Static Tensile Deformation Behavior of a Lean Duplex Stainless Steel Studied by In Situ Neutron Diffraction and Synchrotron Radiation White X-rays

Abstract

To investigate the tensile deformation behavior of a lean duplex stainless steel (S32101) from the viewpoints of plastic deformability among phases or grains, we performed static tensile tests, in situ neutron diffraction, and white x-ray diffraction experiments at room temperature. In the static tensile tests, the S32101 steel displayed a larger uniform elongation and a better tensile strength–uniform elongation balance than a commercial SUS329J4L duplex stainless steel. A larger uniform elongation of S32101 is associated with the macroscopic work hardening behavior that a work hardening rate higher than the flow stress can maintain up until high true strains. From the experimental results of synchrotron radiation white x-ray diffraction experiments, the hard phase of S32101 was changed from the ferrite (α) phase to austenite (γ) one during tensile deformation. This led to a larger stress partitioning between the phases at the latter stage of deformation. From the experimental results of in situ neutron diffraction, it was found that the stress partitioning of the γ phase in the S32101 was the largest among the present results. Therefore, the larger work hardening rate of S32101 can be explained by the large stress partitioning of the γ phase, that between γ and α phases and γ volume fraction.

1. Introduction

Duplex stainless steels consist of approximately equal amounts of ferrite (α) and austenite (γ) phases and exhibit a good combination of mechanical properties and corrosion resistance. These steels are widely used in many industrial facilities, such as chemical plants and hydropower plants. Duplex stainless steels with a higher level of nitrogen, termed the second generation of duplex stainless steels, are now in widespread use and are classified as standard grade, super duplex grade, and lean duplex grade.^1,2,3) Lean duplex stainless steels that have a low Ni and Mo content are balanced by having an increased N and Mn content to assure a balanced microstructure with approximately equal amounts of α and γ phases.¹⁾ Lean duplex stainless steels with a higher strength need to be developed to extend their application range. The widespread use of such steels would enable both cost and weight savings, and thus, lean duplex stainless steels are expected to be increasingly used as an alternative to conventional austenitic steels. To achieve this goal, it is necessary to investigate in detail the plastic deformation behavior of duplex stainless steels, such as their mechanical properties and strength–elongation balance.^1,4)

A number of studies involving neutron and synchrotron x-ray diffractions have been carried out to investigate the deformation behavior of various steels and alloys.^5,6,7,8) Neutron and synchrotron x-ray diffraction measurements, because of their high penetration abilities, can help ascertain bulk strains in engineering materials. It is also possible to monitor changes in diffraction profiles during tensile and compressive deformation by in situ diffraction experiments under loading.^5,8,9,10) In engineering materials, inhomogeneous plastic deformation occurs because of the difference in stress among the constituent grains or phases.^9,10,11) The deformability of the grains or phases can be determined from their neutron and synchrotron x-ray diffraction profiles, and these forms of diffractions are closely related to macroscopic deformation.^5,9,10) Figure 1 is a schematic illustration of stress–strain curves for a duplex stainless steel and its constituent phases.^5,12) In duplex stainless steels, the plastic deformation behavior is different between the α and γ phases, which can be usually classified, respectively, as the hard phase and soft phase. When applied stress is removed after tensile deformation, a tensile residual stress is observed in the hard phase, and a compressive residual stress is observed in the soft phase.^4,5) Plastic deformability can also be investigated by stress partitioning between the [hkl] family grains in each phase or between the constituent phases.^5,10,11) Because the crystal structures of the α phase and γ phase are different in duplex stainless steels, tensile deformation behavior can be discussed from the standpoints of stress partitioning not only between the γ and α phases but also among [hkl] family grains at the constituent phases.^4,5) In this study, we performed static tensile tests and in situ neutron diffraction experiments during tensile deformation as well as synchrotron radiation white x-ray diffraction experiments on a lean duplex stainless steel. The static tensile properties and work hardening behavior of the steel were discussed from the viewpoints of inhomogeneous plastic deformation between the α phase and γ phase or among [hkl]-oriented family grains based on the experimental results.

