Effect of Temperature on Stress–Strain Curve in SUS316L Metastable Austenitic Stainless Steel Studied by In Situ Neutron Diffraction Experiments

Abstract

In situ neutron diffraction experiments during tensile deformation were conducted to investigate the effect of temperature on the tensile properties of JIS-SUS316L steel from the phase stresses of austenite (γ) and ferrite (α) phases and the kinetics of deformation-induced martensitic transformation (DIMT). The 0.2% proof stress and tensile strength increased with decreasing deformation temperature, and the maximum uniform elongation was reached at 223 K. The temperature of the maximum uniform elongation in metastable austenitic stainless steels is related to the mechanical stability of γ and showed good correlation with the Ni equivalent. The estimated phase stress of γ at a given true strain increased with decreasing temperature; however, the temperature dependence of the twinning-induced plasticity effect of the γ phase was small. The phase stress of α was almost independent of temperature between 138 K and 223 K. The effect of temperature on the mechanical properties of the SUS316L steel was largely affected by the transformation-induced plasticity effect, which was related to the kinetics of DIMT according to in situ neutron diffraction experiments.

1. Introduction

The mechanical properties of metastable austenitic stainless steels can be improved by the transformation-induced plasticity (TRIP) effect.^1,2,3,4,5) The deformation-induced martensitic transformation (DIMT) behavior of austenite (γ), as represented by the change in the volume fraction of deformation-induced martensite with increasing applied strain, plays an important role in the TRIP effect. There have been many studies on the relationship between the TRIP effect and DIMT behavior from the viewpoints of chemical composition,^6,7) deformation temperature,^3,4,5,8,9) and strain rate.^5,8,10,11) These studies clarified various deformation behaviors as well as DIMT in γ. In some types of austenitic steels with relatively low stacking fault energy, deformation twins form in γ during plastic deformation.^12,13,14) Tsakiris et al.¹²⁾ discussed the mechanical behavior of a high-Mn austenite steel for which a wide range of temperatures at which both deformation twinning and DIMT was observed. Nakada et al.¹³⁾ studied the effect of deformation twins on DIMT behavior using cold-rolled and cold-drawn JIS-SUS316 steels. These findings indicate that the DIMT and deformation twinning behavior significantly depend on the deformation temperature.^12,14,15) It is important to clarify the quantitative relationships among the temperature dependence on mechanical properties, DIMT, and deformation twinning. Various studies on calculations of the stress–strain curve considering deformation twinning and DIMT have also been reported.^{14,15,16,17,18)} In a previous study,¹⁹⁾ Tsuchida investigated the description of the true stress (σ_t)–true strain (ε_t) curve of a JIS-SUS316L steel at room temperature on the basis of Remy’s study,¹⁶⁾ which proposed a quantitative model to estimate the strengthening by twin deformation. The σ_t–ε_t curve calculated using changes of the volume fraction of deformation twins during tensile deformation showed agreement with the measured one. However, it is difficult for this model to consider both the effects of deformation twinning and DIMT on the σ_t–ε_t curve. One possible way to extract the effect of DIMT from the simultaneous occurrence of these two phenomena is to perform in situ neutron diffraction experiments. In neutron diffraction experiments,^20,21,22,23) the phase stresses of γ and ferrite (α) phases and the DIMT behavior in metastable austenitic stainless steels during tensile deformation can be estimated. The phase stress, which is estimated using the phase strain, is the averaged stress of a constituent phase in multi-phase steels and is generally calculated using the difference in the lattice constant relative to that of a strain-free reference.^20,21,22) The temperature dependence of the mechanical properties of SUS316L steel can be discussed by analyzing the effect of the deformation temperature on the phase stress and the DIMT behavior.

Therefore, in this study, we conducted in situ neutron diffraction experiments during tensile deformation at various deformation temperatures in SUS316L steel. The temperature dependences of the mechanical properties of SUS316L steel were discussed from the viewpoints of the phase stresses of γ and α (deformation-induced martensite) and the DIMT behavior was analyzed using the in situ neutron diffraction experiments. Furthermore, the σ_t–ε_t curves at various temperatures were calculated using a micromechanic model incorporated the phase stresses and DIMT behavior. From the calculated σ_t–ε_t curves at various temperatures, the important factors affecting the temperature effect of the mechanical properties in SUS316L steel were discussed.

