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Edge cracking prevention in 2507 super duplex stainless steel by twin-roll strip casting and its microstructure and properties
Journal of Materials Processing Tech. 266 (2019) 246–254
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Journal of Materials Processing Tech. journal homepage: www.elsevier.com/locate/jmatprotec
Edge cracking prevention in 2507 super duplex stainless steel by twin-roll strip casting and its microstructure and properties ⁎
Yan Zhao , Yuan Wang, Shuai Tang, Weina Zhang, Zhenyu Liu
T
⁎
State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang, 110819, PR China
A R T I C LE I N FO
A B S T R A C T
Associate Editor: M. Merklein
In the present work, the microstructure, mechanical properties and corrosion resistance of 2507 super duplex stainless steel (SDSS) fabricated by conventional hot-rolling and strip casting were investigated, respectively. Large deformation in the course of conventional hot-rolling destroyed the coherency of ferrite/austenite interphase and leaded to a deviation from the ideal K-S orientation relationship, which accelerated the precipitation of sigma phase at the phase interface in 2507 SDSS. These brittle particles provided numerous potential void nucleation sites and promoted the propagation of cracks along the ferrite/austenite interface, resulting in severe cracking during conventional hot-rolling. High-quality thin strip without edge cracking was manufactured by strip casting and the edge cutting-off of hot-rolled products was avoided, leading to a significant reduction in cost compared to the conventional process. Meanwhile, the mechanical properties and corrosion resistance of 2507 SDSS fabricated by this new technology reached the same standard of conventional productions. High Temperature-Short Time (HTST) heat treatment was applied in the preparation of the final product, which significantly improved the strength of 2507 SDSS.
Keywords: 2507 super duplex stainless steel Edge cracking Conventional hot-rolling Strip casting Heat treatment
1. Introduction Super duplex stainless steels (SDSS) with a two-phase microstructure comprised of approximately equal amounts of austenite and ferrite have been widely used in highly corrosive environments, such as marine and petrochemical industries due to their attractive combination of high strength and good corrosion resistance (Chail and Kangas, 2016). Compared with conventional duplex stainless steel, SDSS has greater additions of Cr, Mo and N to improve the corrosion resistance. As one of the SDSS grades, SAF 2507 with a pitting resistance equivalent number (PREN = wt% Cr + 3.3 wt% Mo + 16 wt% N) value higher than 42 has been receiving wide interest by Cabrera et al. (2003); Kingklang and Uthaisangsuk (2017); Pardal et al. (2009) and Ramkumar et al. (2014). However, it possessed a high deformation resistance and poor thermal conductivity due to the high content of alloying elements, resulting in serious difficulties in hot working as delineated by Cabrera et al. (2003); Kingklang and Uthaisangsuk (2017) reported that the difference in the thermal expansion coefficients and deformation behaviors between austenite and ferrite under hot working conditions can tend to form edge cracks, leading to additional operations like grinding, discontinuous processing or scraping, increasing the manufacturing costs and restricting the production of SDSS. Meanwhile, Pardal et al. (2009) found that the SDSS was sensitive ⁎
to precipitation of undesired secondary phases at high hot working or heat treatment temperature, such as σ, χ, and nitrides, which severely deteriorated the hot ductility, mechanical properties and corrosion resistance of the material. Much work so far has focused on the high temperature deformation characteristics and precipitation behaviors of SDSS. The flow behaviors and microstructure evolution during hot deformation have been investigated in SAF 2507 by Kingklang and Uthaisangsuk (2017). They built the constitutive model according to the Zener-Hollomon equation to predict the flow stress curves and indicated that the dynamic recrystallization (DRX) of the austenite was dominant for the softening mechanism. Mishra et al. (2017) reported that the softening was more dominant at higher temperature and lower strain rate in 2507 SDSS. They also pointed out that the presence of both dynamic recovery (DRV) and continuous DRX in the ferrite phase while discontinuous DRX in the austenite phase. In all intermetallic phases, the harmfulness of σ phase is the most prominent. The σ phase is brittle and rich in Cr and Mo, which can severely reduce corrosion resistance and toughness of SDSS. Villanueva et al. (2006) showed that the σ phase can precipitate in a short time with a nose temperature of about 900 °C, following a ‘C’ type kinetic curve. The kinetic behavior of σ phase precipitation (Santos and Magnabosco, 2016) and the effects of σ phase on the mechanical properties (Deng et al., 2009), corrosion resistance
Corresponding author. E-mail addresses: [email protected] (Y. Zhao), [email protected] (Z. Liu).
https://doi.org/10.1016/j.jmatprotec.2018.11.010 Received 2 May 2018; Received in revised form 4 November 2018; Accepted 8 November 2018 Available online 12 November 2018 0924-0136/ © 2018 Elsevier B.V. All rights reserved.
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Table 1 Chemical composition of the 2507 super duplex stainless steel, wt%.
Twin-roll cast strip Conventional cast ingot
C
Cr
Ni
Mn
Mo
N
Si
Fe
0.02 0.02
24.9 25.4
6.8 6.9
0.69 0.72
3.95 3.89
0.29 0.28
0.18 0.15
Bal. Bal.
(Martins and Casteletti, 2009) and welding performance (Wang et al., 2013) during aging treatments were investigated in a comprehensive series of studies on SDSS. Twin-roll strip casting is one of the ‘near-net-forming’ technologies, which integrates casting and hot rolling to produce thin strips in the thickness of several millimeters (Santos et al., 2000). It possessed unique advantages in the fabrication of hard deformable metals and can inhibit the risk of forming hot rolling cracks reported by Wang et al. (2009). Edge cracking has become one of the most difficult problems to be tackled in the preparation of thin gauge products for 2507 SDSS. In the present work, thin strip of 2507 SDSS with no edge cracks was fabricated by a laboratory scale twin-roll strip caster. Microstructures of the cast strip at different processing stages were investigated. Meanwhile, the mechanical properties and corrosion resistance were also studied. As a comparison, the similar experiments were performed on 2507 SDSS fabricated by conventional process. The mechanism of edge cracking during conventional hot rolling was explored in detail. 2. Materials and methods 2.1. Material processing Table 1 shows the chemical composition of 2507 SDSS. These steels were cast into strip or ingot. Fig. 1 shows the schematic diagram of two processes in fabrication of 2507 SDSS.
Fig. 2. Surface appearance of (a) CHR, (b) CHR-CR, (c) CS, (d) CSHR and (e) CSHR-CR.
30mm→22mm→16 mm→12mm→8mm→5 mm→4 mm). Subsequently, the hot-rolled plates were cold rolled into 1 mm sheets after annealing treatments and acid pickling. The total reduction during cold rolling was about 75%. The conventional hot-rolled plate and coldrolled plate were named as CHR and CHR-CR, respectively.
2.1.1. Conventional processing Fig. 1(a) shows the schematic diagram of conventional processing to manufacture 2507 SDSS. The cast ingot with thickness of about 80 mm was homogenized at 1250 °C for 2 h and hot rolled into about 4 mm in thickness. The total reduction during hot rolling was about 95% (9 passes were employed during hot rolling, 80mm→60mm→42mm→
Fig. 1. Schematic diagram of two processes in fabrication of 2507 SDSS (a) conventional processing and (b) strip casting. 247
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Fig. 3. The cracks in CHR in Fig. 2(a).
Fig. 4. (a) Optical micrograph and (b) EPMA micrograph observations of cracks in CHR.
248
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Fig. 5. (a) Spherical precipitations and (b) rod-shaped precipitation distributed along the ferrite/austenite interface.
in hot-rolled plate were observed by using a 200 KV field emission transmission electron microscopy (TEM). Thin foils were prepared by mechanical polishing down to 50 μm followed by twin jet electro-polishing at a voltage of 32 V at −25 °C.
Table 2 Chemical composition of the precipitations, ferrite and austenite in CHR (Fig. 5), measured by EDS-TEM, wt%.
