High-Speed Erichsen Testing of Grain-Refined 301LN Austenitic Stainless Steel Processed by Double-Reversion Annealing

In this section, the initial microstructures prior to DRA are introduced to reveal the influence of the subsequent DRA process. The as-received structure shows a coarse grain structure with an average grain size of 16 µm, as shown in Figure 2(a). A small fraction of α′-martensite, ~ 4 pct, coexisted with and was distributed inside the austenitic grains. This martensite formation in the initial structure is attributed to skin-pass deformation conducted by the supplier, i.e., final surface rolling with a slight deformation of 5 to 7 pct. The material was subjected to cold-rolling deformation with a 55 pct reduction and subsequent reversion annealing at 690 °C for 60 seconds using the CIH line (see Figure 1). This single-reversion annealing (SRA) treatment promoted a fine-grained austenitic structure, as shown in Figure 2(b). The magnified view in Figure 2(c) shows that the grain structure achieved with SRA is inhomogeneous, since some regions of coarse grains of size 4 µm, highlighted by the yellow circles, are enhanced. The grain size distribution of the grain structure in Figure 2(b) is shown in Figure 2(d). Most of the grains in the SRA structure, ~ 75 pct, have a size ≥ 2 ~ µm. Consequently, the average grain size of the grain structure achieved by SRA processing is 2.7 µm.

The microstructural features of the microstructures resulting from DRA were analyzed by EBSD mapping. Figure 3 displays the microstructure of the 301LN steel reversion annealed at T750-0.1 by DRA. Ultrafine grains along with a small fraction of retained austenite grains are observed, as indicated by the yellow circles in Figure 3(a). The corresponding orientation map, i.e., inverse pole figure (IPF), shows that the structure achieved by DRA at T750-0.1 has a textured orientation related to the grain size regime, as shown in Figure 3(b). The elongated retained austenite grain regime exhibits a 〈101〉 crystallographic orientation (green color); however, the new fine-grain regime shows 〈111〉 and 〈001〉 crystallographic orientations. The high-magnification image of Figure 3(c) shows a homogeneous and equiaxed ultrafine-grained structure with sharp and clear high-angle boundaries (HAGBs, > 15 deg). The grain size distributions of the UFG structure achieved at T750-0.1 are shown in Figure 3(d). It shows that 75 pct of the new grains are smaller than 0.5 µm, and 20 pct of the grains are in the 0.5–1 µm range. The estimated average grain size is 0.48 µm. The volume fraction of martensite remaining in the DRA T750-0.1 structure, as measured using XRD, was approximately 3.6 pct.

However, at a longer DRA time (1 second) and higher reversion temperature (800 °C), the microstructure is characterized by fine-equiaxed grains along with some retained elongated grains, as shown in Figure 4(a). The corresponding orientation map exhibits a similar textured orientation related to the grain size regime since the elongated retained austenite grain regime exhibits a 〈101〉 crystallographic orientation, as shown in Figure 4(b). Grain structure variation with different grain size regimes is promoted, as shown by the high-magnification image in Figure 4(c). The grain size distributions of the structure obtained by DRA at T800-1 are shown in Figure 4(d). It shows that 63 pct of the new grains are smaller than 0.5 µm, and 25 pct of the grains are in the 0.5–1 µm range. These grain size distributions are slightly higher than those achieved at T750-0.1. Hence, the estimated average grain size of DRA T800-1 is 0.53 µm, slightly larger than that of the DRA T750-0.1 structure. Less nonreverted α′-martensite coexisted in the structure subjected to DRA, approximately 2 pct.

As the reversion temperature increases to 900 °C, the microstructure after DRA exhibits homogenous fine-equiaxed austenitic grains, as shown in Figure 5(a). A striking feature of this structure is the disappearance of elongated retained austenite grains due to recrystallization. The orientation map shows that most of the new grains have a predominantly 〈101〉 textured orientation, as shown in Figure 5(b). A magnified image of the structure after DRA shows an ultrafine-grained structure with few large grains, as shown in Figure 5(c). The grain size distributions of the structure achieved after DRA at T900-1 are shown in Figure 5(d). DRA T900-1 contains fewer ultrafine grains of sizes < 0.5 µm than DRA T750-0.1 and DRA T800-1, with approximately 55 pct of the new grains being smaller than 0.5 µm, and 30 pct of the grains being in the 0.5–1 µm range. Consequently, the average grain size of DRA T900-1 is 0.75 µm, slightly larger than that of the DRA T800-1 structure.