2. Experimental Procedures

In this study, a commercial JIS-SUS316L steel (0.015C, 0.56Si, 0.92Mn, 12.2Ni, 17.6Cr, 2.06Mo by mass%) with a thickness of 1.5 mm was used. The SUS316L steel sheet was annealed at 1373 K for 600 s followed by water quenching.¹⁹⁾ Figure 1 presents an optical micrograph of the SUS316L steel. The average austenite grain size was measured to be 12.2 μm using the conventional linear intercept method. From the sheet, tensile test specimens with a gage length of 25 mm and gage width of 5 mm were prepared, and static tensile tests were conducted with an initial strain rate of 3.3 × 10⁻⁴ s⁻¹ at various temperatures between 123 and 373 K using a gear-driven-type Instron machine.^4,5,23) The test temperature was controlled using a constant-temperature bath.

Fig. 1.

Optical micrograph of JIS-SUS316L steel.

Test samples deformed by various ε_t were prepared at 223 K and 123 K and were measured using electron back scattering diffraction (EBSD) in a scanning electron microscope (SEM). The EBSD pattern was analyzed as the possible phases were set as both face-centered cubic (FCC) and body-centered cubic (BCC), and the scanning pitch was 0.1 μm. The kernel average misorientation (KAM), which is the average misorientation angle of a given point with all its 1st. nearest neighbors, was calculated with the limitation of a maximum misorientation of 5°.

In situ neutron diffraction experiments during tensile deformation were performed at TAKUMI, the time-of-flight (TOF) neutron diffractometer for engineering sciences at the Materials and Life Science Experimental Facility (MLF) of J-PARC.^21,22,23) Tensile test specimens for the neutron diffraction experiments with a gage length of 25 mm, a gage width of 4 mm, and thickness of 1.5 mm were prepared. The test specimen was mounted horizontally in a loading machine equipped with TAKUMI at J-PARC, and diffraction patterns in the axial (tensile) and transverse (normal to tensile direction) directions were measured simultaneously using two 90°-scattering detector banks. The neutron diffraction patterns for both the axial and transverse directions were measured at detector areas integrated ±15° horizontally and ±15° vertically. Tensile deformations were conducted in a step-load-controlled manner with a 300-s hold for the elastic regions and in a continuous manner with an initial strain rate of 2 × 10^–5 s⁻¹ for the plastic regions. In this study, tensile deformation at various test temperatures between 138 and 373 K was performed using the 100 K cooling system for loading experiments of TAKUMI.

The lattice strains, phase strains, and volume fractions of the γ and α phases were calculated from the diffraction patterns. The microstructure of the present SUS316L steel before deformation is γ single phase and the γ phase is transformed to deformation-induced martensite during tensile deformation. In this study, the data for the α phase were considered to correspond to those of the deformation-induced martensite.²³⁾ The lattice strain (ε_hkl) is given by the difference in the lattice plane spacing relative to that of a strain-free reference using the following equation based on peak analyses:^20,21,22)   

$${\epsilon }_{hkl}=\frac{d_{hkl} - d_{hkl}^0}{d_{hkl}^0}.$$(1)

Here, d_hkl is the lattice spacing during tensile deformation and $d_{hkl}^0$ is the lattice spacing before deformation. $d_{hkl}^0$ of γ was obtained by peak analyses before deformation. Because α does not exist before deformation in the present SUS316L steel, it is difficult to analyze $d_{hkl}^0$ using peak analysis. In this study, $d_{hkl}^0$ of α was estimated using the following equation:²³⁾   

$$V_{\gamma }{\epsilon }_{\gamma }^{r-ph.}+V_{\alpha }{\epsilon }_{\alpha }^{r-ph.}=0,$$(2)

where ${\epsilon }_{\gamma }^{r-ph.}$ and ${\epsilon }_{\alpha }^{r-ph.}$ are the residual phase strains of γ and α, respectively, and V_γ and V_α are the volume fractions of γ and α, respectively. The Eq. (2) is based on the concept that a macroscopic stress can be calculated by the phase stress and the volume fraction for each phase and is generally used for the stress measurement by neutron diffraction method.²⁴⁾ The value of ${\epsilon }_{\alpha }^{r-ph.}$ is calculated using ${\epsilon }_{\gamma }^{r-ph.}$, V_γ, and V_α such that the lattice constant of α before deformation ($a_{\alpha }^0$) can be calculated from the calculated lattice constant of α. $d_{hkl}^0$ of α for each grain was calculated using $a_{\alpha }^0$. To obtain ${\epsilon }_{\gamma }^{r-ph.}$, V_γ, and V_α, the tensile deformation was performed in a stepwise manner at several ε_t, and the diffraction pattern for the unloaded specimen was measured during temporary stops for 600 s. The phase strains of γ and α were summarized as the average values of ε_hkl for each phase. Using the phase strains in the axial and transverse directions, the phase stress for the tensile direction (${\sigma }_{axial}^{ph.}$) of each phase was estimated using the following equation:^22,23)   