A B C D E (ferrite) F (austenite)
Cr
Mo
Ni
Mn
Fe
32.76 36.30 30.96 27.08 27.41 24.96
6.14 6.21 10.20 0.19 4.74 3.40
3.69 3.15 4.40 7.23 4.55 7.20
1.15 0.46 1.75 0.55 1.35 1.74
56.01 54.45 52.56 64.56 60.56 62.29
3. Results 3.1. Cracks in conventional hot-rolled plate Fig. 2 shows the photographs of the strips in 2507 SDSS fabricated by conventional hot rolling and twin-roll strip casting. Severe edge cracks were formed in the conventional hot-rolled plate (named as CHR), as presented in Fig. 2(a). The edge cracking of CHR plate firstly appeared on the fourth pass of hot rolling. The CHR with edge cutting treatment was cold rolled into the final thickness of 1 mm (named as CHR-CR) after intermediate softening and acid pickling, Fig. 2(b). The CHR-CR became very hard and brittle after the deformation of about 75%, which was difficult to continue to deform. It needed to be softened again to manufacture thinner products with a thickness below 1 mm. By contrast, no cracks were observed in the cast strip of 2507 SDSS. The cast strip was hot rolled and cold rolled into the final thickness of 0.5 mm with good surface quality, and no obvious edge cracks have been formed during rolling, Fig. 2(c)∼(e). Fig. 3 shows the micrographs of regions containing cracks in CHR in Fig. 2(a). Severely elongated structures were observed in these micrographs, with dark colored and light colored areas representing the ferrite and austenite phases, respectively. A microstructure analysis in the regions adjacent to the cracks was performed in order to study the location of the cracks with respect to ferrite and austenite. The cracks originating from the edge of the hot-rolled plate would be propagated along the interface between ferrite and austenite, Fig. 3(a). With an increase in the level of rolling deformation, the major cracks surrounded with some tiny branches of micro-cracks penetrated deeper into the steel plate, Fig. 3(b). Fig. 4(a) shows some shear localization and cracking in the conventional hot-rolled plate (CHR) of 2507 SDSS. The shear stresses appeared in duplex stainless steel due to the presence of two phases with different mechanical properties. The location of the cracks with respect to the austenite and ferrite indicated their intergranular characteristics. The hardness measurement results of different constituent phases in the CHR were marked in Fig. 4(b). The hardness of the austenite phase was similar to that of the ferrite phase. However, some particles with rich Cr and Mo were observed at the interface of the ferrite and austenite. The hardness of these particles was far higher than that of the ferrite or austenite phase. The great difference in hardness between brittle particles and the ferrite-austenite matrix was one of the important reasons
2.1.2. New processing (Strip casting) Fig. 1(b) shows the schematic diagram of the new processing (strip casting) to fabricate 2507 SDSS. A vertical type twin-roll strip caster was used. The steel was melted in an induction furnace under the N2 shielding. The steel liquid was poured into a preheated tundish to flow through a nozzle into the water-cooled casting rollers and was cast to a strip with the thickness of about 2.5 mm. The speed of strip casting was set to be 0.25 m/s. After having been homogenized at 1050 °C for 3 min, the cast strip was hot rolled from 2.5 mm to 2.0 mm with the reduction of about 20%. Subsequently, the hot rolled strip was cold rolled into the final thickness of approximately 0.5 mm with the total cold-rolling reduction of about 75% after acid pickling. The cast strip, hot-rolled strip and coldrolled strip were named as CS, CSHR and CSHR-CR, respectively. The cold-rolled strips were solid solution treated and then were cut into different shapes and sizes for microstructure, mechanical and corrosion tests, respectively. 2.2. Microstructure characterizations, mechanical properties and corrosion resistance The microstructures of experimental steels at different processing stages were observed by optical microscope (OM) under suitable magnifications. The micrographs near cracks were observed by OM and EPMA. The hardness of different constituent phases was measured by FM-700 micro-hardness Tester. Tensile specimens with gage length of 25 mm and width of 6 mm were machined parallel to the rolling direction, and tensile tests at room temperature were performed at the strain rate of 2 × 10−3s-1. The pitting corrosion resistance of the experimental steels was investigated in 3.5% NaCl solution by using potentio-dynamic measurements at room temperature with the scan rate of 0.33 mV/s. EBSD mapping was conducted on the plane (RD × ND) in the as-annealed specimens. The acceleration voltage and step size of EBSD analysis was 20 kV and 150 nm, respectively. The precipitations 249
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Fig. 6. Microstructure at different processing stages (a) CHR, (b) CHR-CR, (c) CS, (d) CSHR, and (e) CSHR-CR. RD is rolling direction and ND is normal direction.
approximately 50% austenite and 50% ferrite were observed in CHR. Both the ferrite and austenite phases were elongated along the rolling direction together with a banded structure, Fig. 6(a). Meanwhile, there was a presence of sigma phase precipitations observed at the interface between ferrite and austenite in CHR. The CHR was cold rolled with the reduction of about 75% and it exhibited a significantly refined structure, Fig. 6(b). Fig. 6(c) shows the typical microstructure of the original cast strip (CS). The ferrite matrix was decorated with different forms of austenite such as grain boundary austenite (GBA) formed at the priorferrite grain boundaries, intragranular austenite (IGA) precipitated in the ferrite grains and widmanstatten austenite (WA) grew into the ferrite grains from the GBA. The CSHR was similar to the CS in the microstructure, but its grains were refined due to the one-pass rolling deformation, Fig. 6(d). The microstructure of the cold-rolled strip fabricated by new processing (CSHR-CR) was similar to the product fabricated by conventional processing (CHR-CR). The ferrite and austenite phases were elongated along the rolling direction and the banded structure could not be completely removed even after the subsequent solution treatment, Fig. 6(e). The EBSD phase maps of the present SDSS fabricated by new processing (CSHR-CR) with different heat treatments were shown in Fig. 7. For convenience, the solid solution treatment at 1150 °C for 20 s and the treatment at 1050 °C for 150 s are named as High Temperature-Short
for the initiation of cracks during conventional hot-rolling for 2507 SDSS. Figs. 5(a) and (b) show the typical TEM micrographs in the domains adjacent to the cracks in the CHR. The corresponding chemical composition of the precipitations, ferrite and austenite was measured by EDS-TEM, as presented in Table 2. There were massive of sigma phase precipitations with rich Cr and Mo distributed along the grain boundary or phase interface. Sigma phase often formed at ferrite/austenite interfaces through a heterogeneous nucleation process, then grew toward ferrite, which was attributed to the high Cr and Mo contents and high diffusion rate in ferrite (Ramirez et al., 2003). It had a higher concentration of Cr and Mo than the ferrite or austenite phase, thus, leading to a depletion of these elements in the surrounding regions. The precipitation of sigma phase was hard and brittle, which increased mismatch of hardness between sigma phase and matrix at the phase interface, leading to the deterioration in the ductility of the material. The effects of precipitates on hot-workability would be discussed in detail later. 3.2. Microstructures at different processing stages Fig. 6 shows the optical micrographs of 2507 SDSS at different processing stages. The obvious duplex structures with a mixture of 250
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Fig. 7. Microstructure of (a)∼(c) LTLT-sample and (d)∼(f) HTST-sample obtained by EBSD.
Fig. 8. (a) Engineering stress-engineering strain curves and (b) strength and elongation of 2507 SDSS fabricated by different processing (CHR-CR and CSHR-CR) with solid solution treatment at 1050 °C for 150 s.
(yellow), and deformed (red) regions in ferrite and austenite of LTLTsample, indicating that ferrite and austenite have been almost completely recrystallized. However, there were quite a few substructures in the HTST-sample, Fig. 7(e) and (f).
