Influence of the microstructure on the plastic behaviour of duplex stainless steels

Influence of the microstructure on the plastic behaviour of duplex stainless steels

Materials Science and Engineering A 528 (2011) 2259–2264 Contents lists available at ScienceDirect Materials Science and Engineering A journal homep...

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Materials Science and Engineering A 528 (2011) 2259–2264

Contents lists available at ScienceDirect

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

Influence of the microstructure on the plastic behaviour of duplex stainless steels Alberto Moreira Jorge Júnior a,∗ , Gedeon Silva Reis b , Oscar Balancin a a b

Department of Materials Engineering, Federal University of São Carlos, Rod. Washington Luís, km 235, CEP 13565-905 São Carlos, SP, Brazil Federal Institute for Education, Science and Technology - IFMA, Av. Getúlio Vargas, 04, CEP 65025-001 São Luis, MA, Brazil

a r t i c l e

i n f o

Article history: Received 26 July 2010 Received in revised form 23 October 2010 Accepted 29 November 2010 Available online 5 December 2010 Keywords: Duplex stainless steels Microstructures Plastic flow behaviour Hot torsion tests

a b s t r a c t Microstructure change of ␣ (ferrite) + ␥ (austenite) two-phase structure in duplex stainless steels deformed by hot torsion tests is briefly analyzed. Two types of stainless steels containing different volume fractions of ferrite and austenite were torsion deformed at temperatures ranging from 900 to 1250 ◦ C. Steel A (25.5Cr–4.9Ni–1.6Mo) contained Creq /Nieq = 4.8 and steel B (22.2Cr–5.6Ni–3Mo) contained Creq /Nieq = 3.5 bring about different microstructures and flow stress behaviour. The results show that the shape of the flow stress curves depends on the material and on deformation conditions. Four different flow curve shapes were observed. At high temperatures, steel A displayed a plastic behaviour typical of ferritic stainless steels. As the deformation temperature decreased, the flow curves presented peak stresses at low-temperature deformation. When the austenite particles are distributed coarsely in the matrix (steel B), the plastic flow curve displays a stress peak separating extensive regions of hardening and softening. When both phases have the same volume fractions, the microstructure is characterized by percolation of the two phases in the samples, and the plastic flow curve takes on a very distinctive shape in hot torsion tests. The role of the microstructure present during deformation on the shape of the flow stress curves is analyzed. © 2010 Elsevier B.V. Open access under the Elsevier OA license.

1. Introduction In the past decades, duplex stainless steels, which have a twophase ferrite (␣) + austenite (␥) microstructures, have attracted great interest with its cost-saving combination of high strength and improved resistance to general and localized corrosion, stresscorrosion cracking, abrasion and wear [1–5]. Their properties have been continually improved mainly by optimization of alloying, hot and cold working and heat treatment [6]. The formation of two phases is determined by the preferential partitioning of the alloying elements that make up the ferritic and austenitic phases, mainly the chromium, nickel and molybdenum contents. Numerous microstructural changes can occur as a result of solute partitioning during thermal and thermomechanical treatments of these steels, with the transformation of ferrite into austenite being the main structural change that occurs during hot processing. This transformation occurs through the nucleation and growth of austenite particles. These particles, which are shaped like Widmanstätten-type plates [6,7], follow the orientation relationship of Kurdjumov–Sachs [8,9] and are hence coherent with the ferritic matrix.

∗ Corresponding author. Tel.: +55 16 33518531; fax: +55 16 33615404. E-mail address: [email protected] (A.M. Jorge Júnior). 0921-5093 © 2010 Elsevier B.V. Open access under the Elsevier OA license. doi:10.1016/j.msea.2010.11.087

Duplex stainless steels require special care in metallurgical processing due to the characteristics of the phases that are present during hot deformation, the mechanisms of hardening and softening act differently in each of the phases [10–12]. Ferrite is softened significantly by recovery upon deformation, rapidly forming a subgrain substructure [13]. Austenite has a larger hardening region with greater accumulation of internal energy, leading the material to dynamic recrystallization [14]. When the two phases are deformed jointly, the deformation is not distributed uniformly [15]. Initially, deformation concentrates in the ferrite, which is the softer phase. As deformation proceeds, the internal deformation gradients decline due to the transfer of stresses and strains from the matrix to the austenite and to the action of softening mechanisms, such as dynamic recovery and recrystallization and grain boundary sliding. As a consequence of the complex plastic behaviour presented by biphasic materials, different shapes of the plastic flow curve have been found during the hot deformation of duplex stainless steels [9,16–18], including curves of materials that recover dynamically, curves with peak stresses at low deformations, or curves with a threshold stress after hardening, followed by continuous softening. The purpose of this work is to evaluate the plastic behaviour of duplex stainless steels during hot deformation based on hot torsion tests, in a systematic study that correlates the microstructure with the shape of the plastic flow curve.

