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Keywords:

  • duplex stainless steel;
  • hot tensile test;
  • phase transformation

Abstract

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

The aim of the present work was to evaluate initial secondary austenite phase microstructural changes (δ [RIGHTWARDS ARROW] γ′) in 2205 duplex stainless steel with N content of 0.36 wt% during hot tensile test. The results showed there is no secondary austenite microstructural change in zero deformation degree. The secondary austenite (γ2) phase transformation occurred only in the cooling process and under hot deformation condition. Initial secondary austenite (γ′) occurred within the ferrite matrix at the temperature range of 950–1150 °C. The precipitation of initial γ′ increased with the increase in aging time and deformation degree. With increasing time, a progressive transformation of initial γ′ to γ2 occurred. The orientation relationship between the initial γ′ phase and the δ matrix was different at varying high temperatures. A significant change in the hot deformation condition has been observed as a consequence of secondary austenite reformation. These results indicated the importance of control over the shape and volume fraction of secondary austenite phases on high nitrogen duplex stainless steels to improve hot ductility.


Introduction

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

Duplex stainless steel (DSS) has been used for a variety of applications in marine construction, chemical industries, and power plants.[1] The advantageous properties of DSS are based on a special equilibrium of the DSS microstructure. In typical cases, DSS contains about 40–60% ferrite and austenite phases. These good properties depend on the two-phase microstructure consisting of approximately equal amounts of austenite and ferrite.

Due to the high alloying content, several intermetallic precipitations and phase transformations occur during heat treatment or welding. Most of these phase transformations are concerned with ferrite because element diffusion rates are about 100 times faster in ferrite than in austenite.[2, 3] Several precipitation reactions can occur after ageing in the temperature range of 600–1000 °C, leading for example to such phases as sigma (σ), Chi (χ), γ2, M23C6, Cr2N, etc.[4, 5] Several researchers have studied secondary phase precipitation in DSS and its effect on corrosion resistance and mechanical properties.[6, 7] Among these precipitates, the sigma phase has been particularly studied since it can cause a drastic deterioration in mechanical properties of the weld metal, especially in toughness. The transformation of ferrite into austenite as the main structural change occurs during the cooling process. This transformation occurs through the nucleation and growth of austenite particles.[8] For example, during welding, the first dissolution of austenite takes place during heating, followed by grain growth in the ferrite region, and finally, reformation of austenite during cooling.[2, 9, 10] The most important metallurgically process in duplex stainless steels is the eutectic decomposition of ferrite to sigma phase and secondary austenite which can happen due to thermal effect. Despite the existence of studies on γ2, previous research mostly focused on explaining its influence on corrosion properties, especially the welding passes of DSSs.[11] Nitrogen, an austenitic stabilizing element, if added, can help stabilize the austenitic phase in DSS. By increasing nitrogen, the volume fraction of austenite phase increases from 1000 to 1350 °C. In the current work, a new DSS of 22Cr–5Ni–3Mo with N content of 0.36 wt% was developed to further reduce the cost. However, nitrogen addition is related to the microstructures and effects of the γ2 phase during hot rolling. Thus far, there is still a lack of basic understanding about the early stages of the initial secondary austenite phase transformation process and its effect on hot ductility during hot deformation. To determine the optimum hot rolling conditions, the current work focused on the initial secondary austenite phase transformation and microstructural change of high nitrogen DSS during hot tensile test.

Material and Methods

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

High nitrogen DSSs used in the present study is shown in Table 1. Cylindrical tensile specimens 10 mm in diameter and 90 mm in height were prepared. High-temperature tensile tests were carried out at temperatures ranging from 700 to 1350 °C with 5–60% deformation degree on the Gleeble-1500 thermal mechanical simulator. Two different heating methods were used. In method 1, the specimen was heated directly to a temperature between 700 and 1350 °C with average steps of 50 °C. In method 2, the specimen was heated to 1350 °C at first and then cooled to the testing temperature (1350–700 °C with average steps of 50 °C) at about 10 °C s−1. The tests were conducted at the strain rate of 1 s−1. To investigate the microstructure evolution during deformation, the specimens were quenched from the testing temperature in water immediately after deformation to a certain amount of strain. Some of the samples were also isothermally aged at 950 and 1100 °C for times of 0–960 s after deformation.

Table 1. Chemical composition of the steel used (mass fraction in %)
ElementCSiMnPSCrNiMoN
Content0.0170.721.040.0310.00122.895.473.320.36

The morphological and crystallographic characteristics of initial secondary austenite at different temperatures, aging times, and deformation degrees were investigated using optical microscopy, scanning electron microscopy, and transmission electron microscopy, respectively. Energy dispersive X-ray spectroscopic (EDS) analysis was conducted to determine the chemical composition of various phases. For the purpose of revealing the orientation relationship between adjacent grains of austenite, axis/angle pairs were derived from the corresponding selected area electron diffraction patterns.