Time (HTST) heat treatment and Low Temperature-Long Time (LTLT) heat treatment, respectively. The magenta and yellow green patches in the phase maps were ferrite and austenite, respectively, Figs. 7(a) and (d). The phase maps showed the microstructure consists of a bamboolike structure of coarse ferrite grains elongated along the rolling direction and equiaxed austenite grains. The ferrite grains were always coarser than the austenite grains due to the faster recrystallization and growth kinetics of ferrite than austenite regardless of the heat treatment (Kumar et al., 2017). The area of each elongated ferrite grain in the contiguous bands was converted to that of the equiaxed grain in order to facilitate the grain size measurement. The average grain size of ferrite and austenite in the LTLT-sample was measured to be about 3.55 μm and 2.56 μm, respectively. However, the grain size of the HTST-sample was effectively refined, which was measured to be about 1.96 μm for ferrite and 1.15 μm for austenite, respectively. Figs. 7(b) and (c) show the distribution of the recrystallized (blue), substructured
3.3. Mechanical properties Fig. 8(a) shows the engineering stress-engineering strain curves of 2507 SDSS fabricated by different processing. The corresponding mechanical properties, such as yield strength (YS), ultimate tensile strength (UTS) and total elongation (TE), were compared in Fig. 8(b). The YS, UTS, and TE of CSHR-CR solid solution treated at 1050 °C for 150 s have been measured to be about 675 MPa, 930 MPa and 26%, respectively, which were similar to those of CHR-CR with the same solid solution treatment. Therefore, the mechanical properties of 2507 SDSS fabricated by strip casting can satisfy the corresponding performance 251
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Fig. 9. (a) Engineering stress-engineering strain curves, (b) yield strength, (c) ultimate tensile strength and (d) total elongation of CSHR-CR with different solid solution treatments.
4. Discussion 4.1. Edge cracking prevention by strip casting For the 2507 SDSS, a high deformation resistance and low thermoplasticity due to its high contents of alloying elements (such as Cr, Mo and N) and the incoordination in deformation of two phases would lead to some difficulties in the conventional hot-rolling (Mishra et al., 2017). Severe edge cracks formed during hot-rolling due to their inherently low hot-ductility bring in the cutting-off of edges of final products, leading to an increase in cost. Kingklang and Uthaisangsuk (2017) reported that the hot-ductility of 2507 SDSS depends on many factors, such as chemical composition, deformation temperature, strain rate, phase proportions, precipitations, nature of the interphase boundaries, and softening mechanisms in the constituting phases. Thus, a true understanding of hot-workability during hot-rolling was essential for manufacturing high quality sheets of 2507 SDSS. Fig. 11(a)-(c) show the schematic illustration of cracking taking place in the conventional hot-rolling of 2507 SDSS. The process of conventional hot-rolling began with a reheating stage at about 1250 °C (Duprez et al., 2002). Though increasing the holding time can induce a degree of spheroidization of austenite, a significant fraction of austenite remained undissolved at the reheating temperature and the Widmanstätten microstructure inside the δ-ferrite grains was preserved to some extent (Iris and Suzanne, 2009). Pettersson et al. (2017) reported that the most ferrite/austenite interphase boundaries obey a K–S orientation relationship of (011)δ//(111)γ and [11 1¯] δ//[10 1¯]γ at the beginning stage of hot-rolling. Badji et al. (2008) found that the coherency of ferrite/austenite interphase can inhibit the precipitation of intermetallic compounds. The widmanstätten austenite would be transformed into a fibrous appearance along the rolling direction with increasing hot-rolling reductions. The austenite appeared more dispersedly distributed in the ferritic matrix and the coherency of ferrite/austenite interphase was destroyed, leading to a deviation from the ideal K–S orientation relationship and the formation of random interphase boundaries between the ferrite and austenite (Patra et al., 2016). Lots of precipitations at the interface between ferrite and austenite were observed during conventional hot-rolling in Figs. 4 and 5. Sigma phase precipitation was formed often at δ/γ interfaces through nucleation process. The heterogeneous nucleation of sigma phase depended on the chemical driving force and the interfacial energy. On the one hand, the
Fig. 10. Potentiodynamic polarization curves of specimens for 2507 SDSS fabricated by different processing (CHR-CR and CSHR-CR) with solid solution treatment at 1050 °C for 150 s.
index when compared with the conventional processing. Fig. 9(a) shows the engineering stress-engineering strain curves of CSHR-CR with different solid solution treatments. The corresponding mechanical properties were compared in Fig. 9(b)-(d). The YS and UTS of HTST-sample were measured to be about 750 MPa and 957 MPa, which were about 15% and 3% higher than that of the LTLT-sample, respectively. The heat treatment of HTST had a greater effect on improving YS than on improving UTS of 2507 SDSS. However, the HTSTsample exhibited the TE of about 22%, slightly lower than that of the LTLT-sample. 3.4. Pitting corrosion resistance Fig. 10 shows the potentiodynamic polarization curves of 2507 SDSS fabricated by different processing. The test was carried out in a 3.5% NaCl solution at room temperature. The curve of the CHR-CR solid solution treated at 1050 °C for 150 s nearly coincided with CSHR-CR, indicating that CHR-CR and CSHR-CR had a similar pitting corrosion resistance. The pitting corrosion potential (Epit) of 2507 SDSS in 3.5% NaCl solution was measured to be about 1.0 V. 252
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Fig. 11. Schematic illustration of the microstructure evolution and cracking damage taking place in the hot-rolled processing of 2507 SDSS (a)∼(c) conventional hotrolling, (d) and (e) one-pass hot-rolling after strip casting.
microstructure evolution taking place in the one-pass hot-rolling after strip casting of 2507 SDSS. The hot-rolling compression ratio (h0/hf, h0 and hf is the initial and final thicknesses of hot bands, respectively) of this new processing was measured to be about 1.25, which was far lower than that of conventional hot-rolling (about 20.00). The Widmanstätten microstructure inside the δ-ferrite grains can be preserved after the one-pass hot deformation and the coherency of ferrite/austenite interphase was not destroyed, which was beneficial to inhibiting the precipitation of intermetallic compounds. Therefore, the edge cracking during hot-rolling was completely eliminated by a small reduction of hot deformation in the new processing.
hot-rolling deformation forced the austenite phase to stretch along the rolling direction and increased the length of the interphase between ferrite and austenite, which provided more effective nucleation positions of brittle phase precipitations (Martin et al., 2012). Meanwhile, the massive deformation energy storage during hot-rolling provided the driving force for the nucleation of precipitates (Yutaka and Hiroyuki, 1999). On the other hand, the hot rolling deformation destroyed the coherency of ferrite/austenite interface and made the orientation between ferrite and austenite deviate from the K-S relationship due to the combined effect of strain and softening mechanisms. The deviation from the K-S relationship leaded to an increase in the interface energy, which decreased the activation energy barrier for sigma phase precipitation and accelerated the precipitation of brittle intermetallic compounds (Haghdadi et al., 2017). The edge of the upper and lower surfaces of the slab needed to bear more deformation during the conventional hot-rolling, which increased the additional tensile stress on the surface and promoted the formation of edge or surface microcracks. Metallographic characterization revealed that damage preferentially nucleated near brittle phase particles at the austenite/ferrite interface, as presented in Fig. 4. These precipitations destroyed the continuity of microstructure and provided more original source of microvoids. Meanwhile, they weakened the bonding strength of the δ/γ interface and deteriorated its plastic deformation ability. The interface between matrix and brittle phase precipitation was easily separated under the action of stress, thus, the microviods formed at the interface. Moreover, the accumulation of chromium at the δ/γ interface resulted in the chromium solid solution hardening of the interface, leading to the stress concentration and microvoids nucleation at the δ/γ interface during hot-rolling deformation. The microvoids coalesced into cracks along the ferrite/austenite interface with increasing hot deformation. Therefore, the crack originated from the edge of the upper and lower surfaces of the slab and extended to the center until it ran through the whole thickness direction. The precipitates of brittle intermetallic compounds at the ferrite/austenite interface provided numbers of potential void nucleation sites and the microvoids coalesced into cracks in the subsequent deformation. The cracks originating from edge of slab expanded along the ferrite/austenite interface and caused the more serious cracking. Fig. 11(d) and (e) show the schematic illustration of the
4.2. Effect of the HTST treatment on improving strength The YS and UTS of HTST-sample were about 15% and 3% higher than that of the LTLT-sample, respectively, as presented in Fig. 9. The strength depended strongly on grain size. The relationship of strength and grain size can be described by the Hall-Petch relation, Eq. (1). σy = σ0 + kd−0.5
(1)
Where σy is the yield stress, σ0 is the lattice friction stress, k is a constant and d is the average grain size (Zhang et al., 2012). According to the Hall-Petch relation, the strength of steel was improved by refining its grains. EBSD analysis of the experimental specimens in Fig. 7 showed that the fraction of substructures increased when adopting the HTST treatment, indicating that the enhancement in the mechanical properties was also caused by increasing dislocation density (correlated to the increasing substructures). When the dislocations moved to the grain boundaries, the barrier of grain boundaries must be overcome before the deformation can be transferred from one grain to another. The smaller the grains, the more grain boundaries and the more significant hindrance to dislocation movement. Therefore, the HTST treatment refined the grains of 2507 SDSS and reserved part of substructures, resulting in an improvement in its strength compared with the LTLT treatment. The EBSD analysis presented that the degree of recrystallization for the HTST sample was lower than that for the LTLT sample, causing the internal stress not to be completely eliminated. The residual internal stress in substructures or deformed-structures leaded 253
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to a little decrease in the plasticity of HTST specimen. Therefore, the HTST sample exhibited brittle fracture with only very reduced localization before fracture in comparison to the stress-strain curve of the LTLT sample.