$${\sigma }_{axial}^{ph.}=\frac{E^{ph.}}{(1+{\nu }^{ph.})(1-2{\nu }^{ph.})}\left[(1-{\nu }^{ph.}){\epsilon }_{axial}^{ph.}+2{\nu }^{ph.}{\epsilon }_{transverse}^{ph.}\right].$$(3)

Here, E^ph. and ν^ph. are the elastic constant and Poisson’s ratio, respectively, for each phase.^20,21,25,26) V_γ and V_α were calculated using the integrated intensities obtained from in situ neutron diffraction experiments.^21,22)

3. Results and Discussion

3.1. Effect of Temperature on Tensile Properties of SUS316L Steel

Figure 2 presents nominal stress–nominal strain curves of the SUS316L steel at various temperatures. The 0.2% proof stress (0.2% PS) and tensile strength (TS) increased with decreasing temperature, and the maximum uniform elongation (U.El) was observed at 223 K. Figure 3 shows the 0.2% PS, TS, and U.El as functions of temperature in the SUS316L steel. The temperature dependence of the mechanical properties of the SUS304 steel (0.05C, 0.45Si, 0.98Mn, 8.2Ni, 18.2Cr, by mass%)⁴⁾ is also shown. The 0.2% PS and TS at a given temperature were smaller in the SUS316L steel than in the SUS304 steel. The U.El data of both steels showed significant curves, with the maximum value appearing at one temperature. The maximum U.El was also smaller in the SUS316L steel, and the temperature at the maximum U.El of the SUS316L steel was lower than that of the SUS304 steel. Figure 4 shows σ_t and work-hardening rate (dσ_t/dε_t) as functions of ε_t in the SUS316L steel at various temperatures. dσ_t/dε_t at a given ε_t increased with decreasing temperature, and that at 223 K decreased gradually during tensile deformation and indicated a larger value of dσ_t/dε_t until ε_t of about 0.5. dσ_t/dε_t below 173 K stopped decreasing and began to increase again.^4,5,14) Such temperature dependence of dσ_t/dε_t has been reported in various metastable austenitic stainless steels^2,3,4,5) and is closely related to their mechanical properties.

Fig. 2.

Nominal stress–nominal strain curves of SUS316L steel obtained from tensile tests at various temperatures. (Online version in color.)

Fig. 3.

0.2% proof stress, tensile strength, and uniform elongation as functions of deformation temperature in SUS316L steel. (Online version in color.)

Fig. 4.

True stress (σ_t) and work-hardening rate (dσ_t/dε_t) as functions of true strain at various temperatures in SUS316L steel. (Online version in color.)

Figures 5 and 6 present the SEM–EBSD inverse pole figure and phase maps at various ε_t at 223 K and 123 K. The volume fraction of deformation twins at a given ε_t increased with decreasing temperature. The black line denotes the high-angle grain boundaries with misorientation angles higher than 15°. Some of the high-angle boundaries with straight morphology in the γ phase could be distinguished as deformation twin boundaries. For SUS316L steel, the DIMT was observed to occur below 243 K according to X-ray diffraction experiments,¹⁹⁾ and V_α at a given ε_t increased with decreasing temperature. Most of the deformation-induced martensite was observed to form from the γ grains with deformation twins. It is well known that plastic deformation introduces the deformation substructure²⁷⁾ and that the local misorientation angle represented by KAM increases with increasing applied plastic strain.²⁸⁾ Figure 7 shows the average KAM of the γ phase as a function of ε_t at 123 K and 223 K. The average KAM at a given ε_t was larger at 123 K. This result was expected from Fig. 7, showing that the work hardening of the γ phase becomes slightly larger with decreasing temperature, indicating a slight change in the work hardening behavior of these two specimens.

Fig. 5.

Inverse pole figure maps and phase maps at various true strains (ε_t) deformed at 223 K. (Online version in color.)

Fig. 6.

Inverse pole figure maps and phase maps at various true strains (ε_t) deformed at 123 K. (Online version in color.)