Fig. 1.

Schematic illustration of stress–strain curves for duplex stainless steel, γ and α phases.

2. Experimental Procedures

2.1. Materials and Microstructure Observations

A lean duplex stainless steel (S32101) and a JIS-SUS329J4L (329J4L) steel were used in the experiments. The chemical compositions of the steels are summarized in Table 1. Ingots of the two steels were vacuum melted and reheated at 1453 K for 3.6 ks, hot rolled from 168 to 5 mm, and water quenched. The plates were then cold rolled to the final thickness of 3 mm (S32101) or 2 mm (329J4L) after being heat treated at 1323 K for 30 s. These sheets were then annealed at 1323 K for 30 s followed by air cooling. The microstructures were examined by an optical microscope and electron backscattering diffraction (EBSD) and the deformed microstructures were also observed by transmission electron microscopy (TEM). The volume fractions of the γ and α phases before tensile deformation were calculated by x-ray diffraction experiments^13,14) and the average grain sizes were measured by the linear intercept method. To investigate the element partitioning between γ and α, an electron probe micro-analyzer (EPMA) was used on samples taken from the heat-treated plates under conditions of an acceleration voltage of 15 kV and a probe current of 325 nA.

Table 1. Chemical compositions of the S32101 and SUS329J4L steels and those for the austenite (γ) and ferrite (α) phases (mass%).

		C	Si	Mn	P	S	Ni	Cr	Mo	Cu	N
S32101	specimen	0.022	0.57	4.96	0.022	0.0004	1.51	21.22	0.33	0.23	0.217
	γ	–	0.53	5.28	–	–	1.86	19.9	0.33	0.21	0.336
	α	–	0.6	4.85	–	–	1.34	21.9	0.45	0.17	0.056
SUS329J4L (S31260)	specimen	0.018	0.49	0.69	0.028	0.0005	6.95	24.88	3.04	0.17	0.133
	γ	–	0.53	0.74	–	–	8.2	23.1	2.68	0.18	0.19
	α	–	0.58	0.69	–	–	5.69	25.6	3.67	0.14	0.01

2.2. Static Tensile Tests

Tensile test specimens with a gauge length and gauge width of 25 mm and 5 mm were prepared from each duplex stainless steel. Static tensile tests were conducted at 296 K with an initial strain rate of 3.3 × 10^–4 s^–1. Test samples deformed by various true strains were also prepared for the synchrotron radiation white x-ray diffraction experiments, to determine the residual lattice strains for the γ and α phases in the steels.¹⁴⁾

2.3. In Situ Neutron Diffraction Experiments

In situ neutron diffraction measurements (wavelength: 0.19 nm) were performed during stepwise tensile testing using the angular dispersive neutron engineering diffractometer (RESA-1) in the JRR-3 (Japan Research Reactor No. 3) at the Japan Atomic Energy Agency (JAEA).^9,10,15) For the neutron diffraction experiment, tensile test specimens with a gauge length and gauge width of 80 mm and 3 mm, respectively, were prepared. The sampling volume for neutron diffraction was approximately 20 mm³ in the parallel portion of the tensile test specimen. The specimens were deformed by stepwise tensile loading and the {111}, {200}, and {311} diffraction profiles for the γ phase, and the {110}, {200}, and {211} diffraction profiles for the α phase were measured at each loading step. The collected diffraction patterns were analyzed using a single-peak fitting approach to determine the peak position, i.e. lattice spacing (d_hkl) for each diffraction. The lattice strain ε_hkl was determined by measuring change in the lattice spacing (d_hkl) from a stress-free reference ( $d_{hkl}^0$ ), according to Eq. (1):^5,11)   

$${\mathit{\text{ε}}}_{hkl}=\frac{d_{hkl}-d_{hkl}^0}{d_{hkl}^0}=-cot\left({\theta }_{hkl}\right)\Delta {\theta }_{hkl}.$$(1)