corrosion and mechanical properties of UNS S31803 duplex stainless steel. Corros. Sci. 51, 2969–2975. Duprez, L., Cooman, B.C., Akdut, N., 2002. Deformation behavior of duplex stainless steel during industrial hot rolling. Steel Res. 73, 531–538. Haghdadi, N., Abou-Ras, D., Cizek, P., Hodgson, P.D., Rollett, A.D., Beladi, H., 2017. Austenite-ferrite interface crystallography dependence of sigma phase precipitation using the five-parameter characterization approach. Mater. Lett. 196, 264–268. Iris, A.-A., Suzanne, D.-M., 2009. Great Britain and the United States. Duplex Stainless Steels. ISTE Ltd and John Wiley & Sons, Inc., pp. 31–39. Kingklang, S., Uthaisangsuk, V., 2017. Investigation of hot deformation behavior of duplex stainless steel grade 2507. Metall. Mater. Trans. A 48, 95–108. Kumar, A., Khatirkar, R.K., Chalapathi, D., Kumar, G., Suwas, S., 2017. Microstructure and texture development during cold rolling in UNS S32205 and UNS S32760 duplex stainless steels. Metall. Mater. Trans. A 48, 2349–2362. Martin, G., Yerra, S.K., Bre´chet, Y., Ve´ron, M., Mithieux, J.-D., Che´hab, B., Delannay, L., Pardoen, T., 2012. A macro- and micromechanics investigation of hot cracking in duplex steels. Acta Mater. 60, 4646–4660. Martins, M., Casteletti, L.C., 2009. Microstructural characteristics and corrosion behavior of a super duplex stainless steel casting. Mater. Charact. 60, 150–155. Mishra, M.K., Balasundar, I., Rao, A.G., Kashyap, B.P., Prabhu, N., 2017. On the high temperature deformation behavior of 2507 super duplex stainless steel. J Mater. Eng. Per. 26, 802–812. Pardal, J.M., Tavares, S.S.M., Fonseca, M.C., Souza, J.A., Corte, R.R.A., Abreu, H.F.G., 2009. Influence of the grain size on deleterious phase precipitation in super duplex stainless steel UNS S32750. Mater. Charact. 60, 165–172. Patra, S., Ghosh, A., Kumar, V., Chakrabarti, D., Singhal, L.K., 2016. Deformation induced austenite formation in as-cast 2101 duplex stainless steel and its effect on hot-ductility. Mater. Sci. Eng. A 660, 61–70. Pettersson, N., Wessman, S., Hertzman, S., Studer, A., 2017. High-temperature phase equilibria of duplex stainless steels assessed with a novel In-situ neutron scattering approach. Metall. Mater. Trans. A 48, 1562–1571. Ramirez, A.J., Lipold, J.C., Brandi, S.D., 2003. The relationship between chromium nitride and secondary austenite precipitation in duplex stainless steels. Metall. Mater. Trans. A 34, 1575–1597. Ramkumar, K.D., Thiruvengatam, G., Sudharsan, S.P., Mishra, D., Arivazhagan, N., Sridhar, R., 2014. Characterization of weld strength and impact toughness in the multi-pass welding of super-duplex stainless steel UNS 32750. Mater. Des. 60, 125–135. Santos, C.A., Spim Jr, J.A., Garcia, A., 2000. Modeling of solidification in twin-roll strip casting. J. Mater. Process. Technol. 102, 33–39. Santos, D.C.D., Magnabosco, R., 2016. Kinetic study to predict sigma phase formation in duplex stainless steels. Metall. Mater. Trans. A 47, 1554–1565. Villanueva, D.M.E., Junior, F.C.P., Plaut, R.L., Padilha, A.F., 2006. Comparative study on sigma phase precipitation of three types of stainless steels: austenitic, super ferritic and duplex. Mater. Sci. Technol. 22, 1098–1104. Wang, S.H., Liu, Z.Y., Zhang, W.N., Wang, G.D., Liu, J.L., Liang, G.F., 2009. Microstructure and mechanical property of strip in Fe-23Mn-3Si-3Al TWIP steel by twin roll casting. ISIJ Int. 49, 1340–1346. Wang, Y.Q., Han, J., Wu, H.C., Yang, B., Wang, X.T., 2013. Effect of sigma phase precipitation on the mechanical and wear properties of Z3CN20.09M cast duplex stainless steel. Nuc. Eng. Des. 259, 1–7. Yutaka, S.S., Hiroyuki, K., 1999. Preferential precipitation site of sigma phase in duplex stainless steel weld metal. Scripta Mater. 40, 659–663. Zhang, X.D., Hansen, N., Gao, Y.K., Huang, X.X., 2012. Hall-Petch and dislocation strengthening in graded nanostructured steel. Acta Mater. 60, 5933–5943.
5. Conclusions In the present work, the edge cracking mechanism of 2507 SDSS during conventional hot-rolling was discussed and the high quality thin strips were manufactured by strip casting. The mechanical properties and corrosion resistance of 2507 SDSS fabricated by these two different processes were investigated. The main conclusions are obtained as follows. (1) The brittle intermetallic compounds precipitated at the phase interface in 2507 SDSS during conventional hot-rolling provided numberous potential void nucleation sites, which accelerated the propagation of cracks along the ferrite/austenite interface and caused serious cracking. (2) The edge cracking was completely avoided during one-pass hotrolling after strip casting through a small reduction of hot deformation. Meanwhile, the mechanical properties and corrosion resistance of 2507 SDSS fabricated by this new technology can reach the standard of conventional production. (3) The High Temperature-Short Time (HTST) heat treatment significantly improved the strength of 2507 SDSS, especially improving its yield strength. Acknowledgments This work was supported by the National Natural Science Foundation of China with the contracts of U1460204, U1660117 together with Baosteel Co., and National Natural Science Foundation of China (51774083). References Badji, R., Bouabdallah, M., Bacroix, B., Kahloun, C., Bettahar, K., Kherrouba, N., 2008. Effect of solution treatment temperature on the precipitation kinetic of σ phase in 2205 duplex stainless steel welds. Mater. Sci. Eng. A 496, 447–454. Cabrera, J.M., Mateo, A., Llanes, L., Prado, J.M., Anglada, M., 2003. Hot deformation of duplex stainless steels. J. Mater. Pro. Tech. 143, 321–325. Chail, G.C., Kangas, P., 2016. Super and hyper duplex stainless steels: structures, properties and applications. Procedia Struct. Integr. 2, 1755–1762. Deng, B., Wang, Z., Jiang, Y., Sun, T., Xu, J., Li, J., 2009. Effect of thermal cycles on the
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Journal of Materials Processing Tech. journal homepage: www.elsevier.com/locate/jmatprotec
Edge cracking prevention in 2507 super duplex stainless steel by twin-roll strip casting and its microstructure and properties ⁎
Yan Zhao , Yuan Wang, Shuai Tang, Weina Zhang, Zhenyu Liu
T
⁎
State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang, 110819, PR China
A R T I C LE I N FO
A B S T R A C T
Associate Editor: M. Merklein
In the present work, the microstructure, mechanical properties and corrosion resistance of 2507 super duplex stainless steel (SDSS) fabricated by conventional hot-rolling and strip casting were investigated, respectively. Large deformation in the course of conventional hot-rolling destroyed the coherency of ferrite/austenite interphase and leaded to a deviation from the ideal K-S orientation relationship, which accelerated the precipitation of sigma phase at the phase interface in 2507 SDSS. These brittle particles provided numerous potential void nucleation sites and promoted the propagation of cracks along the ferrite/austenite interface, resulting in severe cracking during conventional hot-rolling. High-quality thin strip without edge cracking was manufactured by strip casting and the edge cutting-off of hot-rolled products was avoided, leading to a significant reduction in cost compared to the conventional process. Meanwhile, the mechanical properties and corrosion resistance of 2507 SDSS fabricated by this new technology reached the same standard of conventional productions. High Temperature-Short Time (HTST) heat treatment was applied in the preparation of the final product, which significantly improved the strength of 2507 SDSS.
