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Vol. 6. Num. 3.July - September 2017
Pages 203-302
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Vol. 6. Num. 3.July - September 2017
Pages 203-302
Original Article
DOI: 10.1016/j.jmrt.2017.01.001
Open Access
Texture evolution and phase transformation of 25Cr-6Mo-5Ni experimental duplex stainless steel during hot and cold rolling
Mohammad Masoumia,b,
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Corresponding author.
, Francisco Evarsito Uchôa Reisa, Mirela Oliveira de Castroa, Miloslav Béresa, Hamilton Ferreira Gomes de Abreua
a Universidade Federal do Ceará, Centro de Tecnologia, Departamento de Engenharia Metalúrgica e de Materiais, Fortaleza, CE, Brazil
b Universidade de São Paulo, Escola Politécnica, Departamento de Engenharia Metalúrgica e de Materiais, São Paulo, SP, Brazil
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Tables (1)
Table 1. Chemical composition (wt.%) of experimental 25Cr-6Mo-5Ni stainless steel.

An experimental as-cast 25Cr-6Mo-5Ni stainless steel has been solution annealed at 1250°C and subjected to hot and then cold rolling. X-ray diffraction, optical microscopy and electron backscatter diffraction were used to investigate the effect of hot and cold deformation on the phase transformation and texture evolution. The results revealed that dominant {100}//ND and {110}//ND texture components of martensite is originated by shear strain generated between rolls and sheet surface. The Kernel average misorientation augmented significantly with increased strain and decreased deformation temperature. The internal grain structure becomes more heterogeneous with the dislocation piles up preventing dislocations movement. High localized stresses were developed at grain boundaries due to different deformation of individual grains, which enhanced martensitic transformation in these regions.

25Cr-6Mo-5Ni stainless steel
Phase transformation
Kernel average misorientation
Full Text

High alloy duplex stainless steels, commonly referred as super duplex, have high corrosion resistance comparable to high Mo content super austenitic steels [1]. Both the super duplex and the super austenitic stainless steels are ranked by an empirical formula that calculates the pitting resistance equivalent number (PRN=%Cr+3.3% Mo+20%N) [2]. The super austenitic stainless steels present higher PRN values (45) most because of the higher Mo content ∼6wt.%. Their drawback is a high Ni content, which increases cost significantly. Thus, there is a need for a lower cost alternative, i.e. duplex stainless steel with increased Mo content and high PRN values.

A suitable balance between chemical composition, processing and post treatment plays significant role in the development of new 25Cr-6Mo-5Ni stainless steel. Some investigations [3–5] reported that higher Mo content increases PRN and improves naphthenic acid corrosion resistance, while the body-centered cubic (BCC) structure restricts the workability. However, the retained austenite can improve the mechanical properties of the steels by strain-induced martensitic transformation. Therefore, the high Mo-Ni stainless steels can be used widely in offshore oil and gas exploration in addition to chemical and petrochemical application due to high resistance to pitting and crevice corrosion, resistance to corrosion fatigue, high strength and good weldability.

In general, ferritic stainless steels have inferior mechanical properties in comparison to austenitic [6]. In contrast, the ferritic stainless steels are considered to resist better to stress corrosion cracking [7]. The 25Cr-6Mo-5Ni stainless steels have a good combination of mechanical properties and corrosion resistance due to their dual crystal structure. The mechanical properties of duplex stainless steels strongly depend on various factors such as microstructure (volume fraction and morphology of retained austenite, grain size), ferritic phase mechanical properties, crystallographic texture, grain boundary and local strain distribution. The retained austenite grains that transform into martensite during deformation increases ductility of material by consumption of stored energy.

Zaid et al. [8] reported that austenite shows dominant brass texture, whereas strong alfa-fiber was developed in ferrite during rolling. Liu et al. [9] studied the plastic deformation mechanism in duplex stainless steel. They discovered deformation heterogeneity in phases and showed that ferrite phase γ-fiber components were more intense during multi-pass hot compressions.