The lattice strain can be described by a change in diffraction angle as shown in above equation. From the difference of ε_hkl calculated by Eq. (1), we can discuss stress partitioning between the hard and soft grains for the γ and α phases.¹⁰⁾

2.4. Synchrotron Radiation White X-ray Diffraction Experiments

Synchrotron radiation white x-ray diffraction experiments were performed by using the BL14B1 at the SPring-8.⁸⁾ Here, the test samples with various true strains described in 2-2 were used and the diffraction profiles, which can be measured in all specimens with various true strains, were discussed. The synchrotron radiation white x-ray beam was incident at the center area of the tensile specimens, and the profile data were obtained by using an energy dispersive x-ray diffraction technique. The transmitted diffracted x-ray was detected by using a solid-state detector at a fixed angle of 10 degrees. In this case, residual lattice strain can be calculated by the change in the diffraction energy as follows:⁸⁾   

$${\mathit{\text{ε}}}_{hkl}=\frac{E_{hkl}^0-E_{hkl}}{E_{hkl}},$$(2)

where $E_{hkl}^0$ and E_hkl are diffraction energies before and after tensile loading, respectively. Because tensile test specimens with various strains were used in the white x-ray diffraction experiments, residual lattice strains were calculated by Eq. (2). In this case, residual phase strain (ε_ph.) can also be calculated by using the estimated lattice parameters for the γ and α phases.¹⁶⁾   

$${\mathit{\text{ε}}}_{ph.}=\frac{a_{ph.}-a_{ph.}^0}{a_{ph.}^0},$$(3)

where $a_{ph.}^0$ and a_ph. are the lattice parameters before and after tensile loading, respectively.

3. Results and Discussions

3.1. Microstructures

Figure 2 shows optical micrographs of S32101 (a) and 329J4L (b) steels. The morphology of γ and α are banded and appear elongated in the rolling direction.^1,4) The volume fractions of γ calculated by x-ray diffraction experiments for S32101 and 329J4L were 51% and 39%, respectively. The average grain sizes measured by the linear intercept method for S32101 and 329J4L were 6.4 and 4.2 μm, respectively. The volume fraction of γ in S32101 was approximately 10% greater than in 329J4L, and the average grain size of 329J4L was a little smaller than that of S32101. Figure 3 shows scanning electron microscopy (SEM) and orientation imaging microscopy (OIM) images for the γ and α phases in S32101 and 329J4L steels in the TD plane. The grain size was different between the γ and α phases, and the average grain size for the γ phase was smaller than that for the α phase in both steels.

Fig. 2.

Optical micrographs of the S32101 (a) and SUS329J4L (b) steels.

Fig. 3.

SEM and OIM images for austenite and ferrite phases in the S32101 (a) and SUS329J4L (b) steels.

3.2. Static Tensile Properties

Figure 4 shows the nominal stress–strain curves (a) and the true stress–strain curves (b) for S32101 and 329J4L. The mechanical properties are summarized in Table 2. The 0.2% proof stress and tensile strength for S32101 were smaller and the uniform and total elongations were larger than those for 329J4L. As shown in Fig. 4(b), a higher work hardening rate results in a larger uniform elongation of S32101. This seems to be associated with the difference of volume fractions of γ and α.^5,12) At the same time, no stress-induced martensitic transformation from γ to α was observed in either of the steels. Figure 5 shows the relationships between tensile strength and uniform elongation in various duplex, austenitic, and ferritic stainless steels.^10,14,15,17) The dashed lines in Fig. 5 are contour lines of the product of tensile strength and uniform elongation. S32101 showed a better tensile strength–uniform elongation balance than other duplex stainless steels (SUS329J4L, SUS329J1). This is attributed to the good uniform elongation of S32101. It is very interesting that the strength–elongation balance of S32101 is almost the same as that of JIS-SUS310S, which is a stable austenitic stainless steel. Figures 4 and 5 indicate that S32101 showed a better combination of tensile strength and uniform elongation than other duplex stainless steels because of a larger uniform elongation. The reason for the good uniform elongation of S32101 is that a larger work hardening rate can be maintained up until high strains.

