Effects of Cold Rolling and Strain-Induced Martensite

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Effects of Cold Rolling and Strain-Induced Martensite Formation in a SAF 2205 Duplex Stainless Steel MAR CO BR ED A, KATYA BR U N ELLI, F R AN CESCO GR AZZI, ANTONELLA SCH ER ILLO, and IR EN E CALLIAR I Duplex stainless steels (DSSs) are biphasic steels having a ferritic-austenitic microstructure that allows them to combine good mechanical and corrosion-resistance properties. However, these steels are sensitive to microstructural modifications, such as ferrite decomposition at high temperatures and the possibility of strain-induced martensite (SIM) formation from coldworked austenite, which can significantly alter their interesting features. In the present work, the effects of cold rolling on the developed microstructural features in a cold-rolled SAF 2205 DSS and the onset of martensitic transformation are discussed. The material was deformed at room temperature from 3 to 85 pct thickness reduction, and several characterization techniques (scanning and transmission electron microscopy, X-ray diffraction, hardness measurements, and time-of-flight-neutron diffraction) were employed in order to fully describe the microstructural behavior of the steel. Despite the low stacking fault energy of DSS austenite, which contributed to SIM formation, the steel was found to be more stable than other stainless steel grades, such as AISI 304L. Rolling textures were similar to those pertaining to single-phase materials, but the presence of the biphasic (Duplex) microstructure imposed deformation constraints that affected the developed microstructural features, owing to phases interactions. Moreover, even if an intensification of the strain field in austenite was revealed, retarded SIM transformation kinetics and lower martensite amounts with respect to AISI 304L were observed. DOI: 10.1007/s11661-014-2646-x The M inerals, Metals & Materials Society and ASM International 2014

I.

INTRO DUCTION

D U PLEX stainless steels (DSSs) are a category of high-alloyed steels characterized by a biphasic austenoferritic (c/a) microstructure obtained from a proper solution treatment after the forming operations. The presence of an equal volume fraction of the phases provides the best combination of mechanical and corrosion-resistance properties, making DSSs very interesting materials, especially for structural and special applications in aggressive environments.[1] However, owing to the presence of the metastable austenitic phase and to the instability of ferrite at high temperatures, these steels are sensitive to diffusive and diffusionless phase transformations. The eutectoidic decomposition of ferrite in the temperature range of 523 K to 1273 K (250 C to 1000 C) and its nitrogen-supersaturated condition are the main causes for precipitation of dangerous secondary phases,[1–5] limiting the employment of D SS at high

MAR CO BR EDA and KATYA BR U N ELLI, Post-Doc R esearchers, and IR ENE CALLIAR I, R esearcher, are with the Industrial Engineering D epartment (D II), U niversity of Padova, Via Marzolo 9 int. 4, 35131 Padova, Italy. F R AN CESCO G R AZZI, R esearcher, is with the Consiglio N azionale delle R icerche (CN R ), Istituto dei Sistemi Complessi, Via M adonna del Piano 10, Sesto F iorentino, 50019 Florence, Italy. Contact e-mail: [email protected] ANTONELLA SCHER ILLO, Instrument Scientist, is with the Science and Technology F acility Council, ISIS Neutron Source, D idcot OX11 0QX, U .K. Manuscript submitted July 23, 2014. M ETALLU R G ICAL AN D M ATER IALS TRAN SACTION S A

temperatures. F urther, the possibility of strain-induced martensite (SIM) formation from cold-worked austenite cannot be neglected if the phase is not adequately stabilized. Secondary phase precipitation from ferrite has been the subject of several studies,[1–5] but less attention has been paid to SIM formation from D SS austenite. Conversely, kinetics and mechanisms of SIM have been extensively investigated in metastable austenitic stainless steels (ASSs), in which it was found to occur at relatively low strains at a wide range of temperatures.[6–15] R ecently, the austeniteto-SIM transformation was studied in low-alloyed D SSs grades (lean D SSs), in order to quantitatively describe SIM formation [16,17] and improve the mechanical characteristics by inducing the onset of the transformation induced plasticity (TR IP) effect.[18,19] On the other hand, in the higher-alloyed D SS grades, austenite is more stable and SIM has been more or less directly observed.[20–22] Austenite stability depends on the amount of solubilized alloying elements, and the involved deformation mechanisms can be associated with its staking fault energy (SF E), which is related to the chemical composition of the phase at a fixed temperature;[6] in isothermal conditions, the tendency to SIM formation increases with the decrease of SF E.[8] At room temperature, the martensitic transformation is governed by strain-induced mechanisms; therefore, its onset can be primarily related to plastic strain rather than the acting stress.[9] H owever, an increase in stress triaxiality has been found to promote the extent of SIM, becoming a function of both plastic strain and stress state.[23,24]