The texture evolution in 25Cr-6Mo-5Ni stainless steels is affected by two mechanisms, plastic deformation and martensitic phase transformation. Although, the preferred crystallographic orientations of both austenite and martensite can play a significant role in the physical and mechanical properties, to date less attention has been paid. Therefore, in this work, the texture evolutions during both hot and cold rolling were investigated as a function of thickness reduction to describe the phase transformation and crystallographic orientation changes. A better understanding of structure changes during deformation is required for proper assessment of service life, cost efficient design and failure analysis.

2Experimental procedure

The material used in the current study is a modified super 25Cr-6Mo-5Ni stainless steel produced in an electric induction furnace. The chemical composition is listed in Table 1. The experimental as-cast ingot was solution annealed at 1250°C for 20min followed by water quenching to prevent the formation of hard and brittle precipitation, i.e. Cr-rich phase. In the solution annealed condition, ferritic matrix with retained austenite consisting of both metastable Widmanstätten and blocky austenite was revealed, Fig. 1.

Table 1.

Chemical composition (wt.%) of experimental 25Cr-6Mo-5Ni stainless steel.

Si  Mn  Ni  Mo  Cr  Nb+Ti 
0.014  1.18  0.58  5.2  6.07  25.1  0.03  0.03  0.09 
Fig. 1.

Optical micrograph of solution annealed sample at different magnifications (F, ferritic; WA, Widmanstätten austenite; BA, blocky austenite).

The solution annealed ingot was cut to 6mm thick sample and subjected to hot rolling at 1250°C using Stanat model TA-315 rolling machine with the rolling speed of 0.2ms−1 to reach a 50% reduction (t=3mm) and was water quenched. Further cold-rolling with 50 and 67% reduction was conducted parallel to the direction of the hot rolling. In this respect, the rolling was performed with the rolling speed of 0.2ms−1 in 15 passes to achieve the final thickness of 1.5 and 1.0mm, respectively. In this study, for simplicity, samples are referred as 50% and 67% cold rolled.

Then, the volume fraction of retained austenite of all specimens was measured by X-ray diffraction (XRD) technique using Philips X-Pert diffractometer equipped with Cu Kα radiation. Integrated intensities of the {110}, {200}, {211} and {220} peaks of ferrite and the {111}, {200}, {220} and {311} peaks of austenite were used for calculation of volume fraction of retained austenite using X-Pert software.

Macro-texture measurements have been carried out at the surface of the rolled sheets perpendicular to normal direction (ND). Three incomplete pole figures, {110}, {200}, and {211} for ferrite and four {111}, {200}, {220} {311} were measured using Phillips X-Pert goniometer equipped with Cu Kα radiation. The orientation distribution function (ODF) of each sample was determined from the measured pole figures using the MTEX software [10]. Bunge's angles were adopted to describe orientations, and the φ2=45° section of Euler space was used to display the computed ODFs.

The electron backscatter diffraction (EBSD) measurements were conducted in a plane perpendicular to the transverse direction of the sheet. EBSD data were acquired using Phillips XL30 scanning electron microscope operating at acceleration voltage of 20kV, sample tilt angle of 70°, working distance of 12mm and 0.5μm step size. HKL Channel 5 software was used to analyze the data. The Kernel average misorientation was calculated from EBSD data to evaluate and map local plastic strain and recrystallization degree in different regions of each sample.

3Results and discussion3.1X-ray diffraction

Fig. 2 shows the volume fraction of both the FCC phase (retained austenite) and the BCC phase (containing ferrite and transformed martensite) determined from XRD patterns. In solution annealed condition, volume fraction of retained austenite was found to be around 33vol.% with no significant change in hot-rolled condition. With progressing deformation degree at room temperature, the volume fracture of retained austenite decreased from 33vol.% to 13 and 10vol.% in 50 and 67% cold rolled specimens, respectively. This reduction in volume fraction can be explained by strain induced martensitic transformation. It is expected that dislocation density increased significantly during cold deformation. Thus separated dislocation walls can enhance transformation of retained austenite into martensite.

Fig. 2.

Volume fraction of BCC and FCC phases during heat treatment and post deformations determined by XRD technique.

