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Vol. 9. Issue 1.
Pages 124-132 (January - February 2020)
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Vol. 9. Issue 1.
Pages 124-132 (January - February 2020)
Original Article
DOI: 10.1016/j.jmrt.2019.10.037
Open Access
Cold rolling performance for austenitic stainless steel with equilibrium and non-equilibrium microstructures
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Kangda Haoa, Ming Gaob,
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mgao@mail.hust.edu.cn

Corresponding author.
, Run Wua
a The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China
b Wuhan National Laboratory for Optoelectronics (WNLO), Huazhong University of Science and Technology, Wuhan 430074, China
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Table 1. Tensile properties, where TS is tensile strength, EL is elongation.
Abstract

Aiming to provide effective theoretical guidance for continuous rolling, the cold rolling with reduction of 45% was carried out for the equilibrium and non-equilibrium microstructures of austenitic stainless steel (ASS). The microstructural evolution, texture components and tensile properties were investigated. The results showed that the fraction of martensite transformed from austenite is 21.1% for the equilibrium base metal (BM), while that for the non-equilibrium weld metal (WM) achieves 32.0%. Whether employing rolling or not, the fraction of misorientation larger than 15° within the BM is much higher than that of the WM because of the large amount of twins. Moreover, the texture of the rolled BM is mainly composed of Goss component with small amount of S and Brass, while that of the WM is composed of mainly S and small amount of Brass, with formation of Copper component. The tensile strength of the rolled BM and WM reaches approximately 1300 MPa, the elongation rate of the BM decreases from 53% to 5.7%, while that of the WM decreases from 32% to 3.5%. The microstructure evolutions of the BM and WM during rolling were compared, and the microstructure-mechanical properties relation was established and discussed based on martensitic nucleation.

Keywords:
Austenitic stainless steel
Cold rolling
Laser weld
Microstructure
Texture component
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1Introduction

Rolling is one of the most key technologies for plate manufacturing, because it can eliminate defects of shrinkage and inclusion in casting process, and its plates have advantages of compact structure, fine grains and good mechanical properties [1–3]. Among which, continuous rolling by welding two sheets together is an effective method to increase capacity and market competitiveness for iron and steel enterprises [4,5]. By applying severe rolling stress on the plate, both of the base metal (BM) and the weld metal (WM) microstructures vary significantly, and the weld quality is then crucial for high-efficiency continuous rolling [6].

Due to good mechanical properties and corrosion resistance, austenitic stainless steel (ASS) is widely used in the fields of nuclear, medical equipment and chemical machinery, and its microstructural evolution and mechanical properties during rolling are the focus of the current researches [7–9]. The dislocation multiplication proceeds rapidly by rolling of the ASS, the free energy increases because of the high strain energy, which promotes the austenite transforming to martensite [10]. Zhang et al. (2018) confirmed that the α' martensite was difficult to be directly transformed from austenite, but experienced transitional phase of ε martensite. Quantitative relation model between martensite fraction and rolling reduction was finally established based on the experimental results [11]. In cold rolling of AISI 304 ASS, it was found that work hardening effect occurred accompanied with the formation of martensite, with its forming rate slowed down by increasing the rolling reduction [12]. By cold rolling of metastable ASS with maximum reduction of 50% at 0 °C and 24 °C, it was proved that the martensitic transformation was more likely to be induced by lower temperature, and the difference in hardness increased with the increase of rolling reduction [2]. It was also found that the martensite fraction within the ASS reached 31.5% by reduction of 40%, the tensile strength increased from 680 to 1444 MPa, but the tensile ductility decreased from 26.6% to 1.8% [13]. In addition, during reversion annealing of cold rolled AISI 304/316 austenitic stainless steel, three distinct stages were identified as the reversion of strain-induced martensite to austenite, the primary recrystallization of the retained austenite, and the grain growth process. Among which, the kinetics of the reversion and recrystallization processes were enhanced at higher temperatures, while the grains grew up at the end of recrystallization [14–16].