Keywords: 2507 super duplex stainless steel Edge cracking Conventional hot-rolling Strip casting Heat treatment
1. Introduction Super duplex stainless steels (SDSS) with a two-phase microstructure comprised of approximately equal amounts of austenite and ferrite have been widely used in highly corrosive environments, such as marine and petrochemical industries due to their attractive combination of high strength and good corrosion resistance (Chail and Kangas, 2016). Compared with conventional duplex stainless steel, SDSS has greater additions of Cr, Mo and N to improve the corrosion resistance. As one of the SDSS grades, SAF 2507 with a pitting resistance equivalent number (PREN = wt% Cr + 3.3 wt% Mo + 16 wt% N) value higher than 42 has been receiving wide interest by Cabrera et al. (2003); Kingklang and Uthaisangsuk (2017); Pardal et al. (2009) and Ramkumar et al. (2014). However, it possessed a high deformation resistance and poor thermal conductivity due to the high content of alloying elements, resulting in serious difficulties in hot working as delineated by Cabrera et al. (2003); Kingklang and Uthaisangsuk (2017) reported that the difference in the thermal expansion coefficients and deformation behaviors between austenite and ferrite under hot working conditions can tend to form edge cracks, leading to additional operations like grinding, discontinuous processing or scraping, increasing the manufacturing costs and restricting the production of SDSS. Meanwhile, Pardal et al. (2009) found that the SDSS was sensitive ⁎
to precipitation of undesired secondary phases at high hot working or heat treatment temperature, such as σ, χ, and nitrides, which severely deteriorated the hot ductility, mechanical properties and corrosion resistance of the material. Much work so far has focused on the high temperature deformation characteristics and precipitation behaviors of SDSS. The flow behaviors and microstructure evolution during hot deformation have been investigated in SAF 2507 by Kingklang and Uthaisangsuk (2017). They built the constitutive model according to the Zener-Hollomon equation to predict the flow stress curves and indicated that the dynamic recrystallization (DRX) of the austenite was dominant for the softening mechanism. Mishra et al. (2017) reported that the softening was more dominant at higher temperature and lower strain rate in 2507 SDSS. They also pointed out that the presence of both dynamic recovery (DRV) and continuous DRX in the ferrite phase while discontinuous DRX in the austenite phase. In all intermetallic phases, the harmfulness of σ phase is the most prominent. The σ phase is brittle and rich in Cr and Mo, which can severely reduce corrosion resistance and toughness of SDSS. Villanueva et al. (2006) showed that the σ phase can precipitate in a short time with a nose temperature of about 900 °C, following a ‘C’ type kinetic curve. The kinetic behavior of σ phase precipitation (Santos and Magnabosco, 2016) and the effects of σ phase on the mechanical properties (Deng et al., 2009), corrosion resistance
Corresponding author. E-mail addresses: [email protected] (Y. Zhao), [email protected] (Z. Liu).
https://doi.org/10.1016/j.jmatprotec.2018.11.010 Received 2 May 2018; Received in revised form 4 November 2018; Accepted 8 November 2018 Available online 12 November 2018 0924-0136/ © 2018 Elsevier B.V. All rights reserved.
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Table 1 Chemical composition of the 2507 super duplex stainless steel, wt%.
Twin-roll cast strip Conventional cast ingot
C
Cr
Ni
Mn
Mo
N
Si
Fe
0.02 0.02
24.9 25.4
6.8 6.9
0.69 0.72
3.95 3.89
0.29 0.28
0.18 0.15
Bal. Bal.
(Martins and Casteletti, 2009) and welding performance (Wang et al., 2013) during aging treatments were investigated in a comprehensive series of studies on SDSS. Twin-roll strip casting is one of the ‘near-net-forming’ technologies, which integrates casting and hot rolling to produce thin strips in the thickness of several millimeters (Santos et al., 2000). It possessed unique advantages in the fabrication of hard deformable metals and can inhibit the risk of forming hot rolling cracks reported by Wang et al. (2009). Edge cracking has become one of the most difficult problems to be tackled in the preparation of thin gauge products for 2507 SDSS. In the present work, thin strip of 2507 SDSS with no edge cracks was fabricated by a laboratory scale twin-roll strip caster. Microstructures of the cast strip at different processing stages were investigated. Meanwhile, the mechanical properties and corrosion resistance were also studied. As a comparison, the similar experiments were performed on 2507 SDSS fabricated by conventional process. The mechanism of edge cracking during conventional hot rolling was explored in detail. 2. Materials and methods 2.1. Material processing Table 1 shows the chemical composition of 2507 SDSS. These steels were cast into strip or ingot. Fig. 1 shows the schematic diagram of two processes in fabrication of 2507 SDSS.
Fig. 2. Surface appearance of (a) CHR, (b) CHR-CR, (c) CS, (d) CSHR and (e) CSHR-CR.
30mm→22mm→16 mm→12mm→8mm→5 mm→4 mm). Subsequently, the hot-rolled plates were cold rolled into 1 mm sheets after annealing treatments and acid pickling. The total reduction during cold rolling was about 75%. The conventional hot-rolled plate and coldrolled plate were named as CHR and CHR-CR, respectively.
2.1.1. Conventional processing Fig. 1(a) shows the schematic diagram of conventional processing to manufacture 2507 SDSS. The cast ingot with thickness of about 80 mm was homogenized at 1250 °C for 2 h and hot rolled into about 4 mm in thickness. The total reduction during hot rolling was about 95% (9 passes were employed during hot rolling, 80mm→60mm→42mm→
Fig. 1. Schematic diagram of two processes in fabrication of 2507 SDSS (a) conventional processing and (b) strip casting. 247
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Fig. 3. The cracks in CHR in Fig. 2(a).
Fig. 4. (a) Optical micrograph and (b) EPMA micrograph observations of cracks in CHR.
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Fig. 5. (a) Spherical precipitations and (b) rod-shaped precipitation distributed along the ferrite/austenite interface.
in hot-rolled plate were observed by using a 200 KV field emission transmission electron microscopy (TEM). Thin foils were prepared by mechanical polishing down to 50 μm followed by twin jet electro-polishing at a voltage of 32 V at −25 °C.
Table 2 Chemical composition of the precipitations, ferrite and austenite in CHR (Fig. 5), measured by EDS-TEM, wt%.