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Table 1 Variations in the volumetric fraction (in %) of austenite in steels A and B as a function of reheating temperature.

Steel A Steel B

900 ◦ C

950 ◦ C

1000 ◦ C

1050 ◦ C

1100 ◦ C

1150 ◦ C

1200 ◦ C

1250 ◦ C

38.8

19.9

14.3

9.4

5.2 53.3

2.4 46.5

1.3 41.4

<1.0 39.6

2. Materials and methods In this work, two types of duplex stainless steels were used having different equivalent Cr/Ni ratios. Steel A (25.5Cr–4.9Ni–1.6Mo) with Creq /Nieq = 4.8, contained a volumetric fraction of austenite of about 40% at ambient temperature, after conventional annealing, while steel B (22.2Cr–5.6Ni–3Mo) with Creq /Nieq = 3.5, contained 50% of austenite in the same conditions. Fig. 1 shows a pseudobinary diagram of the Fe–Cr–Ni system containing 70% of iron, and indicates the constituent phases expected for these steels after reaching equilibrium during reheating at high temperatures. The thermomechanical tests were carried out on a computercontrolled hot torsion-testing machine. In an initial heat treatment, some samples, which had a diameter of 6 mm and an effective length of 20 mm, were reheated to temperatures ranging from 900 to 1250 ◦ C in an infrared radiation furnace coupled to the machine, were held at these temperatures for 1 h, and water quenched. The torsion tests themselves were performed by heating the samples to temperatures ranging from 1000 ◦ C to 1250 ◦ C for 600 s, followed by cooling to the test temperature, which varied from 1250 ◦ C to 900 ◦ C and a strain rate of 1 s−1 . To observe the microstructural evolution occurring during the thermomechanical treatment, water was injected through the quartz tube that surrounded the samples immediately after the end of the soaking time or after deformation. After the standard metallographic procedures, the samples were subjected to attack with a solution composed of 100 ml of distilled water, 20 ml of HCl and 2 g of potassium metabisulfite, and observed in an optical microscope Carl Zeiss – AXIO. Electron Backscattered Diffraction (EBSD) analyses were performed in a Philips XL-30FEG scanning electron microscope

Fig. 1. Pseudobinary diagram of Fe–Cr–Ni containing 70% of iron. The relative location of steels A and B was determined considering the equivalent Cr/Ni ratios. Creq = %Cr + 1.37%Mo + 1.5%Si + 2%Nb + 3%Ti. Nieq = %Ni + 22%C + 14.2%N + 0.31%Mn + %Cu.

equipped with a TSL OIM MSC-2200 equipment. Transmission Electron Microscopy (MET) was carried out by using a Philips CM-120 microscope. Thin foils for TEM observation were prepared by ion milling. 3. Results and discussion Initially, the microstructures of the samples reheated and soaked for 1 h were examined. Table 1 shows the measured values of the proportions of austenite in the ferritic matrix as a function of reheating temperature. Although the ferrite/austenite ratio varied with temperature, steel B always contained a higher volumetric fraction of austenite than steel A in similar reheating conditions. It was found that in steel A, dissolution already started to occur at temperatures slightly above 900 ◦ C, while in steel B, the variation of the volumetric fraction of austenite upon reheating at temperatures below 1100 ◦ C was very small. During torsion tests, phase transformation occurs during cooling from the reheating temperature to the testing temperature. This transformation takes place with nucleation and growth of the austenite phase at grain boundaries or inside ferrite grains, as will be shown below. The plastic behaviour of the steels was evaluated by determining the evolution of stress with strain in the hot torsion tests. Both the stress level and the shape of the plastic flow curve depend on each of the materials and on the deformation conditions. In the tests conducted here, the stress–strain curves assumed one of the four shapes depicted in Fig. 2. Curve 1 corresponds to tests carried out at 1200 ◦ C in steel A, after reheating at the same temperature. Curve 2 was observed in the same material in the test performed at 900 ◦ C after reheating at 1200 ◦ C. Curves 3 and 4 were observed in steel B. Curve 3 was obtained in a test carried out at 1250 ◦ C, after reheating at 1250 ◦ C and Curve 4 resulted from a test performed at a lower temperature; straining it at 1100 ◦ C after reheating at same temperature. Each of the shapes of the plastic flow curve observed in this study is analyzed below together with microstructure evolution.