3.2Texture evolution

In 25Cr-6Mo-5Ni stainless steels, the retained austenite grains deformed first and transformed into BCC martensite under strain. Recently, Akdut et al. [11] have reported that in duplex stainless steel subjected to deformation the ferrite texture became more isotropic. In austenite however, Brass (110)[112] and Goss (011)[110] texture components were developed. Therefore, in the current work texture evolution during deformation were studied in each phase separately.

3.2.1Austenite texture evolution

Slip and twinning are the most important mechanisms of plastic deformation in austenite with low stacking fault energy (SFE) [12,13]. In general, slip occurs on the closed packed planes and in the direction of shortest distance between atoms. The crystallographic texture evolution in FCC austenitic phase during hot and cold rolling can be tracked by orientation distribution function (ODF) analysis and is presented in Fig. 3.

Fig. 3.

ODFs at φ2=45°, FCC austenite texture evolution of (a) solution heat treated, (b) hot rolled, (c) 50% cold rolled, and (d) 67% cold rolled samples.

Although the ODF intensity was weak in all conditions, the {112}//ND, cube (001)[01¯0], transformed copper (113)[11¯0] and (111¯)[21¯1] components developed in solution heat treated sample, Fig. 3a and b shows the evolution of texture during rolling at 1250°C. It is observed that the intensity of texture decreased during hot deformation. The (001) 130 and (110)[11¯3] texture components were developed in this sample. Dynamic recrystallization during hot deformation caused the formation of (001) 130 components with 20° deflection from ideal cube component. The formation of {110}//ND texture fiber, predominant at close to the Brass (110)112 component, was observed and is fully consistent with Refs. [14,15]. The further cold rolling increased intensity of the {001}//ND fiber, with prominent rotated cube (001)110 and (001)150 components, Fig. 3c. Therefore, strain accumulation during deformation at room temperature is responsible for development of {001}//ND textures. Finally, cube component was dominant in 67% cold rolled samples, Fig. 3d. The absence of Copper (112)111 and Brass (110)112 components in retained austenite grains of both hot and cold rolled samples is attributed to the absence of shear bands in the deformed matrix.

3.2.2Ferrite/martensite texture evolution

The texture evolution of BCC ferrite/transformed martensite is presented in Fig. 4. In solution treated sample, a dominant (110)[331] texture was observed. The {110}//ND components were characterized with the highest intensity in all investigated specimens, expect in 67% cold rolled samples where {100}//ND fiber was dominant. The {100}//ND components showed a gradual increase with progressing deformation, reached the highest intensity in 67% cold rolled sample. The {111}//ND fiber texture was weak in all specimens. According to the crystallographic model [16], a BCC lattice can be formed from a FCC lattice by compressing one axis and expanding the other two perpendicular axes. Thus, for the FCC crystal under compression of {100} plane parallel to the rolling direction, the martensite can be developed in the same direction. The texture evolution during cold rolling confirmed that the newly formed BCC grains were aligned with (100) and (110) planes parallel to the normal direction (Fig. 4c and d). It is notable that an increase in {100}//ND ferritic component in cold rolled sample is accompanied with the decreasing in {110}//ND texture in FCC phase [10]. In addition, during cold rolling, available retained austenite at grain interiors was transformed into martensite by specific orientation relationship.

Fig. 4.

ODFs at φ2=45°, BCC ferrite/martensite texture evolution (a) solution-treated, (b) hot rolled, (c) 50% cold rolled, and (d) 67% cold rolled samples.

3.2.3Selective austenite to martensite transformation

If the parent austenite phase has a preferred crystallographic orientation, the product phase is also expected to be textured. The elevated temperature during hot rolling promoted the grain boundaries mobility across the {110} and {111} plans. As a result, the {110}//ND texture fiber was developed during hot rolling. Therefore, γ-fiber emerged together with a strong {110}//ND texture during cold rolling. Cube (001)[010] component that developed in austenite after cold deformation has lowest yield strength and therefore a local weakening effect on the plastic properties [17]. This cube (001)[010] texture component of austenite converted to rotated cube (001)[110], (110)[11¯0] rotated Goss and Goss (110)[001] component in ferrite. Therefore, this texture component was intensified at the last deformation process in both BCC and FCC structure.

