations. In the beginning of the compression tests, an initial degree of deformation was chosen from the results of tensile tests. The collar samples were compressed with the initial degree of deformation and examined for the surface cracks with a stereomicroscope. If a crack was not present in collar of the samples, deformation degree was increased or vice versa. Thus, a limit of the degree of deformation is determined without causing cracks on collar samples.

3. Investigations on precipitation 3.1. Thermodynamic calculations The Thermo-Calc software was used to predict the stability of occurred phases and precipitates. Fig. 2 is an isopleth diagram showing stability of phases as a function of temperature and nitrogen or carbon content. The dashed line in each diagram indicates the relevant nitrogen or carbon amount for steel compositions used here.

Fig. 1. Schematic representation of cylindrical samples for: (a) hot tensile and (b) compression tests.

The results of thermodynamic calculations performed for the temperature range of 600–1600 ◦ C are summarized in Table 2 for the investigated compositions. According to the calculations, intermetallic phases and carbides/nitrides are in equilibrium. Carbides/nitrides are M2 N, M(C,N), M23 C6 , and M6 C. Intermetallic sigma and chi (␹) phases can also precipitate. Typical deformation temperatures lie between 900 ◦ C and 1300 ◦ C for investigated steels. Intermetallic sigma and ␹ phases do not seem critical for austenitic grades 1.4882 and 1.4452, because the calculated solution temperatures are below the minimum deformation temperature of 900 ◦ C. Sigma phase can be seen up to 1003 ◦ C for duplex grade 1.4501. However, carbides/nitrides may exist within the temperature range of deformation. M2 N nitrides are stable in grades 1.4501 and 1.4452. For austenitic grade 1.4882 HNS, M23 C6 and M(C,N) were predicted to be in equilibrium with austenite at 1103 ◦ C and 1344 ◦ C, respectively. 3.2. Homogenization annealing Two series of experiments were performed to evaluate the precipitations of examined steels. The samples were heated to temperatures of 1000 ◦ C, 1100 ◦ C, and 1200 ◦ C, then held at these temperatures for a duration of 15 min in order to dissolve any precipitates, and finally cooled to room temperature at linear cooling rates of 0.003 K/s (slow cooling) or 200 K/s (quenching). The aim of two different cooling rates was to reveal the precipitates with slow cooling and to determine their dissolving temperatures after quenching. 3.2.1. Slow cooling after homogenization annealing In investigations of samples after slow cooling with 0.003 K/s, a variety of nitrogen and carbon-containing precipitates was found by performing SEM (Fig. 3). EDX analysis was employed to determine the composition of the precipitates. Three types of carbides/nitrides M2 N, M(C,N) and M23 C6 and two different intermetallic phases sigma and ␹ were identified. Two kinds of precipitates are observed in austenite: (i) intermetallic phases and (ii) nitrides and carbides [1]. Intermetallic phases are often sigma phase and chi-phase (␹, Fe36 Cr12 Mo10 ). As nitrides/carbides, Cr2 N, M23 C6 , M(C,N) and M6 C may appear [12–15]. Austenitic grade 1.4301 has shown ferrite and austenite phases without any precipitate, as presented in Fig. 3a. The ferrite phase is rich in ferrite forming element Cr, as the austenite is rich in austenite forming element Ni. In duplex steels, sigma phase precipitates in ferrite phase [16] as well as in austenite phase for long annealing periods [15]. In addition, the presence of Mo and Cr in the composition of these steels, can lead to the formation of ␹-phase [17]. SEM/EDX investigations of super duplex grade 1.4501 showed ferrite, austenite, sigma and ␹ phases (see Fig. 3b). The quantitative EDX was performed to estimate the partitioning of the elements in the phases. An enrichment of ferrite forming elements Mo and Cr as well as decreased amount of Ni is observed in ferrite phase. The precipitates are seen in ferrite areas and give chemical composition difference in EDX-analysis. In intermetallic sigma phase, 22 wt.% Cr and 4 wt.% Mo, 8 wt.% Ni with traces of Cu and W are measured. Intermetallic ␹ phase differs from sigma phase with a chemical composition of 25 wt.% Cr