Fig. 7.

Average KAM as a function of true strain at 123 and 223 K. (Online version in color.)

3.2. Relationship between Maximum Uniform Elongation and Mechanical Stability of Austenite in Various Austenitic Stainless Steels

In this section, the temperature dependence of the U.El is discussed. The U.El of metastable austenitic steels generally depends on the DIMT behavior of γ,^2,4,9) and the kinetics of DIMT is closely associated with the mechanical stability of γ. Figure 8 shows the U.El as a function of temperature in various metastable austenitic stainless steels.^4,5,23) The maximum U.El and the temperature at the maximum U.El (T_max) differ for the different steels. The uniform elongations in Fig. 8 are related to only the mechanical stability of γ phase but also the strength of α phase.²³⁾ For SUS316L steel, the maximum U.El was the smallest and T_max was the lowest of the five types of metastable austenitic steels. We discussed the relationship between T_max and the mechanical stability of γ. Here, Ni equivalent (Ni_eq.) and M_d30 were used as the index of mechanical stability of γ.^6,29,30) Figure 9 shows T_max as a function of Ni_eq. (a) and M_d30 (b). Ni_eq. and M_d30 were estimated using the following Sanga et al.²⁹⁾ and Masumura et al.’s³⁰⁾ equations, respectively.   

$$\begin{array}{l}\mathrm{N}{\mathrm{i}}_{eq.}(\%)=\mathrm{N}\mathrm{i}+12.93\mathrm{C}+1.11\mathrm{M}\mathrm{n}+0.72\mathrm{C}\mathrm{r}\\ \mathrm{\hspace{1em}\hspace{1em}\hspace{1em}\hspace{1em}}\mathrm{\hspace{0.17em} }\mathrm{\hspace{0.17em} }+0.88\mathrm{M}\mathrm{o}-0.27\mathrm{S}\mathrm{i}-0.24\mathrm{T}\mathrm{i}-0.07\mathrm{C}\mathrm{o}\\ \mathrm{\hspace{1em}\hspace{1em}\hspace{1em}\hspace{1em}}\mathrm{\hspace{0.17em} }\mathrm{\hspace{0.17em} }+0.19\mathrm{N}\mathrm{b}+0.53\mathrm{C}\mathrm{u}+0.90\mathrm{V}+0.70\mathrm{W}\\ \mathrm{\hspace{1em}\hspace{1em}\hspace{1em}\hspace{1em}}\mathrm{\hspace{0.17em} }\mathrm{\hspace{0.17em} }-0.69\mathrm{A}\mathrm{l}+7.55\mathrm{N}\end{array}$$(4)

  

$$\begin{array}{l}{\mathrm{M}}_{\mathrm{d}30}(\mathrm{K})=756-555\mathrm{C}-528\mathrm{N}-10.3\mathrm{S}\mathrm{i}\\ \mathrm{\hspace{1em}\hspace{1em}\hspace{1em}\hspace{1em}}\mathrm{\hspace{0.17em} }\mathrm{\hspace{0.17em} }-12.5\mathrm{M}\mathrm{n}-10.5\mathrm{C}\mathrm{r}-24.0\mathrm{N}\mathrm{i}-5.6\mathrm{M}\mathrm{o}\end{array}$$(5)

where the unit of those elements is mass%. T_max decreased with increasing Ni_eq. and decreasing M_d30. For the five specimens in Fig. 9, Ni_eq. appears to reasonably summarize T_max of various metastable austenitic steels. The DIMT conditions to obtain better elongation via the TRIP effect in previous studies^1,2,3,4,5,9) are that γ is mechanically stable in the early stage of deformation and the deformation-induced martensite is gradually transformed in the latter part of the deformation. Considering these qualitative conditions,^1,2) Ni_eq. is related to the index that the γ phase is stable in the early stage of tensile deformation. Thus, Ni_eq. appears to be highly correlated to T_max compared with M_d30. The relationship between T_max and the mechanical stability of γ in metastable austenitic steels must be understood for further investigations by increasing the number of specimens.

Fig. 8.

Uniform elongation as a function of temperature in various metastable austenitic stainless steels.^3,4,23) (Online version in color.)

Fig. 9.

Temperature at the maximum uniform elongation as a function of Ni equivalent²⁶⁾ (a) and M_d30²⁷⁾ (b) in various metastable austenitic stainless steels.^3,4,23) (Online version in color.)