Results and Discussion

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

The microstructures before hot tensile test were observed using optical microscopy. The images are shown in Figure 1a. An acicular γ-austenite island is clearly distributed in the δ-ferrite matrix at 950 °C. To understand the secondary austenite phase transformation process of 22Cr–5Ni–3Mo with N content of 0.36 wt%, two different heating methods were examined through the hot tension testing method. For heating method 1, the secondary austenite (γ2) phase was not observed on the micrograph, and the microstructure consisted of only ferrite and austenite in our study. For heating method 2, the secondary austenite phase transformation occurred during cooling from heating temperature to testing temperature. However, when the deformation degree was zero, secondary austenite did not occur. These results show that the secondary austenite formation depends mainly on the cooling and plastic deformation. The effect of secondary austenite phase transformation in high nitrogen DSSs is influenced by deformation temperature, deformation degree, and aging time, among other factors. In the following section, the optical micrograph of heating method 2 will be discussed in detail. Figure 1b–d present a series of optical metallographs obtained from hot tensile deformation at the same strain rate of 1 s−1, aging time of 0 s, different temperature, and deformation degree. The microstructure consists of δ-ferrite, original austenite, and secondary austenite. The ferrite, referred to as δ-ferrite, is the matrix. The acicular austenite island is the original austenite. Fine black particles intragranularly nucleated within the gray-etched ferrite matrix were obtained. These particles were identified as initial γ′ phase by energy dispersive X-ray spectroscopic analysis (EDS) (Table 2). The results indicate that while δ-ferrite is rich in Cr and depleted with regard to Ni, the initial γ′ is rich in Ni and has depleted Mo. The state of ferrite prior to transformation is of primary importance because it influences the nucleation sites for the initial γ′ phase formation. The deformation degree and deformation temperature clearly lead to the appearance of fine black particles inside the δ. These complex factors promote the nucleation of initial γ′ phase in the intragranular heterogeneities.

image

Figure 1. Microstructure of 22Cr–5Ni–3Mo with N content 0.36 wt% at the same strain rates of 1 s−1, ageing time of 0 s, and different temperature and deformation degree, a) 950 °C, ϵ = 0%, b) 950 °C, ϵ = 50%, c) 1000 °C, ϵ = 40%, and d) 1100 °C, ϵ = 15%. The light phase is austenite, fine black particles are initial secondary austenite and gray-etched matrix is ferrite phase.

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Table 2. Chemical composition of γ', δ, and γ phase determined by EDX analysis
PhaseElements (wt%)
SiCrNiMoFe
γ′0.8821.826.862.3068.16
γ0.8222.385.072.7468.46
δ0.8423.124.483.3468.02

The initial γ′ phase of 22Cr–5Ni–3Mo with N content of 0.36 wt% at the same strain rate of 1 s−1 and deformation temperature of 950 °C, as well as different aging times and deformation degrees for heating method 2, are shown in Figure 2. In Figure 2a and b, the initial γ′ particles dispersed inside the δ-ferrite. With an increase in deformation degree, the initial γ′ phase contents increased, whereas the volume fraction of the δ-ferrite phase decreased. In Figure 2c and d, the distribution of the initial γ′ phase inside the δ-ferrite was not uniform. Obvious changes in the microstructure occurred in the size, content, and shape of the initial γ′ phase. Furthermore, the volume fraction of the initial γ′ phase was determined by the aging time. As shown in Figure 2c and d, when the aging time increased to 960 s (950 °C), the initial γ′ phase transformed into the secondary austenite (γ2) phase. The nucleation of the γ2 phase is known to occur through the following transformations: δ [RIGHTWARDS ARROW] γ′ [RIGHTWARDS ARROW] γ2.

image

Figure 2. Microstructure of 22Cr–5Ni–3Mo with N content 0.36 wt% at the same strain rates 1 s−1 and deformation temperature 950 °C, and different ageing times and deformation degree. a) 240 s, 5%, b) 240 s, 25%, c) 240 s, 50%, and d) 960 s, 50%.

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The typical morphology of sympathetically nucleated initial γ′ was investigated. Figure 3a shows the transmission electron microscopy bright field images of the initial γ′ phase and the δ phase at 950 °C, with a strain rate of 1 s−1, deformation degree of 50%, and aging time of 240 s. The initial γ′ phase clearly formed inside the δ phase. To reveal the orientation relationship between neighboring grains of the initial γ′ phase, corresponding selected area electron diffraction patterns were examined. Figure 3b shows the selected area electron diffraction patterns of the initial γ′ phase and the δ matrix. Black areas in the bright field image reveal two sets of electron diffraction spots. The corresponding bright field image with specific diffraction spots of the initial γ′ phase and the matrix are superposed. The initial γ′ phase crystal plane inline image is parallel to the corresponding matrix crystal plane inline image. The initial γ′/δ phase interface is coherent at 950 °C.

image

Figure 3. Transmission electron microscopy bright field images showing initial γ′ phase and δ matrix of 22Cr–5Ni–3Mo with N content 0.36 wt% at 950 °C, strain rates of 1 s−1, 50% deformation degree, and ageing time of 240 s.