Fig. 4.

Nominal stress–strain curves (a) and true stress and work hardening rate as a function of true strain (b) in the S32101 and SUS329J4L steels obtained by the static tensile tests.

Table 2. Mechanical properties of the S32101 and SUS329J4L steels obtained by the static tensile tests at room temperature.

	0.2% proof stress (MPa)	Tensile strength (MPa)	Uniform elongation (%)	Total elongation (%)	Local elongation (%)
S32101	553	812	33.5	60.4	26.9
SUS329J4L (S31260)	620	856	24.8	37.3	12.5

Fig. 5.

Tensile strength–uniform elongation balance in various duplex, austenitic and ferritic stainless steels.^10,14,15,17)

3.3. Stress Partitioning Studied by Neutron and White X-ray Diffraction Experiments

In this subsection, the larger work hardening of S32101, as seen in Fig. 4, is discussed by using the experimental results of in situ neutron diffraction and white x-ray diffraction experiments. Here we focus on three points, as shown in Table 3:

Table 3. Summary of microstructure and deformability of γ and α phases in the S32101 and SUS329J4L steels.

				S32101	SUS329J4L
Microstructures	1	Volume fractions of γ and α phases		γ: 51%, α 49%	γ: 39%, α 61%
	2	Average grain size		6.4 μm	4.2 μm
Deformability of γ and α phases (Stress-strain relationships of γ and α phases)	3	Hard phase	a. Residual lattice strain (Figs. 6,7) b. Lattice strain (Fig.8)	The hard phase was α at the initial stage of deformation. But the hard phase was changed from α to γ at true strains more than 0.12.	α phase is the hard one.
	4	Stress partitioning between the γ and α phases	Residual phase strain (Fig. 7)	The stress partitioning became larger at true strains more than 0.12.	The stress partitioning is decreased with increasing of strain.
	5	Stress partitioning among [hkl]-oriented family grains	a. Difference of lattice strains (Fig. 9) b. Difference of residual lattice strains (Fig.6)	· The stress partitioning for γ phase in the S32101 was the largest in the present results. · The stress partitioning for γ phase was larger than that for α phase in the both steel. · The stress partitioning for α phase was almost the same between the S32101 and the 329J4L.

a. Which phase is harder, the γ phase or α phase ?

b. Stress partitioning between the γ and α phases

c. Stress partitioning between the [hkl] family grains and work hardening behavior in each phase

Figure 6 shows residual lattice strain as a function of true strain for the γ and α phases in S32101 and 329J4L steels as obtained by the white x-ray diffraction experiments. The residual stress for each phase or grain can be discussed by the positive (tensile) residual lattice strain or negative (compressive) strain.^5,18) The residual lattice strains at a true strain of zero means those at the elastic limit. The plastic deformability, not only between the γ and α phases but also among [hkl] family grains at the same phase, can be seen in Fig. 6. The α phase is a hard phase in 329J4L because α phase showed tensile residual lattice strains, whereas the γ phase showed compressive strains. This is consistent with the previous results for duplex stainless steels.⁵⁾ In terms of S32101 on the other hand, the difference in residual lattice strains between the γ and α phases was smaller than that for 329J4L. The residual lattice strains for the α phase decreased whereas those for the γ phase increased with an increase in true strain. This means that the change in flow stress for the γ phase is larger than that for the α phase. This change in residual lattice strains for γ and α phases with true strain also applies for the 329J4L steel. The difference in residual lattice strains in the γ phase of S32101 is also found to be larger than those for other results. The difference in lattice strains among the grains is related to stress partitioning and the work hardening behavior of the phase.^9,10) The larger the stress partitioning, the higher the degree of work hardening. Therefore, the work hardening behavior of the γ phase is larger than that of the α phase in S32101.