Moreover, the rolling texture evolution also attracts much attention [17,18]. There generally exist texture components represented by {110} <001> G (Goss), {110} <112> B (Brass), {123} <634> S and {112} <111> C (Copper) [19]. By rolling of AISI 316 L ASS with 90% reduction, it initially showed components C and B with a spread towards G. The C component diminished while the B and G increased with the progress of rolling, and there formed additional texture of {111} at high strain level with reduction approaching 90% [20]. By rolling of AISI 304 L metastable ASS with 90% reduction, it was also confirmed that the B component was significantly enhanced with the reduction exceeding 80% [21]. By comparing the rolling texture of AISI 304 L with low stacking fault energy and AISI 316 L with medium stacking fault energy, it was found that their texture components were similar at low deformation level, but the rate of texture C transforming to B was lower in AISI 304 L than that in AISI 316 L due to martensitic transformation at high deformation level [22].

The researches above are mainly carried out within equilibrium microstructure of the BM, but that within the non-equilibrium microstructure of the WM has not been reported yet. Generally, some high-temperature phase is remained to room temperature within the WM because of the thermal cycle of the welding, and the grains of the WM are large in size due to low heat conductivity and high thermal expansion coefficient of the ASS, which results in noteworthy difference in microstructure and texture evolutions of the BM and the WM during rolling. Thus, comparison of cold rolling performance of the BM and the WM of the ASS was carried out in this paper, and the related mechanisms were proposed, which is significant to achieve high-efficiency continuous rolling and deepen the rolling theories.

2Experimental methods

The BM used was 2.5 mm-thick AISI 304 ASS (Fe-18.11Cr-7.01Ni-1.21Mn-0.04C, wt. %). As shown in Fig. 1, the laser welding experiment was carried out with a Fanuc M-710 robot and an IPG YLS-6000 fiber laser, under welding parameters of laser power 3 kW, welding speed 2.5 m/min and defocus distance of 10 mm.

Fig. 1.

Experimental set-up used in this study.

(0.2MB).

The sample was 200 mm in length and 100 mm in width, and was rolled by a two-high rolling mill at room temperature. The rolling reduction was 45%, with interval of 15% by three passes. The cross-sectioned samples were prepared to conduct metallography, electron backscatter diffraction (EBSD) and tensile tests by wire-electrode cutting. Among which, the metallurgical samples were chemically etched by a solution containing cupric chloride 5 g, hydrochloric acid 5 mL, ethanol 50 mL and distilled water 50 mL, and were observed by FEI Quanta-200 environmental scanning electron microscope (SEM). The EBSD samples were firstly grinded by sandpaper, and then electrolytically polished by a solution containing perchloric acid 10 mL and ethanol 90 mL under voltage of 25 V for 5 min. The crystal orientation data was collected employing EDAX-TSL OIM-EBSD system equipped on Sirion 200 field emission SEM, with step size of 0.10 μm. The raw date was then cleaned with tolerance of 2 steps for crystal analysis and texture calculation by harmonic expansion. Moreover, the tensile properties were tested according to ASTM E8/(E8M)-11, the result was the average of three samples.

3Results and discussions3.1Microstructure characteristics

As shown in Fig. 2 a&b, the BM is composed of uniform austenite with a large amount of twins, while the WM is in-homogeneously mixed with austenite and δ ferrite. As shown in Fig. 2 c&d, the austenite within the BM is partially transformed to martensite by rolling, the other part is just compressed without martensitic transformation, and the transition and non-transition zone are uniformly interlaced. Relatively, the rolled WM is presented as fibrous structure with enlarged non-transition zone. It can be noted that no crack was found within the weld, indicating the continuous rolling of AISI 304 ASS sheets can be achieved.

Fig. 2.

Microstructure characteristics, (a) base metal (BM) before rolling, (b) weld metal (WM) before rolling, (c) BM after rolling with 45% reduction, (d) WM after rolling with 45% reduction.

(1.12MB).