Fig. 2. Typical plastic flow curves obtained in the tests: (1) sample with an essentially ferritic microstructure; (2) sample with a ferritic matrix and finely dispersed coherent austenite particles; (3) ferritic matrix with coarsely dispersed austenite particles; and (4) duplex microstructure.

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Fig. 3. Optical microstructures of the steel A deformed to fracture associated with the Curve 1 from Fig. 2. (a) Essentially ferritic initial microstructure. (b) Final microstructure, showing old highly deformed grains together with the formation of new grain/subgrain microstructure.

3.1. Essentially ferritic microstructure Analyzing Curve 1 from Fig. 2, the applied stress increases rapidly with the imposed deformation, establishing a steady state after a given amount of deformation. In this condition of deformation, duplex stainless steels with higher equivalent chromium nickel ratios, such as steel A, have a microstructure consisting of a ferritic matrix with a very small volume fraction of austenite particles. Upon reheating at high temperatures, the austenite loses its stability and the microstructure becomes essentially ferritic (Fig. 3a). When ferrite is deformed, the process of softening by dynamic recovery that takes place and simultaneously establishes equilibrium between the generation and elimination of dislocations, after a given amount of deformation. Continuing deformation leads to the formation of new grains through rotation and a gradual increase in the disorientation of the subgrains, in a process known as extended dynamic recovery or continuous dynamic recrystallization [17]. Fig. 3b corresponds to the final deformation state where one can see old highly deformed grains together with the formation of new grain/subgrain microstructure. 3.2. Ferritic matrix with coherent austenite particles Curve 2, from Fig. 2, was observed in the test performed at 900 ◦ C after reheating steel A at 1200 ◦ C. Fig. 4a shows the initial state of this condition. From this picture it can be observed that the cooling

to lower temperatures led to nucleation of Widmanstätten-shaped austenite particles. It has been shown that in duplex steels, the austenite follows the Kurdjumov–Sachs relationship with respect to ferrite. For example, Arboledas et al. [8] found that the (0 1 1)␣ planes are parallel to the (1 1 1)␥ planes, and, consequently, that the [1 0 0]␣ and [1 1 0]␥ directions are also parallel. Thus, the austenitic phase formed by transformation of solid-state ferrite is coherent with the matrix. The Pole Figures (PF), obtained by means EBSD, corresponding to the initial condition prior deformation, are shown in Fig. 5a. From this picture it is clear the parallelism between the above-mentioned planes, confirming the coherence between the transformed austenitic phase and the ferritic matrix. This coherence increases the degree of interaction among precipitates and dislocations. In the initial stage of deformation, the austenite particles inhibit deformation of the matrix, increasing the initial flow stress and the hardening rate. This effect prevails while the austenite particles remain coherent with the matrix. Upon deformation the austenitic particles lost their coherence and aligned in the deformation direction. As deformation proceeds, the process of extended dynamic recovery of the ferritic matrix begins. Initially, subgrains are formed, which are rotated in relation to the austenite particles, losing the orientation relationship. This orientation loss can be observed in Fig. 5b that corresponds to a true deformation of 0.35 and, from where, one can see that the Kurdjumov–Sachs relationship is already no more obeyed. Thus,

Fig. 4. Optical microstructures of the steel A deformed to fracture associated with the Curve 2 from Fig. 2. (a) Initial microstructure showing Widmanstätten-shaped austenite particles. (b) Final microstructure, showing strained austenite particles aligned in the deformation direction and the formation of a significant number of new ferritic grains.

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Fig. 5. Pole Figures obtained by EBSD, indicating the Kurdjumov–Sachs relationship with respect to (0 1 1)␣//(1 1 1)␥ and (1 1 1)␣//(1 1 0)␥ planes. (a) Corresponding to the initial undeformed microstructure with precipitated particles of Widmanstätten-shaped austenite, showing the coherence between austenite and ferrite phases (black arrowhead lines). (b) Corresponding to the true strain of 0.35 showing a pattern with a high degree of misorientation between austenite and ferrite phases.

the process of softening that decreases the flow stress toward a steady state begins before the formation of a significant number of new grains [15]. After softening is complete, the austenite particles are no longer coherent with the matrix. The evolution of stress with deformation takes on the shape of curve 2 and the final microstructure is presented in Fig. 4b.