Fig. 6.

Residual lattice strain as a function of true strain in the S32101 (a, b) and SUS329J4L (c, d) steels obtained by synchrotron radiation white x-ray diffraction experiments.

Figure 7 shows phase strain as a function of true strain for the γ and α phases in S32101 and 329J4L steels. Because the phase strain means the average value of the residual lattice strains of each phase, from Fig. 7 we can discuss the inhomogeneous plastic deformation between the γ and α phases in detail. In the case of 329J4L, α is a hard phase, as shown in Fig. 6, and the difference in phase strain between γ and α decreased slightly with an increase in true strain. On the other hand, in S32101, the phase strain of the α phase was larger than that for the γ phase in the early stage of plastic deformation. However, the difference in phase strains is smaller with true strain. After a true strain of 0.12, the phase strain of the γ phase was larger than that of the α phase and the difference in phase strains, i.e., stress partitioning between γ and α phases, also became larger with an increase in true strain. That is, the hard phase changed from an α phase to a γ phase during tensile deformation in S32101 and the inhomogeneous plastic deformation between the α phase and the γ phase became larger in the latter stage of deformation. In order to discuss phase stress and stress partitioning between γ and α phases in more detail, in situ neutron diffraction experiments during loading using the time-of-flight type engineering neutron diffractometer¹⁹⁾ should be required in the future.

Fig. 7.

Phase strain as a function of true strain for austenite and ferrite phases in the S32101 and SUS329J4L steels.

Figure 8 shows lattice strain as a function of applied stress in S32101 (a) and SUS329J4L (b) steels obtained by in situ neutron diffraction experiments to discuss the work hardening behavior of the γ and α phases. YS in Fig. 8 is the 0.2% proof stress for each specimen. At the initial stage of deformation, a linear relationship was observed between ε_hkl and the applied stress.^5,9) As ε_hkl changed under an increase in applied stress, deviation from the linear relationship was observed prior to YS. P1 in Fig. 8 is the onset stress of the deformation stage, at which point the soft phase is plastically deformed and the hard phase is still deformed elastically^5,9,10) For example, in the γ phase of Fig. 8, ε₁₁₁ increases slowly and ε₂₀₀ increases rapidly with the applied stress, implying that [111]-oriented family grains are soft and that [200]-oriented family grains are hard. Such differences in deformation among [hkl]-oriented family grains are associated with stress partitioning between the hard and soft grains.^10,11) As shown in Fig. 8, the hardest grain in 329J4L is [200]-oriented α grain, while that in S32101 is [200]-oriented γ grain. Meanwhile, the grain that started plastic deformation at P1 was [111]-oriented γ grain for 329J4L and [110]-oriented α grain for S32101. In Fig. 9, the differences of lattice strain between the hard grain and soft grain are summarized in each phase on the basis of Fig. 8. The difference of lattice strain is associated with the stress partitioning or work hardening behavior of the phase. As shown, the stress partitioning for the γ phase are larger than those for the α phase, and the stress partitioning for the α phase is almost the same between S32101 and 329J4L. In terms of the γ phase, the stress partitioning of S32101 was larger than that of 329J4L. That is, the γ phase of S32101 showed the largest stress partitioning. One of the reasons for the higher work hardening of the γ phase in S32101 is associated with nitrogen. As seen in Table 1, the nitrogen content of the γ phase in S32101 is higher than that in 329J4L. Ojima et al.¹⁰⁾ investigated the work hardening behavior in high nitrogen-bearing austenitic steel (HNS) by in situ neutron diffraction and in situ electron backscattering diffraction. The high work hardening of HNS ascribed the larger stress partitioning among [hkl] family grains at small strain and high dislocation density at a larger strain. This is coincident with the present result as seen in Fig. 9 but the larger stress partitioning was maintained up to the maximum load in the S32101. The stress partitioning between γ and α phases also became larger in the latter stage of deformation in the S32101 as seen in Fig. 7. Both stress partitioning among the grains and that between phases seems to be associated with the large work hardening in the S32101. On the other hand, stacking fault energy is decreased with an increase in the nitrogen content.²⁰⁾ The stacking fault energy of the γ phase in S32101 is expected to be lower than that for the 329J4L even if the effects of other alloying content on stacking fault energy are considered.^21,22) Transmission electron microscope (TEM) micrographs of the γ and α phases in S32101 and 329J4L deformed with a true strain of 0.2, see Fig. 10, which is the latter stage of plastic deformation, as shown in Fig. 4. Dislocation-cell structures were observed to have developed during tensile deformation in the α phase in both steels and the γ phase in 329J4L. However, stacking faults were observed in some γ grains of S32101. This means that the stacking fault energy of the γ phase in S32101 is lower than that of 329J4L. The lower the stacking fault energy, the higher the degree of work hardening becomes.^20,21) Therefore, a larger work hardening rate for the γ phase in S32101, as discussed in Figs. 7 and 9, was also indicated by the TEM observations.