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Fig. 2. Calculated phase diagram for grades: (a) 1.4301, (b) 1.4501, (c) 1.4882, and (d) 1.4452.

and 11 wt.% Mo, 2 wt.% W and 4 wt.% Ni. Nitrogen and carbon are dissolved predominantly in austenite, which could not be exactly quantified via EDX-analysis. According to some studies in reference [13,18,19] and calculations in Table 2, M23 C6 and M2 N can precipitate in austenite. However, they were not observed due to inadequate annealing time for their precipitation. M(C,N) can precipitate in titanium or niobium stabilized austenitic steels [1]. Several investigations [12,20–23] on lower carbon grades of HNS have shown that secondary precipitates M23 C6 and M2 N form at grain and twining boundaries as well as appear in lamellar or dispersed form within austenite grains. The lamellar M2 N nitrides are called nitrogen pearlite. In high Nb alloyed grade 1.4882, primary cubic M(C,N) and secondary carbides/nitrides M23 C6 and M2 N can be observed in austenite. Fig. 3c indicates an austenitic microstructure with primary M(C,N) and nitrogen pearlite of M23 C6 /M2 N clusters. 80 wt.% Nb and small amounts of Fe and Cr were measured in M(C,N) phase. M23 C6

and M2 N phases precipitate together and form clusters. Thus, these phases cannot be determined as pure carbide or nitride. However, the cluster has a chemical composition of 50 wt.% Cr, 7% Mn, 5.5% W and 2% Mo. Intermetallic sigma phase was not observed. As can be seen in Fig. 3d, in grade 1.4452, lamellar nitrogen pearlite forms at grain boundaries of coarse austenite grains between ASTM 0 and −2. Nitrogen pearlite phase contains about 25 wt.% Cr and 8 wt.% Mo. According to thermodynamic calculations, small amounts of M6 C and M23 C6 carbides and sigma phase are also stable. However, these phases could not be detected. 3.3. Quenching after homogenization annealing Homogenization annealing showed that precipitations can be dissolved at 1200 ◦ C for grades 1.4301, 1.4501, and 1.4452 (Fig. 4).

Table 2 Calculated solution temperatures of precipitations at equilibrium. Steelgrade

Precipitate Intermetallic phase

Carbide/nitride ◦

1.4301 1.4501 1.4882 1.4452

◦

Sigma phase (Fe,Ni)x (Cr,Mo)y ( C)

␹-phase Fe36 Cr12 Mo10 ( C)

M23 C6 (◦ C)

M6 C (◦ C)

M2 N (◦ C)

M(C,N) (◦ C)

732 1003 754 792

– 817 – –

831 851 1103 800

– – – 879

833 1100 813 1059

– – 1344 –

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Fig. 3. SEM micrographs showing microstructures after slow cooling (1100 ◦ C/15 min/0.003 K/s): (a) austenite microstructure with retained ferrite; 1.4301, (b) duplex ferrite–austenite microstructure containing sigma- and ␹-phases; 1.4501, (c) austenite microstructure with Nb-primary carbonitrides and nitrogen pearlite including M23 C6 ; 1.4882, and (d) austenite microstructure having nitrogen pearlite; 1.4452.

Fig. 4. Microstructures after quenching (1200 ◦ C/15 min/200 K/s): (a) austenite microstructure with retained ferrite; grade 1.4301, (b) duplex ferrite–austenite microstructure; grade 1.4501, (c) austenite microstructure including Nb-primary carbonitrides; grade 1.4882, and (d) austenite microstructure; grade 1.4452.

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Table 3 Summary of metallographic investigations. Steel grade

Microstructure

Precipitates Intermetallic phase

Carbide/nitride

1.4301 1.4501 1.4882

Austenite with retained ferrite Duplex austenite–ferrite Austenite

– Sigma phase ␹–phase –

– – M23 C6 M2 N M(C,N)