A B C D E (ferrite) F (austenite)
Cr
Mo
Ni
Mn
Fe
32.76 36.30 30.96 27.08 27.41 24.96
6.14 6.21 10.20 0.19 4.74 3.40
3.69 3.15 4.40 7.23 4.55 7.20
1.15 0.46 1.75 0.55 1.35 1.74
56.01 54.45 52.56 64.56 60.56 62.29
3. Results 3.1. Cracks in conventional hot-rolled plate Fig. 2 shows the photographs of the strips in 2507 SDSS fabricated by conventional hot rolling and twin-roll strip casting. Severe edge cracks were formed in the conventional hot-rolled plate (named as CHR), as presented in Fig. 2(a). The edge cracking of CHR plate firstly appeared on the fourth pass of hot rolling. The CHR with edge cutting treatment was cold rolled into the final thickness of 1 mm (named as CHR-CR) after intermediate softening and acid pickling, Fig. 2(b). The CHR-CR became very hard and brittle after the deformation of about 75%, which was difficult to continue to deform. It needed to be softened again to manufacture thinner products with a thickness below 1 mm. By contrast, no cracks were observed in the cast strip of 2507 SDSS. The cast strip was hot rolled and cold rolled into the final thickness of 0.5 mm with good surface quality, and no obvious edge cracks have been formed during rolling, Fig. 2(c)∼(e). Fig. 3 shows the micrographs of regions containing cracks in CHR in Fig. 2(a). Severely elongated structures were observed in these micrographs, with dark colored and light colored areas representing the ferrite and austenite phases, respectively. A microstructure analysis in the regions adjacent to the cracks was performed in order to study the location of the cracks with respect to ferrite and austenite. The cracks originating from the edge of the hot-rolled plate would be propagated along the interface between ferrite and austenite, Fig. 3(a). With an increase in the level of rolling deformation, the major cracks surrounded with some tiny branches of micro-cracks penetrated deeper into the steel plate, Fig. 3(b). Fig. 4(a) shows some shear localization and cracking in the conventional hot-rolled plate (CHR) of 2507 SDSS. The shear stresses appeared in duplex stainless steel due to the presence of two phases with different mechanical properties. The location of the cracks with respect to the austenite and ferrite indicated their intergranular characteristics. The hardness measurement results of different constituent phases in the CHR were marked in Fig. 4(b). The hardness of the austenite phase was similar to that of the ferrite phase. However, some particles with rich Cr and Mo were observed at the interface of the ferrite and austenite. The hardness of these particles was far higher than that of the ferrite or austenite phase. The great difference in hardness between brittle particles and the ferrite-austenite matrix was one of the important reasons
2.1.2. New processing (Strip casting) Fig. 1(b) shows the schematic diagram of the new processing (strip casting) to fabricate 2507 SDSS. A vertical type twin-roll strip caster was used. The steel was melted in an induction furnace under the N2 shielding. The steel liquid was poured into a preheated tundish to flow through a nozzle into the water-cooled casting rollers and was cast to a strip with the thickness of about 2.5 mm. The speed of strip casting was set to be 0.25 m/s. After having been homogenized at 1050 °C for 3 min, the cast strip was hot rolled from 2.5 mm to 2.0 mm with the reduction of about 20%. Subsequently, the hot rolled strip was cold rolled into the final thickness of approximately 0.5 mm with the total cold-rolling reduction of about 75% after acid pickling. The cast strip, hot-rolled strip and coldrolled strip were named as CS, CSHR and CSHR-CR, respectively. The cold-rolled strips were solid solution treated and then were cut into different shapes and sizes for microstructure, mechanical and corrosion tests, respectively. 2.2. Microstructure characterizations, mechanical properties and corrosion resistance The microstructures of experimental steels at different processing stages were observed by optical microscope (OM) under suitable magnifications. The micrographs near cracks were observed by OM and EPMA. The hardness of different constituent phases was measured by FM-700 micro-hardness Tester. Tensile specimens with gage length of 25 mm and width of 6 mm were machined parallel to the rolling direction, and tensile tests at room temperature were performed at the strain rate of 2 × 10−3s-1. The pitting corrosion resistance of the experimental steels was investigated in 3.5% NaCl solution by using potentio-dynamic measurements at room temperature with the scan rate of 0.33 mV/s. EBSD mapping was conducted on the plane (RD × ND) in the as-annealed specimens. The acceleration voltage and step size of EBSD analysis was 20 kV and 150 nm, respectively. The precipitations 249
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Fig. 6. Microstructure at different processing stages (a) CHR, (b) CHR-CR, (c) CS, (d) CSHR, and (e) CSHR-CR. RD is rolling direction and ND is normal direction.
approximately 50% austenite and 50% ferrite were observed in CHR. Both the ferrite and austenite phases were elongated along the rolling direction together with a banded structure, Fig. 6(a). Meanwhile, there was a presence of sigma phase precipitations observed at the interface between ferrite and austenite in CHR. The CHR was cold rolled with the reduction of about 75% and it exhibited a significantly refined structure, Fig. 6(b). Fig. 6(c) shows the typical microstructure of the original cast strip (CS). The ferrite matrix was decorated with different forms of austenite such as grain boundary austenite (GBA) formed at the priorferrite grain boundaries, intragranular austenite (IGA) precipitated in the ferrite grains and widmanstatten austenite (WA) grew into the ferrite grains from the GBA. The CSHR was similar to the CS in the microstructure, but its grains were refined due to the one-pass rolling deformation, Fig. 6(d). The microstructure of the cold-rolled strip fabricated by new processing (CSHR-CR) was similar to the product fabricated by conventional processing (CHR-CR). The ferrite and austenite phases were elongated along the rolling direction and the banded structure could not be completely removed even after the subsequent solution treatment, Fig. 6(e). The EBSD phase maps of the present SDSS fabricated by new processing (CSHR-CR) with different heat treatments were shown in Fig. 7. For convenience, the solid solution treatment at 1150 °C for 20 s and the treatment at 1050 °C for 150 s are named as High Temperature-Short
for the initiation of cracks during conventional hot-rolling for 2507 SDSS. Figs. 5(a) and (b) show the typical TEM micrographs in the domains adjacent to the cracks in the CHR. The corresponding chemical composition of the precipitations, ferrite and austenite was measured by EDS-TEM, as presented in Table 2. There were massive of sigma phase precipitations with rich Cr and Mo distributed along the grain boundary or phase interface. Sigma phase often formed at ferrite/austenite interfaces through a heterogeneous nucleation process, then grew toward ferrite, which was attributed to the high Cr and Mo contents and high diffusion rate in ferrite (Ramirez et al., 2003). It had a higher concentration of Cr and Mo than the ferrite or austenite phase, thus, leading to a depletion of these elements in the surrounding regions. The precipitation of sigma phase was hard and brittle, which increased mismatch of hardness between sigma phase and matrix at the phase interface, leading to the deterioration in the ductility of the material. The effects of precipitates on hot-workability would be discussed in detail later. 3.2. Microstructures at different processing stages Fig. 6 shows the optical micrographs of 2507 SDSS at different processing stages. The obvious duplex structures with a mixture of 250
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Fig. 7. Microstructure of (a)∼(c) LTLT-sample and (d)∼(f) HTST-sample obtained by EBSD.
Fig. 8. (a) Engineering stress-engineering strain curves and (b) strength and elongation of 2507 SDSS fabricated by different processing (CHR-CR and CSHR-CR) with solid solution treatment at 1050 °C for 150 s.
(yellow), and deformed (red) regions in ferrite and austenite of LTLTsample, indicating that ferrite and austenite have been almost completely recrystallized. However, there were quite a few substructures in the HTST-sample, Fig. 7(e) and (f).