Fig. 8.

Lattice strain as a function of applied stress in the S32101 and SUS329J4L steels obtained by in situ neutron diffraction experiments.

Fig. 9.

Lattice strain as a function of applied stress in the austenite (a) and ferrite (b) phases in the S32101 and SUS329J4L steels obtained by in situ neutron diffraction experiments.

Fig. 10.

TEM micrographs of γ and α phases in the S32101 (a) and the SUS329J4L (b) deformed with true strain of 0.2.

3.4. Stress–strain Relationships of Duplex Stainless Steels Discussed by a Micromechanic Modeling

On the basis of the experimental results shown in Figs. 6,7,8,9,10, a schematic illustration of stress–strain curves for duplex stainless steel and the γ and α phases for S32101 (a) and 329J4L (b) are summarized in Fig. 11. As shown in Fig. 7, in the case of S32101, the γ phase plays an important role as the hard phase during the latter stage of plastic deformation, and the work hardening behavior of the γ phase is pronounced. In terms of the stress partitioning between γ and α phases, it is larger with an increase in true strain because the work hardening of the γ phase is very large, as shown in Figs. 7 and 9. In the 329J4L, α is the hard phase and the stress partitioning between γ and α phases is slightly decreased with true strain.

Fig. 11.

Schematic illustrations for true stress–strain curves of γ and α phases and duplex stainless steels in the S32101 (a) and the SUS329J4L (b) summarized by the in situ neutron diffraction and white x-ray diffraction experiments.

Next, the true stress–strain curves of the duplex stainless steels are calculated based on the schematic illustration of Fig. 11 to investigate the role of γ phase in the plastic deformation behavior of duplex stainless steels qualitatively. Here, we used the secant method proposed by Weng,²³⁾ which is based on micromechanic models of the Eshelby inclusion theory²⁴⁾ and the Mori–Tanka mean field concept.²⁵⁾ The details of the secant method are described elsewhere.^23,26) In the calculations,²⁶⁾ the true stress–strain curves of the α and γ phases were calculated by the following Swift equation,²⁷⁾   

$$\sigma =a{(b+\mathit{\text{ε}})}^N,$$(4)