1.4452

Austenite

–

M2 N

Grades 1.4301 and 1.4501 have austenite and ferrite phases without precipitates as shown in Fig. 4a and b. For grade 1.4501, intermetallic sigma- and ␹-phases were not seen after the quenching. The microstructure of grade 1.4882 indicates primary Nb(C,N) as in Fig. 4c. These carbonitrides cannot be dissolved up to annealing temperatures of 1200 ◦ C, because of their formation at liquid phase as expected from thermodynamic calculations in Fig. 2. On the other hand, grade 1.4452 has an austenite microstructure without nitrogen pearlite in Fig. 4d. General remarks on precipitation and microstructure are summarized in Table 3. 4. Investigations on hot deformation 4.1. Hot tensile tests The mechanical properties at high temperatures are determined using hot tensile tests. The plots of reduction of area (RA) and maximum load versus test temperatures for all steels examined in the present work are shown in Fig. 5. The RA is considered as a ductility parameter, while strength is measured as a maximum load. The high temperature strength and ductility of HNSs are intimately related with the evolution of precipitates. As seen from Fig. 5a, lower ductility values are noticed for grades 1.4452 and 1.4882. Grade 1.4452 shows a sharp ductility drop at 1100 ◦ C. The ductility drop goes down to a RA of 5% at 950 ◦ C. In grade 1.4882, RA values are lower than 30% up to 1250 ◦ C and drops to about 0% above this temperature due to partial melting. The solidus temperature was measured as 1296 ◦ C with differential thermal analysis, which is compatible with the observed results. They give also higher maximum loads than grades 1.4301 and 1.4501 in Fig. 5b. A microstructure without precipitates gives higher ductility at deformation temperatures. Grade 1.4501 with ferritic–austenitic duplex microstructure showed the best ductility with a minimum % RA of about 65% at 900 ◦ C. Grade 1.4301 steel has slightly lower values in % RA than grade 1.4501, but has a similar trend. The samples were investigated for interpretation of fracture behavior via metallography at SEM. The fractographical examination shows in Fig. 6 that similar high ductility behavior observed for grades 1.4301 and 1.4501 is consistent with the observation of characteristic dimples associated with transgranular ductile fracture in Fig. 6a and b. Large voids present on the fracture surfaces and apparently not associated with second phase particles. In Fig. 6c, a brittle fracture is seen for grade 1.4882. Primary M(C,N) precipitates are visible in fracture surface. In grade 1.4452 with large austenite grains, a typical intercrystalline cracking promoting brittle fracture is observed in Fig. 6d.

General remarks

– Not stable at deformation temperatures M(C,N) is insoluble M23 C6 and M2 N are stable at lower deformation temperatures Stable at lower deformation temperatures

of HNSs. Reference grade 1.4301 steel was not compressed in order to reduce the number of experiments, since it shows similar curve as grade 1.4501 in the results of hot tensile tests. Fig. 7 gives the forging window in the curve of degree of deformation versus deformation temperature. Above the curves, surface cracks appeared on collar of compressed samples. As seen in Fig. 7, the forging window of grade 1.4501 is significantly extended compare to grades 1.4452 and 1.4882 and shows crack-free surfaces after deformation at temperatures between 1050 ◦ C and 1300 ◦ C. In order to achieve a crack free surface after compression, austenitic grades 1.4452 and 1.4882 should not be compressed with a degree of deformation higher than the values between 0.15 and 0.50 depending on deformation temperature. The lowest formability was seen for grade 1.4882, which gives a maximum deformation degree of only 0.25 at 1200 ◦ C. It seems for grade 1.4452 that the samples compressed at temperatures between 1150 ◦ C and 1250 ◦ C show higher formability. The collar samples can be compressed without crack by a deformation degree of 0.4–0.5 in this temperature range.

4.2. Hot compression tests The hot compression tests with collar specimen were performed to establish a forging window of grades 1.4452, 1.4501 and 1.4882

Fig. 5. Temperature dependency of reduction of area (a) and maximum load (b) for materials studied. Strain rate: 1/s.

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Fig. 6. SEM micrographs of fracture surfaces after hot tensile tests with strain rate 1/s: (a) 1.4301 at 1050 ◦ C and (b) 1.4501, (c) 1.4882 and (d) 1.4452 at 950 ◦ C.

Fig. 7. Forging window for grades 1.4501, 1.4882, and 1.4452 of HNSs.