Time (HTST) heat treatment and Low Temperature-Long Time (LTLT) heat treatment, respectively. The magenta and yellow green patches in the phase maps were ferrite and austenite, respectively, Figs. 7(a) and (d). The phase maps showed the microstructure consists of a bamboolike structure of coarse ferrite grains elongated along the rolling direction and equiaxed austenite grains. The ferrite grains were always coarser than the austenite grains due to the faster recrystallization and growth kinetics of ferrite than austenite regardless of the heat treatment (Kumar et al., 2017). The area of each elongated ferrite grain in the contiguous bands was converted to that of the equiaxed grain in order to facilitate the grain size measurement. The average grain size of ferrite and austenite in the LTLT-sample was measured to be about 3.55 μm and 2.56 μm, respectively. However, the grain size of the HTST-sample was effectively refined, which was measured to be about 1.96 μm for ferrite and 1.15 μm for austenite, respectively. Figs. 7(b) and (c) show the distribution of the recrystallized (blue), substructured
3.3. Mechanical properties Fig. 8(a) shows the engineering stress-engineering strain curves of 2507 SDSS fabricated by different processing. The corresponding mechanical properties, such as yield strength (YS), ultimate tensile strength (UTS) and total elongation (TE), were compared in Fig. 8(b). The YS, UTS, and TE of CSHR-CR solid solution treated at 1050 °C for 150 s have been measured to be about 675 MPa, 930 MPa and 26%, respectively, which were similar to those of CHR-CR with the same solid solution treatment. Therefore, the mechanical properties of 2507 SDSS fabricated by strip casting can satisfy the corresponding performance 251
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Fig. 9. (a) Engineering stress-engineering strain curves, (b) yield strength, (c) ultimate tensile strength and (d) total elongation of CSHR-CR with different solid solution treatments.
4. Discussion 4.1. Edge cracking prevention by strip casting For the 2507 SDSS, a high deformation resistance and low thermoplasticity due to its high contents of alloying elements (such as Cr, Mo and N) and the incoordination in deformation of two phases would lead to some difficulties in the conventional hot-rolling (Mishra et al., 2017). Severe edge cracks formed during hot-rolling due to their inherently low hot-ductility bring in the cutting-off of edges of final products, leading to an increase in cost. Kingklang and Uthaisangsuk (2017) reported that the hot-ductility of 2507 SDSS depends on many factors, such as chemical composition, deformation temperature, strain rate, phase proportions, precipitations, nature of the interphase boundaries, and softening mechanisms in the constituting phases. Thus, a true understanding of hot-workability during hot-rolling was essential for manufacturing high quality sheets of 2507 SDSS. Fig. 11(a)-(c) show the schematic illustration of cracking taking place in the conventional hot-rolling of 2507 SDSS. The process of conventional hot-rolling began with a reheating stage at about 1250 °C (Duprez et al., 2002). Though increasing the holding time can induce a degree of spheroidization of austenite, a significant fraction of austenite remained undissolved at the reheating temperature and the Widmanstätten microstructure inside the δ-ferrite grains was preserved to some extent (Iris and Suzanne, 2009). Pettersson et al. (2017) reported that the most ferrite/austenite interphase boundaries obey a K–S orientation relationship of (011)δ//(111)γ and [11 1¯] δ//[10 1¯]γ at the beginning stage of hot-rolling. Badji et al. (2008) found that the coherency of ferrite/austenite interphase can inhibit the precipitation of intermetallic compounds. The widmanstätten austenite would be transformed into a fibrous appearance along the rolling direction with increasing hot-rolling reductions. The austenite appeared more dispersedly distributed in the ferritic matrix and the coherency of ferrite/austenite interphase was destroyed, leading to a deviation from the ideal K–S orientation relationship and the formation of random interphase boundaries between the ferrite and austenite (Patra et al., 2016). Lots of precipitations at the interface between ferrite and austenite were observed during conventional hot-rolling in Figs. 4 and 5. Sigma phase precipitation was formed often at δ/γ interfaces through nucleation process. The heterogeneous nucleation of sigma phase depended on the chemical driving force and the interfacial energy. On the one hand, the
Fig. 10. Potentiodynamic polarization curves of specimens for 2507 SDSS fabricated by different processing (CHR-CR and CSHR-CR) with solid solution treatment at 1050 °C for 150 s.
index when compared with the conventional processing. Fig. 9(a) shows the engineering stress-engineering strain curves of CSHR-CR with different solid solution treatments. The corresponding mechanical properties were compared in Fig. 9(b)-(d). The YS and UTS of HTST-sample were measured to be about 750 MPa and 957 MPa, which were about 15% and 3% higher than that of the LTLT-sample, respectively. The heat treatment of HTST had a greater effect on improving YS than on improving UTS of 2507 SDSS. However, the HTSTsample exhibited the TE of about 22%, slightly lower than that of the LTLT-sample. 3.4. Pitting corrosion resistance Fig. 10 shows the potentiodynamic polarization curves of 2507 SDSS fabricated by different processing. The test was carried out in a 3.5% NaCl solution at room temperature. The curve of the CHR-CR solid solution treated at 1050 °C for 150 s nearly coincided with CSHR-CR, indicating that CHR-CR and CSHR-CR had a similar pitting corrosion resistance. The pitting corrosion potential (Epit) of 2507 SDSS in 3.5% NaCl solution was measured to be about 1.0 V. 252
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Fig. 11. Schematic illustration of the microstructure evolution and cracking damage taking place in the hot-rolled processing of 2507 SDSS (a)∼(c) conventional hotrolling, (d) and (e) one-pass hot-rolling after strip casting.
microstructure evolution taking place in the one-pass hot-rolling after strip casting of 2507 SDSS. The hot-rolling compression ratio (h0/hf, h0 and hf is the initial and final thicknesses of hot bands, respectively) of this new processing was measured to be about 1.25, which was far lower than that of conventional hot-rolling (about 20.00). The Widmanstätten microstructure inside the δ-ferrite grains can be preserved after the one-pass hot deformation and the coherency of ferrite/austenite interphase was not destroyed, which was beneficial to inhibiting the precipitation of intermetallic compounds. Therefore, the edge cracking during hot-rolling was completely eliminated by a small reduction of hot deformation in the new processing.
hot-rolling deformation forced the austenite phase to stretch along the rolling direction and increased the length of the interphase between ferrite and austenite, which provided more effective nucleation positions of brittle phase precipitations (Martin et al., 2012). Meanwhile, the massive deformation energy storage during hot-rolling provided the driving force for the nucleation of precipitates (Yutaka and Hiroyuki, 1999). On the other hand, the hot rolling deformation destroyed the coherency of ferrite/austenite interface and made the orientation between ferrite and austenite deviate from the K-S relationship due to the combined effect of strain and softening mechanisms. The deviation from the K-S relationship leaded to an increase in the interface energy, which decreased the activation energy barrier for sigma phase precipitation and accelerated the precipitation of brittle intermetallic compounds (Haghdadi et al., 2017). The edge of the upper and lower surfaces of the slab needed to bear more deformation during the conventional hot-rolling, which increased the additional tensile stress on the surface and promoted the formation of edge or surface microcracks. Metallographic characterization revealed that damage preferentially nucleated near brittle phase particles at the austenite/ferrite interface, as presented in Fig. 4. These precipitations destroyed the continuity of microstructure and provided more original source of microvoids. Meanwhile, they weakened the bonding strength of the δ/γ interface and deteriorated its plastic deformation ability. The interface between matrix and brittle phase precipitation was easily separated under the action of stress, thus, the microviods formed at the interface. Moreover, the accumulation of chromium at the δ/γ interface resulted in the chromium solid solution hardening of the interface, leading to the stress concentration and microvoids nucleation at the δ/γ interface during hot-rolling deformation. The microvoids coalesced into cracks along the ferrite/austenite interface with increasing hot deformation. Therefore, the crack originated from the edge of the upper and lower surfaces of the slab and extended to the center until it ran through the whole thickness direction. The precipitates of brittle intermetallic compounds at the ferrite/austenite interface provided numbers of potential void nucleation sites and the microvoids coalesced into cracks in the subsequent deformation. The cracks originating from edge of slab expanded along the ferrite/austenite interface and caused the more serious cracking. Fig. 11(d) and (e) show the schematic illustration of the
4.2. Effect of the HTST treatment on improving strength The YS and UTS of HTST-sample were about 15% and 3% higher than that of the LTLT-sample, respectively, as presented in Fig. 9. The strength depended strongly on grain size. The relationship of strength and grain size can be described by the Hall-Petch relation, Eq. (1). σy = σ0 + kd−0.5
(1)
Where σy is the yield stress, σ0 is the lattice friction stress, k is a constant and d is the average grain size (Zhang et al., 2012). According to the Hall-Petch relation, the strength of steel was improved by refining its grains. EBSD analysis of the experimental specimens in Fig. 7 showed that the fraction of substructures increased when adopting the HTST treatment, indicating that the enhancement in the mechanical properties was also caused by increasing dislocation density (correlated to the increasing substructures). When the dislocations moved to the grain boundaries, the barrier of grain boundaries must be overcome before the deformation can be transferred from one grain to another. The smaller the grains, the more grain boundaries and the more significant hindrance to dislocation movement. Therefore, the HTST treatment refined the grains of 2507 SDSS and reserved part of substructures, resulting in an improvement in its strength compared with the LTLT treatment. The EBSD analysis presented that the degree of recrystallization for the HTST sample was lower than that for the LTLT sample, causing the internal stress not to be completely eliminated. The residual internal stress in substructures or deformed-structures leaded 253
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to a little decrease in the plasticity of HTST specimen. Therefore, the HTST sample exhibited brittle fracture with only very reduced localization before fracture in comparison to the stress-strain curve of the LTLT sample.