The texture of the structures processed by DRA at different temperatures and times was evaluated by EBSD analysis. The texture patterns of the structures after DRA are presented by the constant ϕ2 = 0 and 45 deg ODF sections for the austenitic matrix, as shown in Figure 6. The structure after DRA mainly comprises a strong brass component {011}〈112〉. In addition, two texture components, Goss{110}〈001〉 and Cu {112}〈112〉, are present. There is a drastic decrease in the overall intensity of the texture with increasing reversion annealing temperature to 900 °C. The texture of DRA T900-1 is nearly 1/2 of that of the previous structures after DRA. It has been reported that the brass texture component is the major texture pattern in low-SFE materials of austenitic structures, such as austenitic stainless steel and TWIP steels.^[48,49] The strong brass texture component of the DRA T750-0.1 and T800-1 structures appeared to be related to proficient shear banding.^[50]

The tensile properties of the structures subjected to DRA at room temperature were measured at a high strain rate of ~ 1.5 s⁻¹ corresponding to the high punching speed of 50 mm s⁻¹ applied in the Erichsen tests for the same structures. Figure 7(a) shows the engineering stress–strain curves of the CG 301LN and specimens subjected to DRA at different temperatures and times. The mechanical properties extracted from the flow curves, such as the yield strength R_0.2, tensile strength R_m, and total elongation E_tot, are listed in Table I. The DRA T800-1 and T900-1 structures exhibit R_0.2 values of 950 and 770 MPa, respectively, which are significantly higher than the R_0.2 of CG 301LN (≈ 420 MPa). The total elongations of both structures are perfectly adequate, 35 and 40 pct. These values are comparable with those of CG 301LN. However, the DRA T750-0.1 structure reveals a very low elongation of ~12 pct and displays discontinuous yielding. The value of the yield drop, i.e., the difference between the upper and lower yield stresses, is ~100 MPa. This discontinuous yielding of the T750-0.1 structure is attributed to the high dislocation density and the anchored dislocations due to the carbon–dislocation interactions, as explained by Cottrell–Bilby theory.^[51] The occurrence of discontinuous yielding in grain-refined steels is a common phenomenon. As discussed by Song et al.,^[52] discontinuous yielding has been frequently observed in UFG ASSs, while their coarse-grained counterparts exhibit continuous yielding. Thus, a fine submicron/micron grain size is an essential factor. As seen in Figure 7(a), the coarse-grained 301LN steel shows continuous yielding and prominent strain hardening. However, the strain hardening decreases in structures after DRA, and Lüders strain becomes evident at an annealing temperature of 800 °C, i.e., with refined grain size. The same trend is seen in the stress–strain curves reported in Reference 30.

At 800 °C, a slight discontinuous drop in the yield strength of ~10 MPa followed by a Lüders strain plateau without strain hardening is observed. Consequently, the strain hardening at the Lüders plateau strain reaches zero, as shown in Figure 7(b). The slightly discontinuous yielding in the DRA T800-1 structure is attributed to the weak carbon–dislocation interactions. At a higher DRA temperature of 900 °C, the discontinuous yielding was fully eliminated. This is attributed to interstitial C atoms being either absent due to carbide precipitate formation or weakly interacting with the recovered dislocations.^[53]

The temperature rise of the tested specimens was measured during high-strain rate tensile testing using a fixed thermocouple. Due to adiabatic heating, the temperatures increased to 126 and 136 °C at the end of the tensile test for DRA T900-1 and DRA T800-1, respectively. This will be further discussed in Section 4.