Thus, SIM formation and the amount of transformed austenite not only depend on SF E, but must also be related to other parameters, such as local strain, strain rate, strain direction, temperature, and stress state, that can modify the phase response to plastic deformation. In this type of strain-induced mechanism, the nucleation sites are created during plastic deformation, and shear band intersections (stacking faults, mechanical twins, and hcp e-martensite) are the preferential active sites.[9,10] These intersections, promoted by low SF E, act as SIM embryos and the developed martensite is commonly formed via two possible transformation sequences, either c fi e fi SIM or c fi SIM . In this regard, although e-martensite plays a definite role in SIM nucleation, it is not necessary for this purpose, since SIM can nucleate independently, especially when e is not thermodynamically stable.[9] SIM can cause delayed cracking in deep drawn components[22] and can also affect the corrosion resistance of the steel, since the number of active anodic sites in the surface is increased. [25,26] In D SSs, austenite is alloyed differently from ASSs, owing to a lack of the strong c-promoting Ni, suggesting the possibility of SIM formation. However, the presence of the Duplex microstructure can play a not negligible role, modifying the mechanisms of deformation and SIM kinetics in this class of steels. In the present work, the microstructural effects of cold rolling in a 2205 DSS were investigated, and the onset of such transformation is discussed. II.

MATERIALS AND M ETH ODS

The material under study was a SAF 2205 DSS in the form of 10-mm-thick hot-rolled bar, solution annealed at 1323 K (1050 C) and water quenched. The chemical composition of the investigated DSS is reported in Table I, together with that of an AISI 304L ASS, examined to compare the reliability of the SF E calculations. The steel was cold rolled (CR ) at room temperature, involving thickness reductions in the range of 3 to 85 pct, and the various reduction grades were achieved through discrete deformation steps (0.5 mm maximum). The rolling process was free of constraints in the transverse direction, thus allowing the deformation of the material in all directions. As pointed out by other authors,[10,11,23,24] a change in deformation mode influences SIM formation, causing a stress-state dependence Table I.

2205 304L

of transformation kinetics. Therefore, an equivalent effective strain was considered, by measuring single strains in the three directions and combining them using the following relation (von Mises): pffi ffiq ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi ffi 2 eVM ¼ ðe1 e2 Þ2 þ ðe2 e3 Þ2 þ ðe3 e1 Þ2 : ½ 1 3 This criterion typically provides a better correlation between plastic flow and stress state[10] and gives a better description of the effective strain developed during the rolling process. Table II relates the various thickness reductions to the effective strains at which the samples were subjected. The as-received (AR ) and CR samples were prepared following the standard procedure for metallographic investigations (mechanical grinding and polishing), and two etching procedures were performed. A first set of samples was formerly subjected to electrochemical etching at 6 V in a 10 pct solution of oxalic acid and then reetched using a solution composed of 0.10 g of N a 2S2O 5, 5 mL of HCl, and 1 mL of H NO 3 in 60 mL H 2O (solution A), while a second set was chemically etched with the Beraha’s reagent (solution B). The 85 pct CR samples were more deeply investigated using a JEOL* *JEOL is a trademark of Japan Electon Optics Ltd., Tokyo.