3.3. Ferritic matrix with coarsely dispersed austenite particles Curve 3, from Fig. 2, was obtained in a test carried out at 1250 ◦ C, after reheating steel B at 1250 ◦ C. Fig. 6a shows the initial state of this condition. Duplex stainless steels such as steel B, with lower equivalent chromium nickel ratios, have equal proportions of phases in the conditions of use. Thus, the two phases are aligned in the rolling direction of the previous process. Upon reheating at high temperatures, part of the austenite dissolves. For example, when steel B was reheated to 1250 ◦ C, the volumetric fraction of austenite was about 40% (see Table 1). The austenite particles no longer percolated throughout the microstructure, but appeared as particles coarsely dispersed in the ferritic matrix, as can be observed in Fig. 6a.

Fig. 7 shows TEM bright field images of steel B deformed in the conditions of Curve 3, where Fig. 7a corresponds to a true strain of 0.35 and the true straining of 0.55 is shown in Fig. 7b. Typical Selected Area Electron Diffraction Patterns (SAEDP) for ferrite and austenite, obtained in both conditions, are shown Fig. 7a and b, respectively. In duplex stainless steels whose matrix is constituted of ferritic phase, the harder austenitic phase offers greater resistance to deformation than the matrix. Consequently, during the first stages of plastic deformation, as can be seen in Fig. 7a, ferritic phase accommodates most of the strain, causing hardening of the material to be initially controlled by the efficient dynamic recovery of the ferrite, which leads to a constant plastic flow stress (Curve 3). At greater deformations, as can be seen in Fig. 7b, there is a transfer of load from ferrite to austenite, leading to an increase in hardening of the material, which appears in Curve 3 in the form of linear hardening. As a result of the concentration of plastic deformation in the ferrite, Fig. 6b indicates that the matrix flowed around or parallel to the harder austenite particles. Since the austenite particles no percolated throughout the samples, after alignment of the microstructure along the deformation direc-

Fig. 6. Optical microstructures of the steel B deformed to fracture associated with the Curve 3 from Fig. 2. (a) Initial microstructure showing austenitic particles coarsely dispersed in the ferritic matrix. (b) Final microstructure, showing austenite particles aligned along the deformation direction.

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Fig. 7. TEM bright field images and SAEDPs of steel B deformed in the conditions of Curve 3. (a) Corresponding to a true strain of 0.35, showing ferritic phase accommodating most of the strain; and the SAEDP of ferrite. (b) Corresponding to a true strain of 0.55, showing that there is a transfer of load from ferrite to austenite; and the SAEDP of austenite.

tion, straining again occurred preferentially in the matrix and the stress level decreased. In this case, a peak separating an extensive hardening region from an extensive softening region, as shown in Curve 3 of Fig. 2, characterized the flow curve.

3.4. Duplex microstructure Curve 4, from Fig. 2, resulted from a test performed in steel B at 1100 ◦ C. Fig. 8 shows microstructures in different deformation conditions, i.e., in (a) initial condition, and in the following true

Fig. 8. Optical microstructures of steel B deformed to fracture associated with the Curve 4 from Fig. 2, showing microstructures in different deformation conditions, i.e., in (a) initial condition, and in the following true deformations: 0.3 in (b), 0.55 in (c) and 1.1 in (d).

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deformations: 0.3 in (b), 0.55 in (c) and 1.1 in (d). Fig. 8a shows that upon reheating the samples to relatively low temperatures such as 1100 ◦ C, there was no dissolution of austenite. Fig. 8b and c shows that when the samples were deformed at this temperature, both phases were present in almost equal proportions and percolated throughout the sample. Curve 4 shows a small threshold at low deformations, followed by linear hardening up to the peak stress, after which the stress level gradually declined until failure of the test specimen occurred. The difference in the plastic behaviour observed between samples with dispersed particles and lamellae percolating through the microstructure started at the end of the hardening stage. In this stage, the two phases were deformed simultaneously, elongating the austenite particles and thus reducing their thickness. This geometric deformation was discontinuous at deformation levels close to those corresponding to the peak stresses, when a process of fragmentation of the austenite particles began. As the density of small particles increased, the ferrite flowed around these particles and deformation gradually concentrated in the matrix until fracture of the sample occurred. 4. Conclusions 1. When the microstructure of duplex stainless steels is essentially ferritic during hot deformation, the mechanical behaviour is determined by the matrix and the plastic flow curve is typical of materials that undergo extended dynamic recovery. 2. The presence of austenite particles coherent with the ferritic matrix inhibits deformation of the material, increasing the initial flow stress and the hardening rate. With the extended dynamic recovery of the matrix, there is loss of this coherence and thus softening of the material. 3. When the two phases have similar volumetric fraction and appear in the form of lamellae, deformation begins in the ferrite

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