Figure 7(b) shows the corresponding strain-hardening rate (SHR) curves of the structures after DRA. The SHR of the CG structure, plotted in black, is higher than the SHR of the structures after DRA. The soft γ-matrix of CG 301LN is significantly hardened due to the formation of shear bands, i.e., highly strained regions with a high density of dislocations and strain-induced α′-martensite plates formed at shear bands. Hence, the SHR of the soft γ-matrix of CG 301LN is higher than the SHRs of the hard γ-matrix of the structures after DRA. In agreement, Hamada et al.^[33] found that UFG 301LN with a grain size ~ 0.75 µm exhibited a lower SHR than CG 301LN. Mandal et al.^[54] reported that the CG structure of 301L showed substantial strain hardening, whereas fine- and ultrafine-grained structures developed by reversion annealing at 750 °C for 1 to 2 minutes and 700 °C for 10 minutes hardly strain hardened. A striking feature of the SHR of DRA T750-0.1 is the considerable drop below zero. This drop in SHR to a negative value usually indicates the start of necking according to the criterion of necking.^[55,56] However, plastic instability associated with necking does not occur at this strain. This remarkable drop in SHR is attributed to discontinuous yielding, as shown in the corresponding flow curve in Figure 7(a). It is well reported that the discontinuous yielding of reversion-annealed structures induces a sharp decrease in SHR.^[57]

The deformation mechanism during tensile straining of the structures after DRA at a high strain rate of 1.5 s⁻¹ was studied using EBSD examinations. Figure 8 shows the EBSD phase maps of the deformed DRA T750-0.1, DRA T800-1, and DRA T900-1 structures. High-speed deformation is accompanied by α′-martensitic transformation. A striking feature of the high-speed tensile deformation of the structure after DRA is the increased degree of martensitic transformation and an ultrafine-grained structure, as shown in Figure 8(a). A significant fraction of \(\gamma \) martensite transforms to \({\alpha }^{^{\prime}}\) martensite, ~70 pct, in DRA T750-0.1. In the magnified view shown in Figure 8(b), the γ-ultrafine grains are fully transformed into α′-grains. Figure 8(c) shows the deformed microstructure of DRA T800-1. The fraction of α′-martensite decreases to 45 pct. The high-magnification image in Figure 8(d) shows that the α′-martensitic transformation is favorable in the γ-ultrafine grains of size ~500 nm, as indicated by the yellow circles, whereas the α′-martensitic transformation does not occur in the γ-grains of size 1 to 2 µm, as marked by the white circles.

The deformed DRA T900-1 microstructure shows a completely comparable trend in α′-martensitic transformation, which reveals a significant decrease in α′-martensite fraction to 25%, as shown in Figure 8(e). As observed in Figure 8(f), most of the reverted γ grains did not transform into α′-martensite during tensile straining at high speed.

The maximum punch stroke displacement before cracking, i.e., limiting dome height (LDH) or Erichsen index (EI), has traditionally been used as a stretch formability parameter that indicates the ability to manufacture a material with the desired geometry. It is well known that the EI and the applied force in Erichsen cupping testing strongly depend on the sheet thickness and punch diameter.^[58,59] In the present study, the sheet thickness and punch diameter were constant in all tests.

High-speed stretch formability tests of the structures after DRA were conducted at ambient temperature using the conventional Erichsen cupping test procedure. Figure 9(a) presents the punch stroke displacement (X) vs force (F) curves. At the early stage of deformation, the average forces applied to achieve a stroke displacement of 2 mm are 4 and 10 kN for the CG and DRA T800-1 structures, respectively. As deformation proceeds, the applied force increases gradually up to maximum values of 52 and 74 kN for those structures. Thus, the force required for straining is much higher for structures resulting from DRA than for CG 301LN structures. In agreement, Saray et al.^[60] reported that the maximum load at the EI significantly increases for UFG structures of an IF steel processed by equal channel angular pressing.

A striking feature of the deformation behavior of the structures resulting from DRA and CG 301LN under biaxial straining in the Erichsen tests at high punching speeds is the significant variation in the strain-hardening capacity, as indicated by plotting the first-order derivative of the punch force with respect to the displacement, dF/dX vs. X curves, as shown in Figure 9(b).