JEM 200CX transmission electron microscope (TEM ), after electrochemically polishing the material at 248 K (–25 C) and 35 V using a solution composed of 5 pct perchloric acid, 25 pct glycerol, and 70 pct ethanol. The etched microstructures of AR and CR samples were observed using a Leica Cambridge Stereoscan 440 scanning electron microscope (SEM ; Cambridge Instruments Ltd., Cambridge, UK), operating in backscattered-electron mode (BSE) at 29 kV, whereas a Leica DMR light optical microscope (OM ; Leica Microsystems, Wetzlar, G ermany) was employed to investigate the material in the AR conditions after Beraha’s etching. Both SEM -BSE and OM micrographs were edited using image-analysis software for an estimation of the volumetric percentage of the observed phases. Hardness and microhardness tests were performed in order to follow the work-hardening evolution of the phases caused by plastic deformation and to obtain the overall effect of the deformation on the material in terms

Chemical Compositions of SAF 2205 DSS and AISI 304L ASS (Weight P ercent)

C

Si

Mn

Cr

Ni

Mo

P

S

N

0.030 0.033

0.56 0.53

0.89 1.59

22.75 18.11

5.04 7.96

3.19 0.64

0.025 0.026

0.002 0.002

0.18 0.07

Table II.

Correlation Between Thickness Reduction and Effective Strain

Thickness R eduction

3 pct

5 pct

15 pct

35 pct

50 pct

65 pct

85 pct

Von Mises effective strain (eVM )

0.032

0.054

0.170

0.451

0.726

1.099

1.987

M ETALLU R GICAL AN D MATER IALS TR AN SACTIONS A

of global hardening. The measurements involved five indentations in each deformed sample (HV10) and five points for each phase (H V0.025); in both cases, an average value was then considered. The samples were analyzed by X-ray diffraction (XR D) on a Siemens D500 diffractometer (Siemens Corporation, Cherry H ill, NJ) using Cr Ka radiation in the angular range 2h = 60 to 140 deg (step size of 0.05 deg and 3 seconds of counting time for each step), in order to investigate phases arrangement in solution-annealed and cold-worked materials. The same instrument was employed to achieve the diffraction patterns for the SF E calculations (step size of 0.025 deg and 6 seconds of counting time for each step). XRD measurements were performed both on the rolling plane, by orienting the samples toward the rolling direction (RD), and on the plane orthogonal to the rolling one [normal direction (ND )]. M oreover, AR and 85 pct CR samples were subjected to texture analysis on a texture goniometer using a 241-point hexagonal grid for data acquisition. In this latter case, the specimen tilt was limited to the range of 0 to 70 deg, owing to instrumental geometrical limits, thus allowing the achievement of incomplete polar figures for both phases. The stacking fault probability of the austenitic phase was estimated considering the Warren–Averbach approach,[27] founded on c(111) and c(200) peak displacement after plastic deformation. The calculations were performed comparing the annealed (AR ) and the 3 pct cold-worked (3 pct CR) diffraction patterns,[28] using the following expression:[6] D2h ¼ ð2h200 2h111 Þ3pctCR ð2h200 2h111 ÞAR pffi ffi 45 3 1 ¼ tan h200 þ tan h111 a: p2 2

½ 2

The SF E of austenite was determined using the formula developed by Schramm and R eed, [6] generally employed for SF E calculations in ASSs: c¼

K111 x 0 Gð111Þa 0 A pffi ffi p 3

0:37

e250 111 : a

½ 3

In expression [2], the 2hs are given in degrees and were achieved by a Marquardt fit of the diffraction peaks, while a denotes the stacking fault probability, which increases with the density of the generated stacking faults. In Eq. [3], K111 x 0 is a proportionality constant found to be 6.6;[29] a 0 is the austenite lattice parameter; G(111) is the shear modulus for the c(111) planes (assumed to be 7.39 9 1010 N/mm2;[20]) A = 2C 44/(C 11 – C 12) is the Zener anisotropy (considering C 11 = 204.6 G Pa, C 12 = 137.7 GPa, and C 44 = 126.2 G Pa [30]); and he2i 111 is the mean-square microstrain, estimated from XRD peak broadening after cold deformation. As previously reported, for a verification of the SF E results, the same calculations were performed on an AISI 304L ASS (composition in Table I) whose SF E is extensively reported in the literature (~17 to 20 mJ/m2[6]), assuming a shear modulus G(111) = 6.5 9 1010 N /mm2 and a Zener anisotropy A = 3.43.[6] M ETALLU R G ICAL AN D M ATER IALS TRAN SACTION S A