Fig. 5 shows the correlation between texture components during austenite to martensite transformation, considering 24 variants in Kurdjumov–Sachs (K–S) where {111}γ//{110}α, {110}γ//{111}α, are orientation relationships. For example, {100}〈01¯0〉 component in austenite can generate {100}〈11¯0〉, {100}〈1¯1¯0〉, {110}〈11¯0〉 and {110}〈001〉. Since the {100}001 component has the highest Taylor factor among all FCC components, it accumulates high dislocation density [18]. Accordingly, during elevated deformation temperature the cube {100}001 component is known to be the first to undergo dynamic recrystallization, therefore it disappears because accumulates the highest stored energy [19]. As a consequence, it can enhance austenite to martensite transformation. This can be confirmed when comparing the texture in austenite after deformation. Another texture component of interest is the {113}〈11¯0〉 component, which become strongest component in alloys with a low SFE under torsion [20,21]. This implies that the {113}〈11¯0〉 component might selectively undergo martensitic transformation similar to the {100}001 component. According to Fig. 5, the parent austenite component can transform into a BCC Copper component having {113}〈11¯0〉 orientation.

Fig. 5.

Schematic representation of texture component in the φ2=45° section between transformation austenite into martensite considering K–S orientation relationship. The figures identify the (a) parent FCC orientation from which a specific (b) BCC orientation originate adopted from [17].

3.3EBSD studies

To further investigate the effect of hot and cold rolling, EBSD analyses were conducted in initial, hot-rolled and cold rolled samples. Retained austenite and ferrite can easily be indexed by EBSD technique due to the different crystallographic structure. In contrast, martensitic phase is difficult to identify due to high dislocation density in martensite, which results in poor quality of Kikuchi patterns. During sample preparation the retained austenite near the surface transformed into martensite. Indeed, it was observed that the volume fraction of retained austenite determined by EBSD technique was lower than the amount obtained by XRD. This can be explained by differences in penetration depth between these techniques.

In low carbon content (<0.02%) steel, BCC martensite with low tetragonality is present. Due to the lattice distortion in deformed ferrite/transformed martensite it is difficult to match Kikuchi pattern with undeformed crystal leading to differences in quality of band contrast. Fig. 6a shows band contrast map of cold deformed sample. In the inset image, the frequency distribution of band contrast is shown. One can see that two different peaks are present in this sample indicating separation of undeformed ferrite and deformed ferrite/transformed martensite. This is presented separately in Fig. 6b and c. However, to separate between deformed ferrite and transformed martensite using EBSD technique is not possible.

Fig. 6.

Band contrast map of 50% cold deformed sample, (a) EBSD map, (b) ferrite (high band contrast values), and (c) deformed ferrite/transformed martensite (low band contrast values).

The EBSD observation of the solution annealed sample indicated the elongated retained austenite was predominantly found at grain boundaries. Microstructure which consists of elongated grains of retained austenite in BCC ferrite/transformed martensite matrix is very beneficial for improvement of the ductility and strength. Higher stresses are applied at the grain boundaries because each grain deforms by different slip system and grain rotation, which induces higher localized stress at boundary. Shen et al. [20] have reported that the mechanical stability of retained austenite is affected by its morphology. The retained austenite is more stable if is surrounded by martensitic phase, which have higher deformation resistance. This martensitic transformation at the grain boundaries could be ascribed to the higher localized stresses. In contrast, if retained austenite is located at grain boundary it exhibits lower stability.

The inverse pole figure (IPF) maps were used to illustrate grain orientations. The ODFs (Fig. 7a3, b3 and c3) of each phase were calculated separately from IPF maps (Fig. 7a1, b1, and c1). The ODF analysis of initial material showed that (112)[111] (copper) (φ1,ϕ,φ2: 90,35,45°) component was the main orientation of retained austenite, whereas the (221)[212] (φ1,ϕ,φ2: 60,80,45°) was the dominant texture component in ferrite. In hot rolled samples the (332)[113] (φ1,ϕ,φ2: 90,65,45°) was predominant in the FCC phase. Ray et al. [21] showed beneficial effect of this texture component for strain induced martensitic transformation leading to improvement of strength and toughness. Also, (111)[101] (φ1,ϕ,φ2: 50,75,45°) component was observed in hot deformed sample. In cold rolled samples, Fig. 7c3, due to shear strain generated between rolls and sheet surface, retained austenite transformed completely and dominant (110)[110] (rotated Goss) (φ1,ϕ,φ2: 0,90,45°) texture developed. This suggests that the martensitic transformation is related to the applied shear strain.