3.3. In Situ Neutron Diffraction Experiments during Tensile Deformation at Various Temperatures

Figure 10 presents representative neutron diffraction patterns of the SUS316L steel deformed at 138 K in the axial direction. The diffraction patterns in the axial direction confirm the presence of γ as well as α and epsilon (ε) martensite phases in the present SUS316L steel. The presence of the ε martensite phase was not implied in the diffraction patterns in the transverse direction, and the analysis of the axial pattern derived a volume fraction of ε martensite not exceeding 5%. Thus, the tensile deformation behavior of SUS316L steel is discussed using the phase stresses of the γ and α phases. The α phase was not confirmed before tensile deformation at 123 K. Figure 11 shows the phase strains of the γ phase (a) and α phase (b) at various temperatures as a function of ε_t in the axial direction. The phase strain of γ at a given ε_t increased with decreasing deformation temperature, whereas the phase strain of α was almost independent of temperature. For the phase strain of α phase at 223 K, diffraction peaks of α were observed at ε_t > 0.4 and V_α in the axial direction was below 10%. The width of the error bar in phase strain corresponds to the half width of each peak appearing in the diffraction patterns. The relatively large width of the error bar with the phase strain of α at 223 K likely stems from the small volume fraction with small grain size.

Fig. 10.

Examples of diffraction patterns obtained from in situ neutron diffraction experiments during tensile deformation at 138 K in SUS316L steel.

Fig. 11.

Phase strains of austenite phase (a) and ferrite one (b) as a function of true strain in SUS316L steel at various temperatures in the axial direction. (Online version in color.)

Figure 12 shows the ${\sigma }_{axial}^{ph.}$–ε_t relationships for the γ phase (a) and α phase (b) at various temperatures estimated using Eq. (3). In Fig. 12, ${\sigma }_{axial}^{ph.}$ is plotted, and the lines are the calculated ${\sigma }_{axial}^{ph.}$–ε_t curves obtained using the following Ludwik equation:³¹⁾   

$${\sigma }_t={\sigma }_0+K{\epsilon }_t{}^n,$$(6)

where σ₀, K, and n are constants and the values of σ₀, K, and n used in Fig. 12 are summarized in Table 1. For ${\sigma }_{axial}^{ph.}$ for the γ phase, σ₀ increased with decreasing temperature and the temperature dependence of σ₀ was almost in agreement with that of 0.2% PS in Fig. 3. The effect of temperature on the values of K and n in Eq. (6) was independent of temperature. The deformation twins were not observed at 373 K; thus, the difference in the work-hardening component calculated using K ε_tⁿ in Eq. (6) between each temperature and 373 K are correlated to the TWIP effect.^16,19) As observed in Figs. 4 and 12(a) and Table 1, the TWIP effect of the γ phase due to deformation twins was approved, and the U.El was improved approximately 20% from 373 K to 296 K by the TWIP effect.¹⁹⁾ However, the effect of the temperature dependence on the TWIP effect was small. In terms of the α phase, the ${\sigma }_{axial}^{ph.}$–ε_t relationship and the values of σ₀, K, and n were almost independent of temperature. It should be noted that ${\sigma }_{axial}^{ph.}$–ε_t does not change greatly even though the deformation temperature brings different kinetics of DIMT. The ${\sigma }_{axial}^{ph.}$ of the α phase in the SUS316L steel was approximately 2–2.5 GPa, which is almost the same as that in the SUS304 steel.²³⁾ The robustness of the phase stress in deformation-induced martensite against deformation temperature is one of the most significant findings of this work, although the mechanism that provides this robustness has not yet been clarified. According to a study on the relationship between the strength and deformation temperature, the robustness indicates that the strength of the martensite phase is dominantly controlled by athermal stress. It is well known that grain-boundary strengthening provides an athermal stress,^32,33) which leads to the deductive consideration that the robustness implies that the fine grain size of martensite is one of the dominant strengthening mechanisms of the deformation-induced martensite in the examined steel. However, this discussion should be supported by the thermal activation process governing the strength of martensite. The temperature dependences on the ${\sigma }_{axial}^{ph.}$–ε_t relationships of the γ and α phases in the SUS316L steel were different from those of γ and α single phase steels. This seems to be associated with the interaction between γ and α phases and the changes of their volume fractions in metastable austenitic stainless steels. More experimental data is required in order to discuss the ${\sigma }_{axial}^{ph.}$–ε_t relationships for the γ and α phases in metastable austenitic steels.

Fig. 12.