The samples were also analyzed by means of the timeof-flight-neutron diffraction (ToF -ND ) technique in the Italian neutron experimental station (IN ES) located at the pulsed ISIS neutron source. The INES diffractometer [31] is equipped with nine banks of 3He detectors, lying on a circle of 1-m radius from the diffraction center and covering an angular range of 2h = 11 to 170 deg, going from backscattering (bank 1) to forward-scattering (bank 9) conditions. Exploiting the ToF method, each bank provides a diffraction pattern whose resolution depends on the angular position of the bank itself, according to the derivative of the ToF equation, thus leading to the highest resolution in the backscattering one. F or the measurements, samples were wrapped in aluminum tape, in order to avoid undesired reflections from the sample holder, and oriented to achieve the diffraction patterns pertaining to the RD . The ToF -ND data were refined using the R ietveld method and G SAS software,[32] and the peak profile function 4—a convolution between a pair of back-to-back exponentials and a pseudo-Voigt function—was taken into account.[33]

III.

RESULTS AND DISCUSSIO N

A. Stacking Fault Energy The susceptibility to SIM formation via nondiffusive phase transformation from austenite is strongly related to the stability of the austenitic phase and, therefore, to its SF E. The microstructural stability depends on the amount of the solubilized alloying elements in the phase, but also can be related to the microstructure itself, which is biphasic in the steel under analysis. Several empirical relations reported in the literature link the SF E value to the steel chemical composition, [6,27,34] but these were developed for ASS and cannot be considered reliable for DSS, since the amounts of alloying elements are substantially different. F urthermore, it was also highlighted that the validity of these formulas is restricted to a limited range of compositions for ASSs as well.[35] The calculated SF E and stacking fault probability of 2205 DSS and AISI 304L are reported in Table III. F or the austenitic grade, the results are strictly in agreement with the values reported in the literature, while the obtained value for the DSS approaches the 10 mJ/m 2 determined by R eick et al.[20] from direct measurements of extended dislocation nodes. As expected, owing to a lower amount of c-promoting elements than ASS, the results confirmed the 2205 DSS austenite as a low SF E phase, and since SIM formation in AISI 304L is well established, such transformation in the Duplex steel is then expected.

Table III.

2205 304L

XRD P arameters and SFEs of the Steels Under Analysis

106 he2i 111

103 a

103he2i 111/a

SF E (mJ/m 2)

8 9

18 9

0.4 1.0

10.9 17.2

B. M icrostructure The AR material exhibited a well-balanced microstructure, composed of almost equal volume fractions of the phases (Table IV), in which austenite (c) was dispersed inside a ferritic matrix (a). As can be seen in F igure 1(a), the austenitic grains were elongated toward the hot-R D and rather twinned, owing to recrystallization occurred in the solution treatment. CR modified the microstructure by thinning the grains in the orthogonal direction with respect to the rolling plane, causing the formation of a heavily banded microstructure and making DSS strongly anisotropic. On the rolling plane, the grains were increasingly deformed and considerably fragmented at the highest deformation degree (Figure 1(b)). As expected, AR ferrite was slightly harder than AR austenite (F igure 2 and Table IV), owing to the different extents of recovery and recrystallization stages that occurred during the solution treatment. Cold working caused a different hardening of the phases because of their crystal structures. In austenite, as an fcc phase, the strain-hardening rate was very high and a significant work-hardened state was reached at 15 pct deformation (F igure 2). Lower thickness reductions (3 pct CR ) allowed easy deformation of the phase, but not all the grains contributed to this first hardening, and traces of slip bands were only found within annealing twins. In the 5 pct CR samples, the dislocation density was increased and austenite hardness reached values comparable to those of ferrite at the same deformation degree; also in this case, twins were mainly interested in shear band intersections (F igure 3(a)). The presence of slip activity in twins can be ascribed to a their nearly random arrangement with respect to other parts of the same grain, since XR D measurements revealed strong initial textures in the AR material (Section III–C), thus resulting in an increased probability to be favorably oriented to the rolling force. Conversely, the nontwinned grains required further thickness reduction to allow the activation of the primary slip systems, and the low deformation range was not enough to induce substantial hardening. At 15 pct deformation, several slip systems were activated and a great increase in hardness was measured (F igure 2). In 15 pct CR samples, almost all portions of grains contributed to deformation and a wide number of shear band intersections were observed (F igure 3(b)), justifying the measured austenite hardening. Greater thickness reduction (35 to 85 pct) caused a progressive rise in hardness; shear band formation was increasingly pronounced and all grains contributed to the deformation process, denoted by the observed heavy