Generally, the austenite within the AISI 304 ASS is in metastable state, which was reported by Olson [23] and Lee et al. [24]. In this case, the thermodynamic driving force can be supplemented by the rolling stress, which promotes the austenite transforming to martensite. Generally, the martensitic transformation is nucleation controlled, and occurs easily at high energy regions with high dislocation densities. In this study, the twins within the BM are beneficial to release the stress, while the δ ferrite within the WM is broken during rolling, which causes stress concentration by inhibiting dislocation motion. Thus, the microstructural evolutions of the BM and WM during rolling can be illustrated as follows.

The dislocation multiplication develops rapidly during rolling. As shown in Fig. 3a, the stress concentration at the grain boundaries with twins is weakened, and is incapable to induce martensitic transformation, because the dislocations can pass through the boundaries and move along the twins. Then, the austenite is just elongated along the rolling direction. However, the stress concentration is high enough to promote the austenite transforming to martensite at boundaries without twins. Moreover, the transition and non-transition zone are uniformly interlaced because of equilibrium microstructure of the BM.

Fig. 3.

Schematic diagram of microstructure evolution during rolling, (a) BM, (b) WM, where A is austenite, M is martensite, δ-F is δ ferrite.

(1.12MB).

As shown in Fig. 3b, the grain size of the WM is obviously larger than that of the BM, owing to the non-equilibrium heating process of welding and the thermal inertia of the austenite growth. The columnar structure grows from the weld edge to the center because of the temperature gradient, with some dendritic δ ferrite remained and distributed at grain boundaries of the austenite. The martensitic transformation is then induced to a larger extent. On one side, as mentioned above, the transformation occurs easily at the grain boundaries under stress concentration. On the other hand, the δ ferrite is broken during rolling, which acts as barrier for the dislocation motion and promotes the formation of high-energy regions. The number of high-energy nucleation sites for martensitic transformation then increases. However, there also exists austenite not transformed, but with larger size elongated along the rolling direction, and the microstructure uniformity of the WM after rolling is lower than that of the BM, because of their original microstructure characteristics.

3.2EBSD analysis

As shown in Fig. 4, the initial microstructure of the BM and the WM are all composed of austenite, while the δ ferrite is difficult to be distinguished because of the identification accuracy. The martensite fraction within the BM and the WM after rolling are 21.1% and 32.0%, respectively, indicating more sufficient transformation occurred in the WM.

Fig. 4.

Inverse pole figures (IPF) and phase figures, (a) BM before rolling, (b) WM before rolling, (c) BM after rolling with 45% reduction, (d) WM after rolling with 45% reduction.

(2.48MB).

Generally, there are shear directions of [11]γ, [11] γ, [11] γ on {111}γ plane, with the angle of 60°. Thus, the grain misorientation mainly locates at 60° for the BM and WM, as shown in Fig. 5. Among which, there are more boundaries with 60° misorientation for the BM because of the twins, with the fraction of high-angle misorientation larger than 15° (FHM) achieving 94.7%. However, the twins disappear during welding, and the grains grow up, resulting in the FHM decreases to 61.8%. By rolling with 45% reduction, the grains are broken, and a part of austenite transforms to martensite, the FHM thus decrease to 54.5% and 41.1% for the BM and the WM, respectively. It can be noted that the FHM is always lower in the WM than that in the BM, whether employing rolling or not.

Fig. 5.

Misorientation distribution characteristics, (a) the overall misorientation distribution, (b) the fraction of misorientation above 15°.

(0.15MB).

Moreover, the kernel average misorientaion (KAM) maps are shown in Fig. 6, which is related to the deformation ability. The KAM value increases from 0.489° to 1.587° for the BM after rolling, while increases from 0.763° to 1.772° for the WM. Whether employing rolling or not, the KAM value of the WM is larger than the BM, indicating poorer deformation ability of the WM.

Fig. 6.

Kernel average misorientaion (KAM) maps, (a) BM before rolling, (b) WM before rolling, (c) BM after rolling with 45% reduction, (d) WM after rolling with 45% reduction.

(2.86MB).