Estimated phase stress of austenite (a) and ferrite (b) vs. true strain in SUS316L steel obtained from in situ neutron diffraction experiments at various temperatures. (Online version in color.)

Table 1. Values of σ₀, K and n in Eq. (6) for austenite and deformation-induced martensite at various temperatures.

T (K)	Austenite			Deformation-induced martensite
	σ0	K	n	σ0	K	n
138	340	1033.5	0.71	600	3090.3	0.762
173	280
223	240
296	200			–	–	–
373	180	1000.0	0.7

Figure 13 shows V_α as a function of ε_t at 138, 173, and 223 K. The solid lines are the calculated results obtained using the following equation proposed by Matsumura et al.:³⁴⁾   

$$V_{\alpha }=1-\frac{V_{\gamma 0}}{1+(k/q)V_{\gamma 0}{\epsilon }_t{}^q},$$(7)

where V_γ0 is the volume fraction of γ before deformation and k and q are constants.³⁴⁾ The values of k and q at various temperatures are summarized in Table 2. The value of V_α at a given ε_t increased with decreasing temperature. Equation (7) can be used to precisely calculate the DIMT behavior of SUS316L steel. However, from the values of k and q, it is difficult to discuss the effect of deformation temperature on the DIMT behavior. We also investigated the temperature dependence of the DIMT behavior in Fig. 13 using the following equation:^35,36)   

$$ln{V}_{\gamma 0}-ln{V}_{\gamma }=A+k^{\prime }\epsilon ,$$(8)

where A and k′ are constants and k′ is related to the mechanical stability of γ. For larger k′, the mechanical stability of γ decreases.³⁶⁾ The values of k′ and A at various temperatures are also provided in Table 2. As indicated by the dashed lines in Fig. 13, Eq. (8) was also used to calculate the DIMT behaviors, and the value of k′ increased with decreasing temperature.

Fig. 13.

Volume fraction of deformation-induced martensite as a function of true strain estimated using in situ neutron diffraction experiments in SUS316L steel at various temperatures. (Online version in color.)

Table 2. Constants in Eqs. (7) and (8) at various temperatures.

T (K)	Eq. (7)		Eq. (8)
	k	q	k′	A
138	15.3	1.80	2.95	−0.2
173	7.35	1.79	1.77	−0.13
223	14.2	3.11	1.47	−0.34

To investigate the estimated ${\sigma }_{axial}^{ph.}$–ε_t relationships for γ and α phases and the DIMT behavior at various temperatures in detail, σ_t–ε_t curves of SUS316L steel were calculated using the secant method.^18,34,37,38) The details of the calculation method for TRIP steels based on the secant method have been reported in previous studies.^18,34,37,38) In the calculations, Eq. (6) and the values in Table 1 were used for the calculations of σ_t–ε_t curves for γ and α phases, and Eq. (7)³⁴⁾ was used for descriptions of the DIMT behavior. Figure 14 compares the calculated σ_t–ε_t curves with the measured ones at 223 K (a), 173 K (b), and 138 K (c). The calculated σ_t–ε_t curves were roughly agreement with the measured ones. We discussed the effect of temperature on the σ_t–ε_t curve and mechanical properties in SUS316L steel from the ${\sigma }_{axial}^{ph.}$–ε_t relationships and DIMT behavior determined from in situ neutron diffraction experiments. In Fig. 14, the ${\sigma }_{axial}^{ph.}$–ε_t relationships of γ at each temperature are shown as dashed lines. As discussed in Fig. 12 and Table 1, the temperature dependency of the TWIP effect of the γ phase was small, and the ${\sigma }_{axial}^{ph.}$–ε_t relationship of the α phase was almost independent of temperature between 138 K and 223 K. The difference between σ_t and ${\sigma }_{axial}^{ph.}$ of γ in Fig. 14 corresponds to the TRIP effect resulting from DIMT and became larger with decreasing temperature. This is closely associated with the DIMT behavior because the phase stress of α is almost independent of temperature. Considering the phase stress and the DIMT behavior analyzed by in situ neutron diffraction experiments, the 0.2% PS of the SUS316L steel at various temperatures is in agreement with the phase stress of the γ phase, i.e., σ₀ in Eq. (6). V_α and the DIMT behavior have controlling influences on the effect of the deformation temperature on TS and U.El in SUS316L steel.^1,2,3,4,5) When the mechanical properties of SUS316L and SUS304 were compared,⁴⁾ TS at a given temperature was smaller in SUS316L and U.El of SUS316L was larger at temperatures below 273 K. This finding is associated with the mechanical stability of γ,^6,29,30) as discussed in Fig. 9. The difference of TS at the same temperature can be explained by the calculated σ_t from V_α and ${\sigma }_{axial}^{ph.}$ of α. Here, it was assumed that ${\sigma }_{axial}^{ph.}$ of α at each temperature was the same for SUS316L and SUS304. Judging from these results, the DIMT behavior due to the difference in mechanical stability of γ plays an important role in the TRIP effects of SUS316L and SUS304 steels. That is, the effect of temperature on the mechanical properties in the two steels can be roughly determined based on the DIMT behavior. This conclusion is based on evidence obtained from the present study that the phase stress of α was almost the same between the two steels and was almost independent of temperature. However, the effect of temperature on the TWIP effect in the γ phase and the difference in the ${\sigma }_{axial}^{ph.}$–ε_t relationship for γ between SUS316L and SUS304 were not clear. Thus, there is room for further discussion on these issues.