Table IV.

grain fragmentation at the highest thickness reduction (F igure 1(b)). On the other hand, owing to its bcc structure, the activation energy of the primary slip system in the ferritic matrix was higher than that in austenite, leading to a nearly regular HV increment in the entire deformation range (F igure 2). In the rolling plane, ferrite was elongated toward R D, but if compared to austenite, a greater thinning of the grains was observed at mediumhigh thickness reductions, due to the activation of a considerable number of slip systems that allowed greater deformations.

F ig. 1—2205 DSS microstructure in the rolling plane: (a) AR conditions (solution A) and (b) after 85 pct thickness reduction (solution B).

Volume Fractions, H ardnesses, and ToF P arameters of the AR Phases

Phase F raction (Vol Pct)

F errite (a) Austenite (c)

ToF R ietveld Parameters (Bank 1)

OM

ToF

HV0.025

Cell Dimension a (A)

Crystallite Size (A)

Microstrain S 400

Texture Index J

45.5 54.5

43.2 56.8

375 345

2.87906 3.60637

2425 4106

0.991 0.177

9.482 1.393

M ETALLU R GICAL AN D MATER IALS TR AN SACTIONS A

F ig. 2—Effect of cold rolling on hardness and microhardness, average values (maximum standard deviations: 5 H V10 and 3 H V0.025).

Since deformation of D SSs is dominated by austenite,[19] the behavior of the biphasic steel under study is modified with respect to single-phase materials, and the H V10 measurements were conditioned by the high workhardening rate of the fcc phase, especially in the low deformation range (Figure 2). The concurrent presence of large volumes of two different crystal structures altered the micromechanisms of deformation, because neighboring grains imposed restraints on each other. It was shown that in each DSS phase, deformation takes place homogeneously and the strain is transmitted through the interfaces by common slip systems.[36] However, as the number of dislocations in a boundary increases, the effectiveness of the boundary as an obstacle to the movement of dislocations also increases. In D SS, austenite is the ‘‘second phase’’ in the D uplex microstructure, in the form of elongated aggregates confined by grain boundaries, and its deformation is restricted to an area that is limited by the surrounding ferrite. In the case of low and medium thickness reductions (up to 50 pct), ferrite well accommodated the deformed austenitic grains, whereas with an increase in the applied strain, grain boundaries became harder as a result of dislocations piling up, constraining the deformation of austenite and causing its considerable grain fragmentation at high thickness reduction (Figure 1(b)). On the contrary, owing to its continuous character, the ferritic matrix suffers less such restrictions and the imposed strain can find easier deformation paths as thickness reduction proceeds, allowing the activation of a considerable number of slip systems at high deformations and causing a greater thinning of the phase. The previous investigation suggested the possibility of SIM nucleation for thickness reduction that overcame 5 to 15 pct, where the number of shear band intersections was increased. As a matter of fact, SEM observations on the Beraha’s etched samples revealed some features pointing out the possibility of such transformation. Above 15 pct thickness reduction, the reagent revealed some structural modifications inside the austenitic regions, which tended to expand as deformation M ETALLU R G ICAL AN D M ATER IALS TRAN SACTION S A

F ig. 3—Slip bands and intersection traces in austenite (solution A): (a) 5 pct CR sample and (b) 15 pct CR sample.