5. Conclusions Three grades 1.4501, 1.4882, and 1.4452 of HNSs were characterized and compare to grade 1.4301 steel using thermodynamic calculations, metallographic investigations and hot deformation simulations. The results obtained can be summarized as follows: Various precipitations of intermetallic phases and carbides/nitrides were identified in the studied steels. Intermetallic sigma and chi phases were found in the ferrite in grade 1.4501. However, intermetallic phases are not stable at deformation temperatures. M2 N nitride and M23 C6 carbides have been detected in grades 1.4452 and 1.4882 of HNSs. These precipitates may present at lower deformation temperatures up to 1100 ◦ C. In grade 1.4882, M(C,N) particles have high thermodynamic stability and cannot be dissolved at annealing temperatures up to 1200 ◦ C. At deformation temperatures, these particles present in microstructure. Thermo-

dynamic calculations were also consistent with the experimental results. Some predicted phases were not observed because of prolonged equilibrium conditions or their small amounts. Hot formability of HNSs was investigated by hot tensile and compression in the temperature range of 900 ◦ C and 1300 ◦ C at the strain rate of 1/s. Hot ductility curves of tensile tests showed that hot formability is critical for grades 1.4882 and 1.4452. A forging strategy is normally considered regarding to the fact that the strengthening precipitates shall be dissolved after heating up to deformation temperatures. However, grade 1.4882 is alloyed with 2 wt.% Nb in order to increase N solubility, which leads to the formation of hardly soluble M(C,N) precipitates having a negative effect on ductility. The hot ductility of grade 1.4452 is determined by the precipitation of M2 N and grain coarsening during hot deformation. M2 N precipitates as nitrogen pearlite at temperatures below 1100 ◦ C. A forging window was determined by observing crack formation during hot compression. The results of hot compression tests were consistent with the results of tensile tests. Grade 1.4501 has a large forging window compare to grades 1.4882 and 1.4552. Acknowledgements This paper is based on work founded by Federal Ministry of Economics and Technology (BMWi) through the AiF (Arbeitsgemeinschaft industrieller Forschungsvereinigungen “Otto von Guericke”). The authors are grateful to subcommittee “Chemisch beständige Stähle” of steel institute VDEh for technical support. References [1] V.G. Gavriljuk, H. Berns, High Nitrogen Steels, Springer-Verlag, Berlin, 1999. [2] M.O. Speidel, P.J. Uggowitzer, Stickstofflegierte Stähle, Ergebnisse der Werkstoff-Forschung, Band 4, ETH-Zürich, 1991.

E. Erisir et al. / Materials Science and Engineering A 528 (2010) 519–525 [3] M.O. Speidel, M. Zheng-Cui, in: Proceedings HNS 2003 High Nitrogen Steels, vdf Hochschulverlag, Zürich, 2003, pp. 63–73. [4] W. Bleck, F. Hippenstiel, R. Kopp, O. Dürr, Untersuchungen zur Warmumformbarkeit neuer stickstofflegierter nichtrostender Stähle/Institut für Eisenhüttenkunde, Institut für Bildsame Formgebung, RWTH Aachen – AiFAbschlussbericht Bd. 11147N, 2000. [5] W. Bleck, E. Erisir, G. Hirt, V. Schneider, Untersuchungen zur wirtschaftlichen Warmumformung neuer hoch stickstofflegierter nichtrostender Stähle in Abhängigkeit vom Stickstoffgehalt, des Oberflächenzustandes und der Ofenatmosphäre/Institut für Eisenhüttenkunde, Institut für Bildsame Formgebung, RWTH Aachen – AiF-Abschlussbericht Bd. 13888N, 2007. [6] Tikhovskiy, H. Brauer, A. Fischer, Mat.-wiss. Werkstofftech. 33 (2002) 448–452. [7] L.A. Norström, S. Malm:, Scand. J. Met. 8 (1979) 31–33. [8] E. Erisir, U. Prahl, W. Bleck, 1st International Conference on Interstitially Alloyed Steels, 2008, IAS 2008. [9] DIN EN 10088-1, Nichtrostende Stähle-Teil 1, Verzeichnis der nichtrostenden Stähle, 2005. [10] TCFE5, TCS Steels/Fe-alloys database (v.5) for Thermo-Calc, 2008. [11] Thermo-Calc TCC Software User’ Guide, Version Q, Stockholm Technology Park, Stockholm, 2005.

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