Fig. 7.

EBSD analysis (a) solution-treated, (b) 50% hot rolled, (c) 50% cold rolled. (a1. b1, c1) IPF, (a2, b2, c2) phases identifications, and (a3, b3, c3) related BCC and FCC ODFs, and (a4, b4, c4) KAM.

IPF maps (Fig. 7a1, b1 and c1) demonstrate that the ferrite grains were coarse and quasi-polygonal in initial condition then decreased gradually with progressing deformation. This trend grain size reduction can be attributed to the increased density of low angle grain boundaries (LAGBs). The misorientation analysis obtained by EBSD demonstrated that the ratio of LAGBs increased from ∼35% in initial material to ∼85 and ∼95% in hot and cold rolled samples, respectively. Then, various deformation bands are formed at grain interior, next dislocations are trapped at subdivide grain boundaries which lead to formation of LAGBs at grain interior as described by Belyakov et al. [22].

Kernel average misorientation (KAM) can be used to monitor dislocation motion including dislocation density inside individual grain during deformation [23,24]. The KAM is represented by the average misorientation between a given point and its nearest neighbor belonging to the same grain and associated with a misorientation less than 15°. An increase in the KAM is closely associated with the cumulative dislocation density. The KAM maps (Fig. 7a4, b4 and c4) obtained from corresponding EBSD maps demonstrates the strain distribution evolution during deformation, where red color indicates highest KAM value. In the solution annealed specimen, the local misorientation was found to be non-uniform distributed where higher intensity was observed near grain boundary regions and was found to be low inside the grain. During hot rolling, the magnitude of the local misorientation increased, while recovery and dynamic recrystallization took place which led to annihilation and reduction of dislocation density. Because of low SFE in austenite, partial recrystallization took place. However recovery was the dominant mechanism in ferrite due to high SFE. In addition the local misorientation increased inside grains and grain boundaries. Micro-bands were formed, indicating that these regions suffered majority of the deformation. In cold rolled sample, the KAM values increased significantly in both at grain boundaries and inside the grains. In this condition, the fraction of high KAM values exceeded 90% and indicated more homogeneous distribution. Furthermore, the retained austenite was completely transformed into martensite. Also, Winther et al. [25] has reported that at higher strain the newly formed martensite have the highest KAM value due to the displacive transformation mechanism. The newly formed martensite has higher hardness in comparison to soft ferrite. In [26] it was demonstrated that crack formation is enhanced by variation in hardness between individual phases. Results obtained in this work can be utilized to enhance precision of material models [27,28].


The effect of hot and further cold deformation on microstructure, phase transformation and texture evolution in the experimental 25Cr-6Mo-5Ni stainless steel were investigated. Following conclusions can be drawn:

  • -

    The microstructure of solution annealed sample consists of ferrite and austenite.

  • -

    With progressing deformation degree at room temperature, the volume fracture of retained austenite decreased from 33vol.% to 13 and 10vol.% in 50 and 67% cold rolled specimens, respectively.

  • -

    The BCC transformed martensite showed dominant {100}//ND and {110}//ND texture components. The dominance of {100}//ND component was attributed to the martensitic transformation whereas {110}//ND component was induced by shear strain generated between rolls and sheet surface.

  • -

    High localized stresses were developed at grain boundaries due to different deformation of individual grains and enhanced martensitic.

  • -

    Kernel average misorientation increased significantly with progressing degree of deformation and reduction in rolling temperature.

Conflicts of interest

The authors declare no conflicts of interest.


The authors acknowledge the Brazilian research agencies CNPq and CAPEs, the Research Board of the Federal University of Ceará for the financial support and Laboratório de Caracterizac¸ ão de Mate-riais (LACAM) and Analytical Center (CT-INFRA/MCTI-SISNAD) for the provision of research facilities of this work.

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