proceeded (F igure 4). H owever, cold working has been found to render a deformed area more anodic,[25,26] leading to possible etching differences inside the same phase. Therefore, the highlighted regions inside the c grains clearly gained a greater anodic potential than the surrounding austenite; unfortunately, this cannot be related strictly to martensite laths. In fact, among them, there could be some austenitic regions that did not transform into SIM but that gained a different anodic potential than other parts of the same grain. These features have been observed already in a 2101 lean DSS[16] and, even in this case, a direct attribution to SIM laths has not been possible. N evertheless, the micrographs were edited on image-analysis software, and the estimated SIM amounts are reported in Section III–F , together with neutron diffraction data. C. X-Ray Diffraction The diffractogram of the AR material along the R D was characterized by very narrow peaks, owing to the solution treatment, and height of both a(200) and c(220) reflections with respect to the a(110)/c(111) principal

degree can be correctly ascribed to texture development rather than to SIM formation, as instead reported in R eference 38. In both 85 pct CR phases, the huge peak broadening denoted the heavy grain fragmentation induced by cold working (F igure 5(a)). The use of XR D is helpful in revealing SIM in ASSs, but in D SSs, this analysis cannot clearly reveal if this transformation takes place, owing to the superimposition of ferrite and SIM peaks.[14] F urther, the strong initial orientation of the samples and the developed rolling textures hindered XR D pattern analysis, making phase quantification unreliable. However, if SIM occurred, the presence of huge c peaks in the 85 pct CR samples underlined that only a small portion of austenite transformed into martensite. D. Time-of-Flight-Neutron Diffraction

F ig. 4—Structural modifications in austenite (solution B): (a) 35 pct CR sample and (b) 65 pct CR sample.

peaks underlined a strong initial texture in both phases (F igure 5). AR ferrite was found to mainly possess a {100} cubic orientation, almost equally spread toward h110i and h100i directions (F igures 5 and 6), whereas the presence of a dominant {110}h100i Goss-type component was revealed in AR austenite (F igures 5 and 6). In austenite, the polar figure of c(111) reflections revealed the concurrent presence of a partial nearly random texture on the background, attributable to twin formation (refer to the normalized intensity with a value of 2 in F igure 6(b)). Cold deformation caused grain rearrangements, and after maximum thickness reduction, phases acquired strong rolling textures similar to those pertaining to single-phase materials. In 85 pct CR samples, ferrite mainly arranged toward the {100}h110i rotated cubic texture and a spread in the developed orientations was observed, in agreement with the previous work of R ys and Witkowska,[37] since cold rolling was performed in the same direction as the previous hot-rolling process. In austenite, the Goss texture was lost in the 85 pct CR samples and the {110}h112i Brass-type texture was developed (F igures 5 and 6). In this regard, the disappearance of the c(200) peak at the highest deformation

ToF -N D measurements can overcome the limits imposed by XR D, owing to the high penetrating power of neutrons and to the simultaneous acquisition on nine diffraction banks. As can be seen from Table IV, ToF ND phase fractions in the AR material were very close to those estimated by image analysis, reflecting the wellbalanced AR microstructure of the DSS under study. The changes in volumetric fraction of the phases were determined by multipattern refinement, considering the average value provided by GSAS after the analysis of the entire set of diffraction banks. F rom the refinements, an increment of the bcc phase volume fraction was revealed (Section III–F ), pointing out that the SIM transformation took place in 2205 DSS. A significant increasing in bcc phase was observed starting from 35 pct CR samples (eVM = 0.451), underlining that DSS austenite was more stable if compared to austenite in AISI 304L, where SIM formation at room temperature occurs from the early stages of deformation, even in plane strain conditions.[10,21] Cold deformation caused grain refinement, and the corresponding reduction can be estimated from the broadening of the ToF -N D peaks. The crystallite sizes of the AR phases—intended as the average phase volumes free of defects—are reported in Table IV; from a comparison with the observed microstructure, the obtained values were smaller than expected. This is due to an instrumental upper limit on the correctly detectable grain size (100 nm [31]), whereas smaller values can be reliably determined. H ence, in AR and 5 pct CR samples, the reported average size must be considered from a qualitative point of view, being the real value larger than the refined one. Both phases globally exhibited similar trends as the thickness reduction increased (F igure 7(a)), and low deformations (5 pct) caused the strongest reduction in defect-free areas. F urther thickness reductions progressively reduced the crystallite size in both phases, reaching a plateau after 35 pct deformation. The global trend is in agreement with several results reported in the literature,[39,40] even though both phases reached constant size values, which were shifted toward very low levels of strain if compared to other materials and were interested by heavier grain fragmentation. This strong and anticipated grain fragM ETALLU R GICAL AN D MATER IALS TR AN SACTIONS A