In the early stage of punching travel, the onset of deformation is governed by both Young’s modulus and the strength of the structure. Hence, the DRA T750-0.1 structure exhibits the highest dF/dX at the beginning of punch travel, as indicated by the arrow. Each dF/dX-X curve exhibits a minimum peak. This peak is related to the start of yielding and plastic deformation. CG 301L exhibits the lowest peak value at early displacement in accordance with it having the lowest R_0.2. The structures resulting from DRA display higher minimum peak values and higher displacements, which can be attributed to the superior R_0.2 of the structures resulting from DRA, as seen from the tensile properties in Figure 7(a). dF/dX increases with further increasing displacement up to a maximum. In this stage, the deformation characteristic is membrane stretching and straining, where the specimen sheet becomes thinner under the effect of biaxial tensile stresses acting on the dome wall.^[61] The effect of the refined grain structure obtained by DRA on the membrane thinning regime under biaxial deformation is obvious. For instance, the CG 301LN structure exhibits a long regime of 8 mm (shown by the blue arrows), whereas the DRA T800-1 structure displays a shorter regime of 4.6 mm (shown by the yellow arrows). This will affect microstructural evolution during Erichsen cupping testing, as discussed in Section 4. After further punching beyond the maximum peak, the force decreases with displacement due to crack initiation and propagation on the dome surface, promoting deformation localization. Subsequently, fracture occurred at the hemispherical dome, resulting in different EI values. This means that the fracture strains of the studied structures are different since the true strain (ε_T) at the fracture of the hemispherical dome is correlated with the EI according to Eq. [1]^[62]:where \({r}_{\text{P}}\) is the punch radius (10 mm in the present work). Hence, for CG 301LN and DRA T800-1 structures, EI = 12.5 and 10.6 mm, and the true failure strains are 0.44 and 0.37, respectively.

The EI values from the high-speed Erichsen tests (HSs) are plotted as a function of the R_0.2 of the structures resulting from DRA. In Figure 10, the EI values measured with low-speed tests (LSs), taken from Reference 37, are given for comparison. Although grain refinement in the structures resulting from DRA significantly improves the mechanical strength, their EI values are almost comparable with those of the coarse-grained structure. In addition, with increasing punch speed from quasi-static to high speed, the EIs of the studied structures are only slightly decreased. For instance, by promoting the grain refinement of DRA T800-1, R_0.2 increased by 140 pct, while the corresponding EI decreased by 10 pct at a low punching speed. Similarly, at a high punching speed, the R_0.2 of DRA T800-1 increased by 125 pct, and the EI decreased by 15 pct.

The EBSD image quality (IQ) map together with the corresponding phase map of the deformed CG structure during high-speed Erichsen testing is shown in Figure 11. From the variation in IQ in Figure 11(a), the formation of dislocation substructures within the strained regions inside the grains can be concluded. According to the literature, the most significant feature of the uniaxial deformation of CG 301LN is shear bands.^[2,63] Since the strain field of these bands is significantly high due to the high density of dislocations, martensitic transformation is induced in these bands. Hamada et al.^[64] found that deformation strain is localized at the grain boundaries of TWIP steels during biaxial deformation in Erichsen testing. The phase map in Figure 11(b) shows that the martensitic transformation was strongly activated under biaxial straining in Erichsen testing so that the austenitic structure of the CG 301LN steel was almost completely transformed to α′-martensite, approximately 98.8 pct.

Figure 12(a) shows the IQ map of DRA T800-1 after high-speed Erichsen deformation, revealing distinct local differences caused by existing martensitic and austenitic regions, as seen from the EBSD phase map (Figure 12(b)). Thus, there is ~ 42 pct martensite and 58 pct austenite in the strained structure. The martensite fraction is higher than the 33% measured after tensile testing, indicating that higher martensite fractions formed in biaxial straining than in uniaxial tensile straining, as reported earlier.^[43] On the other hand, this shows that the structure grain-refined by DRA is less prone to martensite transformation during the high-speed Erichsen test. This is attributed to the longer thinning deformation regime, 8 mm, of the CG structure under stretching and straining, whereas the DRA T800-1 structure exhibited a shorter thinning deformation regime of 4.6 mm (see Figure 9(b)).

The austenite and martensite contents of the Erichsen-tested specimens were also measured by XRD. The corresponding X-ray diffraction patterns are presented in Figure 13. The martensite and austenite phase fractions of the CG 301-HS (HS=high speed) specimen were 98.8 and 1.2 pct, respectively, whereas the corresponding phase fractions were revealed to be 45.2 pct martensite and 54.8 pct austenite in the DRA T800-1 (UFG-HS) specimen.