Fig. 14.

Calculated true stress–true strain curves and measured ones in SUS316L steel at 223 K (a), 173 K (b), and 138 K (c). Here, the estimated phase stress of the austenite phase is also shown as a dashed line. (Online version in color.)

4. Conclusions

In this study, in situ neutron diffraction experiments during tensile deformation were conducted at various temperatures to elucidate the effect of temperature on the tensile properties of JIS-SUS316L steel from the phase stresses of γ and α phases and the DIMT behavior. The main conclusions were as follows:

(1) In the tensile tests of the SUS316L steels between 123 and 373 K, the 0.2% proof stress and tensile strength increased with decreasing temperature and the uniform elongation reached a maximum at 223 K.

(2) The temperature at the maximum uniform elongation in metastable austenitic stainless steels is dependent on the mechanical stability of γ and shows good correlation with the Ni equivalent.

(3) The estimated phase stress of γ at a given ε_t increased with decreasing temperature; however, the effect of temperature on the work-hardening behavior of γ was small. The phase stress of α was almost independent of temperature between 138 K and 223 K. The calculated σ_t–ε_t curves using the estimated phase stresses and the DIMT behavior were roughly in agreement with the measured ones.

Acknowledgments

The authors are grateful to Mr. K. Tomita, an undergraduate student at University of Hyogo and Drs. K. Aizawa, S. Harjo, T. Kawasaki, and G. Wu of J-PARC center, Japan Atomic Energy Agency for their help and valuable discussions. The neutron diffraction experiments were performed at the J-PARC/MLF beamline in JAEA (Proposal No. 2018P0010).