the solution treatment that restored the microstructure in terms of lattice strains; the absolute increase in microstrain indexes is reported in F igure 7(b). Austenite was subjected to an increasing microstrain field while, in ferrite, the measured strain nearly approached a saturation value at medium-high deformation degrees. These different behaviors again can be attributed to the D uplex microstructure and, therefore, to the constrained state of austenite, which determined a progressive rising of the strain field in the fcc phase. On the contrary, owing to its continuous character, the ferritic matrix was freer to deform, resulting in minor intensification of the developed stress field. R efinements also provided an estimation of the texture index J, defined as the ratio of the total volume of grains to the volume of the nonoriented grains, taking into account the spherical-harmonic approach.[33] The indexes in the AR phases are reported in Table IV, while their evolution during cold working is shown in F igure 7(c). In AR conditions, ferrite was strongly oriented, while austenite was within typical values for mediumtextured materials, as also revealed by XR D measurements in both phases. Comparing F igures 5(a) and 7(c), the reorientation of ferrite is clearly evident; cold working caused a rearrangement of the bcc grains, and starting from a (double) cubic texture in the AR material, ferrite passed through intermediate rotations reaching a heavily rolling-oriented state at 85 pct thickness reduction. The situation was different in austenite, for which the texture index was found to increase in the entire deformation range. F or the fcc phase, a change in the curve slope in F igure 7(c) can be noticed after 50 pct deformation, which can be ascribed to a saturation of the Brass-type texture, due to restraints imposed by the biphasic structure. E. TEM Investigation

F ig. 5—XR D patterns: (a) R D and (b) N D.

mentation can be attributed to compositional factors and to biphasic microstructure. The increased amount of substitutional alloying elements lowers the SF E of both phases, leading to a more effective resistance to dislocation motion and inhibiting extensive dislocation cross-slip; therefore, the formation of dislocation boundaries and the achievement of smaller mean cell sizes were favored. M oreover, the D uplex microstructure determined a reduction of the average dislocationfree paths, contributing to the observed rapid formation of cells and subgrains. Phase microstrains were evaluated by means of the S 400 parameter, which is usually employed in cubic crystals together with the S 220 coefficient.[33] In the case under study, the S 220 parameter was found to be negative, with no associated physical meanings; therefore, only the S 400 parameter was refined. The microstrain indexes of the AR phases are reported in Table IV (in microstrain units[33]) and, as expected, were found to be low, due to M ETALLU R G ICAL AN D M ATER IALS TRAN SACTION S A

N eutron diffraction revealed the possibility of the martensitic transformation in SAF 2205, but the estimated volume fractions were rather low, also after the maximum thickness reduction (Section III–F ). As expected, the deformed microstructures in the 85 pct CR samples were characterized by high dislocation density. In austenite, the highest thickness reduction determined the generation of a cellular structure composed of defect-free cells separated by dislocation walls (F igure 8(a)). The rapid cell formation was eased by the biphasic microstructure and by the more complex interaction mechanisms than in single-phase steels, causing the observed grain fragmentation and substantially lowering the mean defect-free cell size, as also revealed by neutrons. In the 85 pct CR samples, SIM was detected such as regions having a deep dark contrast inside the austenitic grains and surrounded by dislocation agglomerates but without assuming the characteristic lathlike shape (F igure 8(b)), as already reported when large deformations are involved.[10,12] As a matter of fact, the corresponding electron diffraction pattern clearly evidenced the presence of martensitic bcc structure within fcc austenitic grains (F igure 8(b)).

F ig. 6—XR D polar figures of AR phases: (a) a(100) and (b) c(111); 85 pct CR phases: (c) a(110) and (d) c(111).