corrosion and mechanical properties of UNS S31803 duplex stainless steel. Corros. Sci. 51, 2969–2975. Duprez, L., Cooman, B.C., Akdut, N., 2002. Deformation behavior of duplex stainless steel during industrial hot rolling. Steel Res. 73, 531–538. Haghdadi, N., Abou-Ras, D., Cizek, P., Hodgson, P.D., Rollett, A.D., Beladi, H., 2017. Austenite-ferrite interface crystallography dependence of sigma phase precipitation using the five-parameter characterization approach. Mater. Lett. 196, 264–268. Iris, A.-A., Suzanne, D.-M., 2009. Great Britain and the United States. Duplex Stainless Steels. ISTE Ltd and John Wiley & Sons, Inc., pp. 31–39. Kingklang, S., Uthaisangsuk, V., 2017. Investigation of hot deformation behavior of duplex stainless steel grade 2507. Metall. Mater. Trans. A 48, 95–108. Kumar, A., Khatirkar, R.K., Chalapathi, D., Kumar, G., Suwas, S., 2017. Microstructure and texture development during cold rolling in UNS S32205 and UNS S32760 duplex stainless steels. Metall. Mater. Trans. A 48, 2349–2362. Martin, G., Yerra, S.K., Bre´chet, Y., Ve´ron, M., Mithieux, J.-D., Che´hab, B., Delannay, L., Pardoen, T., 2012. A macro- and micromechanics investigation of hot cracking in duplex steels. Acta Mater. 60, 4646–4660. Martins, M., Casteletti, L.C., 2009. Microstructural characteristics and corrosion behavior of a super duplex stainless steel casting. Mater. Charact. 60, 150–155. Mishra, M.K., Balasundar, I., Rao, A.G., Kashyap, B.P., Prabhu, N., 2017. On the high temperature deformation behavior of 2507 super duplex stainless steel. J Mater. Eng. Per. 26, 802–812. Pardal, J.M., Tavares, S.S.M., Fonseca, M.C., Souza, J.A., Corte, R.R.A., Abreu, H.F.G., 2009. Influence of the grain size on deleterious phase precipitation in super duplex stainless steel UNS S32750. Mater. Charact. 60, 165–172. Patra, S., Ghosh, A., Kumar, V., Chakrabarti, D., Singhal, L.K., 2016. Deformation induced austenite formation in as-cast 2101 duplex stainless steel and its effect on hot-ductility. Mater. Sci. Eng. A 660, 61–70. Pettersson, N., Wessman, S., Hertzman, S., Studer, A., 2017. High-temperature phase equilibria of duplex stainless steels assessed with a novel In-situ neutron scattering approach. Metall. Mater. Trans. A 48, 1562–1571. Ramirez, A.J., Lipold, J.C., Brandi, S.D., 2003. The relationship between chromium nitride and secondary austenite precipitation in duplex stainless steels. Metall. Mater. Trans. A 34, 1575–1597. Ramkumar, K.D., Thiruvengatam, G., Sudharsan, S.P., Mishra, D., Arivazhagan, N., Sridhar, R., 2014. Characterization of weld strength and impact toughness in the multi-pass welding of super-duplex stainless steel UNS 32750. Mater. Des. 60, 125–135. Santos, C.A., Spim Jr, J.A., Garcia, A., 2000. Modeling of solidification in twin-roll strip casting. J. Mater. Process. Technol. 102, 33–39. Santos, D.C.D., Magnabosco, R., 2016. Kinetic study to predict sigma phase formation in duplex stainless steels. Metall. Mater. Trans. A 47, 1554–1565. Villanueva, D.M.E., Junior, F.C.P., Plaut, R.L., Padilha, A.F., 2006. Comparative study on sigma phase precipitation of three types of stainless steels: austenitic, super ferritic and duplex. Mater. Sci. Technol. 22, 1098–1104. Wang, S.H., Liu, Z.Y., Zhang, W.N., Wang, G.D., Liu, J.L., Liang, G.F., 2009. Microstructure and mechanical property of strip in Fe-23Mn-3Si-3Al TWIP steel by twin roll casting. ISIJ Int. 49, 1340–1346. Wang, Y.Q., Han, J., Wu, H.C., Yang, B., Wang, X.T., 2013. Effect of sigma phase precipitation on the mechanical and wear properties of Z3CN20.09M cast duplex stainless steel. Nuc. Eng. Des. 259, 1–7. Yutaka, S.S., Hiroyuki, K., 1999. Preferential precipitation site of sigma phase in duplex stainless steel weld metal. Scripta Mater. 40, 659–663. Zhang, X.D., Hansen, N., Gao, Y.K., Huang, X.X., 2012. Hall-Petch and dislocation strengthening in graded nanostructured steel. Acta Mater. 60, 5933–5943.
5. Conclusions In the present work, the edge cracking mechanism of 2507 SDSS during conventional hot-rolling was discussed and the high quality thin strips were manufactured by strip casting. The mechanical properties and corrosion resistance of 2507 SDSS fabricated by these two different processes were investigated. The main conclusions are obtained as follows. (1) The brittle intermetallic compounds precipitated at the phase interface in 2507 SDSS during conventional hot-rolling provided numberous potential void nucleation sites, which accelerated the propagation of cracks along the ferrite/austenite interface and caused serious cracking. (2) The edge cracking was completely avoided during one-pass hotrolling after strip casting through a small reduction of hot deformation. Meanwhile, the mechanical properties and corrosion resistance of 2507 SDSS fabricated by this new technology can reach the standard of conventional production. (3) The High Temperature-Short Time (HTST) heat treatment significantly improved the strength of 2507 SDSS, especially improving its yield strength. Acknowledgments This work was supported by the National Natural Science Foundation of China with the contracts of U1460204, U1660117 together with Baosteel Co., and National Natural Science Foundation of China (51774083). References Badji, R., Bouabdallah, M., Bacroix, B., Kahloun, C., Bettahar, K., Kherrouba, N., 2008. Effect of solution treatment temperature on the precipitation kinetic of σ phase in 2205 duplex stainless steel welds. Mater. Sci. Eng. A 496, 447–454. Cabrera, J.M., Mateo, A., Llanes, L., Prado, J.M., Anglada, M., 2003. Hot deformation of duplex stainless steels. J. Mater. Pro. Tech. 143, 321–325. Chail, G.C., Kangas, P., 2016. Super and hyper duplex stainless steels: structures, properties and applications. Procedia Struct. Integr. 2, 1755–1762. Deng, B., Wang, Z., Jiang, Y., Sun, T., Xu, J., Li, J., 2009. Effect of thermal cycles on the
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