References

• 1)   I.  Tamura: Tetsu-to-Hagané, 56 (1970), 429 (in Japanese).
• 2)   I.  Tamura: Met. Sci., 16 (1982), 245.
• 3)   V. F.  Zackay,  E. R.  Parker,  D.  Fahr and  R.  Busch: Trans. Am. Soc. Met., 60 (1967), 252.
• 4)   N.  Tsuchida,  Y.  Morimoto,  T.  Tonan,  Y.  Shibata,  K.  Fukaura and  R.  Ueji: ISIJ Int., 51 (2011), 124.
• 5)   N.  Tsuchida,  Y.  Yamaguchi,  Y.  Morimoto,  T.  Tonan,  Y.  Takagi and  R.  Ueji: ISIJ Int., 53 (2013), 1881.
• 6)   N.  Tsuchida,  S.  Kawabata,  K.  Fukaura and  R.  Ueji: J. Alloy. Compd., 577 (2013), Suppl., S525.
• 7)   K.  Nohara,  Y.  Ono and  N.  Ohashi: Tetsu-to-Hagané, 63 (1977), 772 (in Japanese).
• 8)   S. S.  Hecker,  M. G.  Stout,  K. P.  Staudhammer and  J. L.  Smith: Metall. Trans. A, 13 (1982), 619.
• 9)   T.  Angel: J. Iron Steel Inst., 177 (1954), 165.
• 10)   Y.  Takagi,  R.  Ueji,  T.  Mizuguchi and  N.  Tsuchida: Tetsu-to-Hagané, 97 (2011), 450 (in Japanese).
• 11)   K.  Spencer,  M.  Veron,  K.  Yu-Zhang and  J. D.  Embury: Mater. Sci. Technol., 25 (2009), 7.
• 12)   V.  Tsakiris and  D. V.  Edmonds: Mater. Sci. Eng. A, 273–275 (1999), 430.
• 13)   N.  Nakada,  H.  Ito,  Y.  Matsuoka,  T.  Tsuchiyama and  S.  Takaki: Acta Mater., 58 (2010), 895.
• 14)   T. S.  Byun,  N.  Hashimoto and  K.  Farrell: Acta Mater., 52 (2004), 3889.
• 15)   D. R.  Steinmetz,  T.  Japel,  B.  Wietbrock,  P.  Eisenlohr,  I.  Gutierrez-Urrutia,  A.  Saeed-Akbari,  T.  Hickel,  F.  Roters and  D.  Raabe: Acta Mater., 61 (2013), 494.
• 16)   L.  Remy: Acta Metall., 26 (1978), 443.
• 17)   O.  Bouaziz and  N.  Guelton: Mater. Sci. Eng. A, 319–321 (2001), 246.
• 18)   N.  Tsuchida and  Y.  Tomota: Mater. Sci. Eng. A, 285 (2000), 346.
• 19)   N.  Tsuchida: CAMP-ISIJ, 28 (2015), 488 (in Japanese).
• 20)   N.  Tsuchida,  T.  Tanaka and  Y.  Toji: ISIJ Int., 60 (2020), 1349. https://doi.org/10.2355/isijinternational.ISIJINT-2019-751
• 21)   S.  Harjo,  Y.  Tomota,  P.  Lukas,  D.  Neov,  M.  Vrana,  P.  Mikula and  M.  Ono: Acta Mater., 49 (2001), 2471.
• 22)   S.  Morooka,  O.  Umezawa,  S.  Harjo,  K.  Hasegawa and  Y.  Toji: Tetsu-to-Hagané, 98 (2012), 311 (in Japanese).
• 23)   N.  Tsuchida,  E.  Ishimaru and  M.  Kawa: ISIJ Int., 61 (2021), 556. https://doi.org/10.2355/isijinternational.ISIJINT-2020-253
• 24)   Y.  Akiniwa: J. Soc. Mater. Sci., Jpn., 54 (2005), 785 (in Japanese).
• 25)   H. J.  Frost and  M. F.  Ashby: Deformation Mechanism Maps, Pergamon Press, Oxford, UK, (1982), 20.
• 26)  The Japan Society of Mechanical Engineering: The Elastic Modulus for Iron and Specimen, Maruzen Press, Tokyo, (1996), 40.
• 27)   N.  Hansen,  X.  Huang and  D. A.  Hughes: Mater. Sci. Eng. A, 317 (2001), 3.
• 28)   R.  Ueji,  N.  Tsuchida,  K.  Harada,  K.  Takaki and  H.  Fujii: IOP Conf. Ser. Mater. Sci. Eng., 89 (2015), 012045. https://doi.org/10.1088/1757-899x/89/1/012045
• 29)   M.  Sanga,  N.  Yukawa and  T.  Ishikawa: J. Jpn. Soc. Technol. Plast., 41 (2000), 64 (in Japanese).
• 30)   T.  Masumura,  K.  Fujino,  T.  Tsuchiyama,  S.  Takaki and  K.  Kimura: Tetsu-to-Hagané, 105 (2019), 1163 (in Japanese).
• 31)   D. C.  Ludwigson: Metall. Trans., 2 (1971), 2825.
• 32)   D.  Jia,  K. T.  Ramesh and  E.  Ma: Acta Mater., 51 (2003), 3495.
• 33)   N.  Tsuchida,  K.  Fukaura,  K.  Nagai and  Y.  Tomota: ISIJ Int., 48 (2008), 1020.
• 34)   O.  Matsumura,  Y.  Sakuma and  H.  Takechi: Scr. Matall., 21 (1987), 1301.
• 35)   M. Y.  Sherif,  C.  Garcia Mateo,  T.  Sourmail and  H. K. D. H.  Bhadeshia: Mater. Sci. Technol., 20 (2004), 319.
• 36)   K.  Sugimoto,  M.  Misu,  M.  Kobayashi and  H.  Shirasawa: ISIJ Int., 33 (1993), 775.
• 37)   N.  Tsuchida,  Y.  Morimoto,  S.  Okamoto,  K.  Fukaura,  Y.  Harada and  R.  Ueji: J. Jpn. Inst. Met., 72 (2008), 769 (in Japanese).
• 38)   G. J.  Weng: J. Mech. Phys. Solids, 38 (1990), 419.