F . Strain-Induced M artensite Quantification Martensitic transformation took place in the 2205 D SS austenite, and despite its low SF E, it occurred to a significantly lower extent than in metastable ASS. The estimated volume fraction of SIM from ToF -ND was around 12 pct at the highest deformation degree, underlining the greater stability at room temperature of D SS austenite compared to ASS. In Figure 9, the volume fraction of SIM obtained from ToF -ND refinements is reported as a function of the applied equivalent strain, together with that derived from SEM micrographs. As previously pointed out, the metallographic investigation suggested the possibility of such transformation, and the estimated volume fractions by image analysis were in very good agreement with the data from neutrons. At room temperature, DSS austenite behaved differently from ASS, and in the present study, SIM was found to form for thickness reductions greater than 15 pct. Although data in the literature are not available for deformations lower than 60 pct, the results

confirmed the investigation performed by Tavares et al.[21] with a similar D SS grade, since SIM volume fractions at medium-high thickness reductions were strictly comparable (F igure 9). SIM transformation is strongly temperature dependent, and all the data concerning ASS evidence the tendency to reach a saturation value, especially at low temperatures. [6] Combining the results of the present study with those of Tavares et al., [21] it is evident that in this steel, the transformation kinetics followed an exponential law rather than a sigmoidal-type curve (F igure 9); austenite was increasingly transformed within the considered deformation range and martensite saturation was not evidenced, even up to 96.6 pct thickness reductions. Therefore, even though in lean DSS martensite kinetics are similar to those in ASS,[17] the proposed regression models for SIM formation [13,23] do not seem to be applicable in this DSS grade, and a deeper analysis must be carried out, taking into account both the effect of temperature and the possible contribution of different microstructures.

M ETALLU R GICAL AN D MATER IALS TR AN SACTIONS A

F ig. 8—TEM micrographs in 85 pct CR austenite: (a) deformation structures and (b) SIM with corresponding electron diffraction pattern (zone axis [001]).

F ig. 7—ToF -N D refinements: (a) crystallite size, (b) microstrain, and (c) texture indexes.

IV.

CO NCLUSIO NS

In the present work, the effects of cold rolling and the onset of SIM in a 2205 D SS were investigated for thickness reductions in the range of 3 to 85 pct. In D SS, austenite is the second phase in the biphasic

M ETALLU R G ICAL AN D M ATER IALS TRAN SACTION S A

F ig. 9—SIM Tavares[21]).

quantifications

(present

work

and

results

from

microstructure, in the form of elongated aggregates almost possessing the same grain size of the ferritic matrix. As expected, cold working refined grain size, but an anticipated fragmentation was observed, considerably reducing the mean defect-free areas when low strains were involved. Besides the high content of substitutional alloying elements, which determines an increased resistance to dislocation motion, this behavior can also be attributed to the peculiar biphasic structure, allowing for a further reduction of dislocation-free paths and easing a rapid formation of cells and subgrains. Moreover, at a fixed thickness reduction, the strain fields at which the phases were subjected were found to be substantially different. The constrained state of D SS austenite caused an intensification in the strain field in the fcc phase, whereas ferrite was freer to deform, owing to its continuous matrix character. The results revealed a low SFE value of austenite in 2205 D SS (~11 mJ/m2), but the phase was found to be less metastable than austenite in AISI 304L. TEM investigation unambiguously revealed the occurrence of SIM as a consequence of cold rolling, and at a fixed equivalent strain, ToF -ND measurements highlighted that its amount was substantially lower than in the austenitic grade. In the Duplex steel, 12 pct austenite transformed into SIM at 85 pct thickness reduction, and besides the known effects of cold working on mechanical properties, it was found to not alter the pitting-corrosion resistance.[41] H owever, it should be noted that the initial austenite texture may have affected the martensitic transformation, inducing a greater stability in the low deformation range and causing a retardation in the observed transformation kinetics. Moreover, even SEM observations on the Beraha’s etched samples revealed the possibility of a martensitic transformation in DSS austenite, and the estimated volume fractions were in very good agreement with the data from neutrons, confirming the reliability of metallographic techniques for SIM detection in DSS. The phases interactions derived from the peculiar D uplex microstructure caused different responses to plastic deformation in 2205 DSS with respect to singlephase materials, owing to interphase mechanisms arising from the concurrent presence of large volumes of different crystal structures. SIM kinetics were different from ASS, and considering the work performed by Tavares et al.,[21] the investigation revealed that the models usually employed to describe the martensitic transformation [13,23] seemed not to be applicable in 2205 D SS, since the SIM amount was always found to be increasing in the entire deformation range.

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