Generally, texture can be formed by directional solidification and rolling deformation. For simplification, only the representative sections φ2 = 0, 45, and 65° are chosen to reveal the texture evolution, as shown in Fig. 7. The austenite within the BM shows random distribution, with maximum intensity of just 4.842. Because of the directional solidification of the welding, the austenite in the WM develops {110} <110> components with intensity of 11.088. By rolling with 45% reduction, the BM develops strong G and S components, and the B component also forms because of the contribution of the twins [19], with their intensity of 20.508, 8.126 and 6.208, respectively. However, there is no G component in the rolled WM, the S component is more likely to form with intensity of 30.656, and the intensity of B component increases to 11.722. Moreover, new component of C is observed with intensity of 12.176. Souza Filho et al. [19] pointed out that the B orientation is beneficial for the martensitic transformation because of the Nishiyama-Wassermann crystallographic orientation relationship, which promotes the formation of martensite.

Fig. 7.

Orientation distribution function (ODF), (a) BM before rolling, (b) WM before rolling, (c) BM after rolling with 45% reduction, (d) WM after rolling with 45% reduction.

(0.78MB).
3.3Tensile properties

As shown in Table 1, the tensile strength of the WM is similar to that of the BM, while its elongation rate is 60% of that of the BM. By rolling with 45% reduction, the strength of the BM and the WM increases to approximately 1300 MPa, while the elongation rate decreases to 5.7% and 3.5%, respectively.

Table 1.

Tensile properties, where TS is tensile strength, EL is elongation.

Weld no.  TS (MPa)  EL (%) 
BM  675 ± 4  53 ± 0.5 
WM  667 ± 6  32 ± 0.4 
BM-45%  1262 ± 9  5.7 ± 0.4 
WM-45%  1311 ± 5  3.5 ± 0.3 

As shown in Fig. 8, the fracture surfaces of the BM and the WM before rolling are presented as dimple because of the microstructure of pure austenite. By rolling with 45% reduction, the martensite is induced by stress, which is difficult to coordinate with the deformation under dislocation force. The martensite is then separated from the martrix or cracks itself during deformation. Moreover, the existence of the δ ferrite within the WM also inhibits the moving of the dislocations and the propagation of the cleavage cracks. Then, cleavage steps are formed because of the change of the crack propagation direction. Thus, the cleavage characteristic within the rolled WM is more obvious than that within the rolled BM.

Fig. 8.

Fracture surfaces, (a) the initial BM, (b) the initial WM, (c) the rolled BM, (d) the rolled WM.

(1.22MB).
4Conclusions

  • (1)

    The continuous rolling of AISI 304 ASS sheets can be achieved by laser welding. Among which, the fraction of stress induced martensite within the rolled weld metal (WM) is much higher than that within the rolled base metal (BM), because the large amount of twins within the BM is beneficial to promote the dislocation motion and decreases the stress concentration, while the broken δ ferrite within the WM inhibits the dislocation motion and causes stress concentration.

  • (2)

    The fraction of misorientation larger than 15° for the BM decreases from 94.7% to 54.5%, while that for the WM decreases from 61.8% to 41.1%. Whether employing rolling or not, it is much higher within the BM than that within the WM.

  • (3)

    The grains of the BM are random distributed before rolling, but {110} <110> component is developed within the WM because of the directional solidification of the welding. By rolling with reduction of 45%, the BM shows Goss component with certain S and Brass components, while the WM develops mainly S component and small amount of Brass, accompanied with the formation of Copper component.

  • (4)

    The weld was obviously work hardened by rolling: the tensile strength (TS) of the BM increases from 675 MPa to 1262 MPa with its elongation rate (EL) decreases from 53% to 5.7%, the TS of the WM increases from 667 MPa to 1311 MPa with its EL decreases from 32% to 3.5%. Relatively, the hardening effect of the WM is stronger than that of the BM.

Conflict of interest

We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled, ‘Cold rolling performance for austenitic stainless steel with equilibrium and non-equilibrium microstructures’.

Acknowledgments

This research is financially supported by the National Natural Science Foundation of China (grant nos. 51775206, 51905391 and 51805182).

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