Furthermore, to quantify the dislocation density of Erichsen-tested CG-HS and UFG-HS specimens, the average lattice strain and crystallite size of both austenite and martensite phases were first determined from the average peak broadening of the (111), (200), (220), (311) and (222) reflections for austenite and the (101), (200), (211) and (202) reflections for martensite using Scherrer’s equation,^[46,65] and the dislocation density of the corresponding phases was estimated by the modified W–H method using the following equation^[47]: where D is the average crystallite size, ΔK is the full width at half maximum (FWHM), K = (2sinθ)/λ, λ is the wavelength (0.15406 nm for Cu-Kα), and θ is the diffraction angle. β = πρM ²b² / 2, where ρ is the dislocation density, M is a constant depending on both the effective outer cutoff radius of dislocations and the dislocation density (usually taken as 2^[66,67]), and b is the magnitude of the Burgers vector (0.254 nm and 0.249 nm for austenite and martensite, respectively^[68,69]). C_h00 is a constant corresponding to the elastic constants of C11, C12, and C44 based on the elastic anisotropy, expressed as A_i = 2C_44/(C₁₁ − C₁₂), H² = (h²k² +h²l² +k²l²)/(h² +k² +l²), where h, k, and l are Miller’s indices of each peak and q is a parameter indicating the dislocation character of the samples. The values of C11, C12, and C44 are 248, 110, and 120 for martensite, respectively, and 209, 133, and 121 GPa for austenite, respectively.^[70,71] To obtain the value of q, the linear plot of {ΔK² −(0.9/D)²}/K² vs H² for all the considered reflections was used. Afterward, the slope of the plot of ΔK² vs K²C_h00(1−qH²) for all the considered reflections was used to obtain the value of β. The values of the calculated lattice microstrain crystallite size, dislocation density and other constants are summarized in Table II.

The calculated microstrain of the Erichsen-tested UFG-HS specimen was found to be relatively higher than that in the CG-HS specimen. This is further corroborated by the calculated dislocation density. The accumulation of a relatively large amount of microstrain in the UFG-HS specimen yields a relatively high degree of strain-induced martensitic transformation from austenite. Note that the dislocation density of both CG-HS and UFG-HS specimens obtained after the high-speed Erichsen test is lower than that obtained after the low-speed Erichsen tests.^[43] This is naturally attributed to adiabatic heat evolution during the high-speed test, resulting in significant dislocation recovery.

SEM micrographs representing the free dome surfaces and fracture surfaces through the thickness of DRA T750-0.1 and DRA T800-1 samples after the Erichsen tests are shown in Figure 14. The surface cracks extend on the dome surface following the paths of the strain localization bands, as shown in Figures 14((a) and (b)). It is generally accepted that the intense deformation of grain-refined 301LN is accommodated at grain boundaries and shear bands formed during stretching deformation in Erichsen testing.^[72] Meanwhile, the orange peel effect associated with stretching deformation affects the surface roughness of the stretched cups.^[73] This is illustrated by the irregularity of the surfaces of the stretched cups. Based on the measurements of the surface roughness parameter R_Z of electropolished CG 301L and DRA T800-1 specimens, the stretched surface of CG 301LN exhibited a considerable R_Z value of 10 µm, whereas the very smooth the surface of DRA T800-1 was characterized by a lower R_Z ~ 3 µm.

The fracture surface through the thickness of DRA T750-0.1 after the Erichsen test is shown in Figure 14(c). Distinct zones of micro/nanoscale dimpled features and relatively flat quasi-cleavage facets with tiny tear ridges, marked by yellow circles, are the features of the fracture morphology. Normally, the fracture surface of the hard-martensitic phase typically exhibits a flat-shaped fracture morphology.^[74] The flat “cleavage-like” pattern is evidence of martensite in DRA T750-0.1. In contrast, the morphology of the fracture surface of DRA T800-1 after the Erichsen test is clearly ductile, displaying dimple features, as shown in Figure 14(d). The dimpled fracture surface and multiple secondary cracks are clearly visible in the SEM image shown in Figure 14(e) taken at higher magnification. The present observations reveal that secondary cracks form along the grain boundaries within the localized deformation bands.

The contours of Von Mises equivalent strain at the end of high-speed tensile tests (the failure point determined from experimental flow curves) are obtained by FEM analysis and are demonstrated in Figure 15. Necking is obvious in both specimens, leading to higher local strains in the middle.