where σ is true stress, ε true plastic strain, a, b and N constants, respectively. The constants of a, b, and N in Eq. (4) used in the calculations are shown in Table 4. Figure 12 shows the calculated σ–ε curves up until the plastic instability of S32101 (a) and 329J4L (b) (black dashed lines) and those for the γ and α phases used in the secant method (red and blue dashed lines). Here, the constants of a, b, and N for the γ and α phases were determined to be the same work hardening rate at the same true strain on the basis of Fig. 11 and the volume fraction of the γ phase was assumed to be 50%. The calculated σ–ε curves shown in black dashed lines were almost the same in Figs. 12(a) and 12(b). The calculated σ–ε curve in the case where the γ volume fraction is assumed to be 40%, which is equal to the experimental result of the 329J4L, is shown by a black solid line in Fig. 12(b). Flow stress is increased and uniform elongation is decreased by increasing the volume fraction of the α phase, which is the hard phase in 329J4L. The differences in the stress–strain curve and the tensile properties between S32101 and 329J4L, as shown in Fig. 4, can be duplicated by the calculated σ–ε curves. This means the volume fraction is an important factor in the tensile deformation behavior of duplex stainless steels. The black solid line in Fig. 12(a) shows the calculated σ–ε curve in the case where the work hardening of the γ phase becomes larger, as shown by the red solid line, which is based on the stress partitioning of the γ phase obtained in Fig. 10. The work hardening of duplex stainless steel also increased by changing the work hardening behavior of the γ phase. Furthermore, both the true stress and true strain at the plastic instability were found to be larger. The calculated results on the basis of Fig. 11 explain the stress–strain curves and tensile properties of S32101 and 329J4L steels qualitatively. By summarizing the obtained results from the viewpoints of the deformability of the γ and α phases, the work hardening behavior of S32101, which is related to the better combination of tensile strength and uniform elongation, can be explained by the increased work hardening of the γ phase and its volume fraction.

Table 4. Constants of a, b and N in the Swift equation used in the calculations of stress–strain curves in Fig. 12.

	phase	a	b	N	In Fig. 12
(a) S32101	γ	1950	0.002	0.4	Red dashed line
		2300	0.002	0.45	Red solid line
	α	1184	0.002	0.092	Blue dashed line
(b) 329J4L	γ	1850	0.002	0.5	Red dashed line
	α	1282	0.002	0.083	Blue dashed line

Fig. 12.

Calculated true stress–strain curves of the S32101 (a) and the SUS329J4L (b) by using the Weng secant method and the flow curves of γ and α phases used in the calculations. Here, the effect of work hardening of γ phase (Vγ) on true stress-strain curve and that of volume fraction of γ phase on flow curve were investigated in (a) and (b), respectively.

4. Conclusions

In this study, static tensile tests, in situ neutron diffraction and white x-ray diffraction experiments were performed at room temperature to investigate the tensile deformation behavior of lean duplex stainless steels from the aspects of plastic deformability among phases or grains. The following conclusions were obtained:

(1) In the static tensile tests, S32101 steel exhibited a larger uniform elongation and a better tensile strength–uniform elongation balance than SUS329J4L steel. This is because a higher work hardening rate can be maintained up until high strains in the S32101.

(2) In the synchrotron radiation white x-ray diffraction experiments, the hard phase was changed from the α phase to γ one during tensile deformation in S32101. This led to a larger stress partitioning between the phases at the latter stage of deformation. From the experimental results of in situ neutron diffraction, it was found that the stress partitioning of the γ phase in the S32101 was the largest among the present results.

(3) The larger uniform elongation of S32101 was associated with macroscopic work hardening behavior. The larger work hardening rate of S32101 is explained by the in situ neutron and white x-ray diffraction experiments, in which stress partitioning between the γ and α phases, work hardening of the γ phase and γ volume fraction, play an important role. This was also verified by the calculations of true stress–strain curve by using a micromechanic modeling.

Acknowledgments

The authors are grateful to Professor R. Ueji of Kagawa University (now at Joining and Welding Research Institute, Osaka University) and Dr. Stefanus Harjo of Japan Atomic Energy Agency for their helps and valuable discussions. This work has been supported by the Inter-University Program for the Joint Use of JAEA Facilities. The synchrotron radiation experiments were performed at JAEA beamline in SPring-8 (Proposal No. 2011A-E31).

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