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Figure 4a shows the transmission electron microscopy bright field images of the initial γ′ phase and the δ phase at 1100 °C, with a strain rate of 1 s−1, deformation degree of 50%, and aging time of 240 s. The initial γ′ phase is a face-centered cubic. A stacking fault can be seen in the initial γ′ phase. Figure 4b and c show the selected area electron diffraction patterns of the initial γ′ phase and the δ matrix, respectively. The orientation relationship between the initial γ′ phase and the δ matrix and the poles of a close-packed plane of inline image and γ′(112)γ′ are parallel to each other. The initial γ′/δ phase interface, however, is not coherent.

image

Figure 4. Transmission electron microscopy bright field images showing initial γ′ phase and δ matrix of 22Cr–5Ni–3Mo with N content 0.36 wt% at 1100 °C, strain rates of 1 s−1, 50% deformation degree, and ageing time of 240 s.

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The curves of the reduction of area against the testing temperature range from 700 to 1350 °C at the strain rates of 1 s−1, and ageing time of 240 s for tensile testing from two different heating methods are shown in Figure 5. In the figure, the optimum hot working temperature range is between 1100 and 1150 °C at the strain rate of 1 s−1. The reduction of area values using heating method 2 was much greater than that of method 1. The hot ductility depends on different factors such as temperature, strain rate, and chemical composition, as well as the different hot deformation behavior of the ferrite and austenite phases. Under the same conditions of temperature, strain rate, and ageing time, the initial microstructure, which is determined by the volume fraction of the phases, as well as their distribution, size, shape, and orientation, is very important.[12, 13] Figure 2d shows the microstructure of the samples after hot deformation at 1100 °C and 1 s−1 by water quenching. Original austenite and secondary austenite grain size are clearly distributed in the ferrite matrix. This phenomenon is related to secondary austenite formation. Compared with traditional 2205 DSSs, the volume fraction of ferrite is decreased by the increase of N content. In the case of 0.36 wt% N, the ferrite volume fraction is lower (42%) than the traditional 2205 DSSs (45%) at 1100 °C.[14] The ferrite volume fraction is lower than the optimal 50%. But the optimum hot ductility temperature range is 1100 °C. Figure 2d and 5 indicate that secondary austenite reformation causes a significant increase in hot ductility. The results indicated the importance of distribution, size, shape, and orientation of austenite phases inside ferrite on hot ductility. In addition, the initial γ′/δ phase interface is not coherent at 1100 °C, which is conducive to austenite and ferrite co-deformation. This study has provided convincing new experimental evidence on the favorable effect of control over the shape and volume fraction of secondary austenite on hot ductility in 22Cr–5Ni–3Mo with N content of 0.36 wt%.

image

Figure 5. The variation of reduction of area as a function of temperature for 22Cr–5Ni–3Mo with N content 0.36 wt% using two different heating methods.

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Conclusions

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

The initial γ′ phase transformations have been monitored as a function of the warm processing parameters in high nitrogen DSSs. The main results can be summarized as follows: The initial γ′ phase transformation occurs only in the cooling process and under hot tensile test condition. The deformation temperature, aging time, and deformation degree affect the initial γ′ phase within the ferrite phase transformation. The secondary austenite precipitates in the temperature range of 950–1150 °C. Increasing aging time and deformation degree favors the precipitation of an initial γ′ formation in the ferrite matrix. The initial γ′/δ phase interface is coherent at 950 °C. However, the initial γ′/δ phase interface is not coherent at 1100 °C. The optimal combination of hot ductility was obtained when precipitation of initial γ′ and secondary austenite formation are stimulated, i.e., at 1100 °C. The secondary austenite formation shows the potential to improve hot working processing for 22Cr–5Ni–3Mo with N content of 0.36 wt%.

Acknowledgements

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References

This work was supported by the National Natural Science Foundation of China (Grant No. 51371123), the Research Fund for the Doctoral Program of Higher Education of China (20131402110003), Graduate Science and Technology Innovation Fund of Shanxi province and Graduate Science and Technology Innovation Fund of Taiyuan University of Technology.

References

  1. Top of page
  2. Abstract
  3. Introduction
  4. Material and Methods
  5. Results and Discussion
  6. Conclusions
  7. Acknowledgements
  8. References