The distributions of true strain along the gage length of the specimen in various instances are plotted in Figure 16. The strain in the CG 301LN specimen remained uniform up to true strains as high as ~ 0.25, although strain localization was observed at the end of the test (0.26 seconds from the beginning). However, localization initiated in the reversion-annealed specimen after ~0.15 s and severely increased 0.2 seconds from the beginning of the test.

The slightly different strain profiles of the two specimens can be explained by their strain-hardening capabilities. As shown in Figure 7(b), the CG specimen showed a high hardening rate of ~ 2000 MPa from the beginning of the test until a true strain of ~ 0.2 was reached. A lower hardening rate was observed for DRA T800-1, although it slightly increased with increasing deformation. On the other hand, previous work has shown that the extent of strain-induced martensitic transformation increases faster in the DRA specimen than in the CG counterpart for quasi-static deformation of the same steel.^[43] Thus, the lower strain-hardening rate encountered here in the high-speed deformation of the UFG material can be at least partly ascribed to its much higher initial yield stress. In other words, the very high lattice frictional stress in this ultrafine-grained specimen diminishes the influence of isotropic hardening. Furthermore, due to its higher yield stress, the specimen subjected to DRA exhibits an adiabatic temperature rise that is initially more pronounced than that exhibited in the CG material. This can greatly affect the SFE of the material and, hence, its mechanical behavior, as will be further discussed in this section.

It is well known that high-speed deformation of materials is an adiabatic process that increases the temperature of the material and, hence, influences its mechanical behavior. The temperature rise (ΔT) can be calculated by a balance of thermal and mechanical energies through the following general relationship: where Q is the generated heat, m is the mass, c is the specific heat capacity, W is the total deformation work applied to the material, and β is usually considered the part of the mechanical work that is transformed into heat (also known as the Taylor–Quinney coefficient). Dividing both sides of Eq. [3] by volume gives where ρ is the density of the material, σ and ε are the flow stress and strain, respectively, and \(\int \sigma \mathrm{d}\varepsilon \) is equal to the deformation work per unit volume.

The temperature contours for the high-rate tension of CG specimens and specimens subjected to DRA at the failure strain, which are calculated by the thermomechanical coupled field FEM analyses, are presented in Figure 17. Expectedly, a severely localized increase in temperature was seen in the necked region of both specimens, while outside the neck, a temperature increase of 70 °C to 80 °C was calculated.

Using FEM analyses, the average temperature and strain values over the whole gage length of the two samples were obtained and compared with the experimentally recorded data (Figure 18). Beyond certain values of strain, the experimental curves exceed the temperature values calculated by the FEM. This can be related to the heat released due to strain-induced martensitic transformation, which is an exothermic reaction in the investigated steel.^[75]

Hence, for the correct calculation of the temperature rise through Eq. [4], the enthalpy of this reaction (ΔH_r) should also be considered. In other words, in the case of metastable microstructures such as those studied in the present work, energy conservation for the adiabatic deformation of the material can be written as follows: where the coefficient \({\beta}^{\prime}\) is as follows:

Thus, to consider the influence of martensitic transformation, \({\beta}^{\prime}\) can be substituted for β in Eq. [4]. In the present work, the experimentally measured temperature increase (ΔT) was used along with the plastic work per unit volume \((\int \sigma {\text{d}}\varepsilon )\) of the gage zone (calculated by the FEM) to obtain \({\beta}^{\prime}\) via Eq. [4]. The result is presented in Figure 19.

In stable microstructures, less than 10 pct of mechanical energy is restored as metallurgical defects in the material, and the remaining energy arises as thermal energy so that β takes values of 0.9 to 1. However, for the present investigation, with increased deformation, excess heat is released due to phase transformation, and hence, the coefficient β′ increases to values higher than unity. This increase in c' for the microstructure of the specimen subjected to DRA is sharp at a true strain of ~ 0.1. However, the β′ of the CG microstructure shows a gradual increase, which is in agreement with the high and definitely more stable strain hardening of this microstructure, as shown previously in Figure 7(b). The sudden release of heat and temperature increases that occurs in the specimen subjected to DRA increases the SFE and stabilize austenite.^[76] This may be more pronounced in the center part of the specimen, which undergoes flow localization.