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  • 1. Materials Science and Engineering A317 (2001) 32 – 36 www.elsevier.com/locate/msea Microstructure and damage initiation in duplex stainless steels S. Bugat, J. Besson *, A.-F. Gourgues, F. N’Guyen, A. Pineau Laboratoire de Metallurgie Mecanique, UMR CNRS 7633, Centre des Materiaux, Ecole des Mines de Paris BP 87, 91003, E6ry Cedex, France ´ ´ ´ Abstract The damage nucleation of a duplex stainless steel is investigated. Electron Back Scatter Diffraction (EBSD) technique is used to correlate local phase morphology with crystallographic properties. In situ tensile tests are performed to characterize strain fields and to monitor sites of damage nucleation. These observations are correlated with crystallographic orientations and finite element calculations. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Duplex stainless steels; Micromechanical behavior; Crystal plasticity; Damage nucleation 1. Introduction This paper deals with the experimental characterization of the micromechanical behavior and damage nucleation of an aged duplex (austenite/ferrite) stainless steel. This steel is used as a model two-phase material, as the hardness of the ferritic phase may be increased by thermal aging. This also leads to significant decrease of ductility and fracture toughness. The investigated material was provided as a centrifugally cast pipe containing about 20% of ferrite. Its chemical composition is given in Table 1. This material has already been investigated in [1 – 3]. This pipe was aged at 400°C during 700 h. The material has a coarse microstructure consisting of large basaltic ferritic grains (d phase) formed during the early stages of solidification. Austenite grains (g phase) appear by solid phase transformation as cooling proceeds leading to an interconnected duplex structure. The final microstructure consists of laths having a width of about 10 mm as shown in Fig. 1a. Both phases are related by near Kurdjumov –Sachs (K – S) crystallographic relationships [4]: one {111}g parallel to one {110}d and one Ž110g parallel to one Ž111d in those close packed planes. * Corresponding author. Tel.: + 33-160-763037; fax: +33-160763150. E-mail address: besson@mat.ensmp.fr (J. Besson). Thermal aging causes ferrite to cleave when the material is deformed. As both phases are percolated, the cavities initiated from the cleavage cracks in the ferritic phase have to plastically grow into the austenitic phase, leading to the final fracture of the material. It has been shown that damage is anisotropic (due to the orientation of the cleavage cracks) and heterogeneously distributed [1,3]. Highly damaged zones can be found with the surrounding material remaining fully undamaged (Fig. 1b). In Fig. 1b, it can also be noticed that the phase morphology of the damaged area differs from the morphology of the surrounding material. This suggests that phase morphology, possibly related to local crystallography, plays an important role in the damage process. The aim of this paper is to investigate the correlation between the phase morphology and the crystallographic properties and to model damage initiation. In situ tensile tests were also performed in a Scanning Electron Microscope (SEM) to observe the cleavage crack initiation sites and to correlate them with the phase morphology (Section 3). 2. Correlation between phase morphology and crystallographic properties In order to correlate the local morphology of both phases to their crystallographic properties, comparisons between light micrographs and Electron Back Scatter Diffraction (EBSD) maps were used. 0921-5093/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 1 9 6 - 0
  • 2. S. Bugat et al. / Materials Science and Engineering A317 (2001) 32–36 33 Table 1 Chemical composition of the investigated steel (wt.%) C S P N Si Mn Ni Cr Mo Cu Co Nb+Ta Al Fe 0.036 0.008 0.021 0.051 1.06 0.89 9.70 21.25 2.50 0.16 0.05 0.1 0.02 Bal 2.1. Use of EBSD The EBSD technique [5,6] was used to determine the local crystallographic properties of both phases. It allowed us to acquire orientation maps for a plane sample, to identify grains and sub-grains, and to study morphology, orientation, and in the case of a multiphase material, the phase geometry. In this last case, specific correlation techniques are used to obtain the grains for one given phase. The back scatter device was mounted on a DSM982 Gemini SEM. More details can be found in [7]. Data were acquired every 55 mm in both directions. To rebuild the grains from raw data, a specific procedure was designed, so that, for each phase, the set of acquired data points could be dilated to fill the whole sample. This procedure was applied to two plane and three notched samples. An example of data processing is given in Fig. 2 and Fig. 3. A light micrograph of the ferrite and austenite laths (Fig. 2) is used to compare the local phase morphology with the EBSD results. Fig. 3a shows the acquired data points: austenite is represented by black dots and ferrite by white dots. Fig. 3b shows the reconstructed austenite grains; Fig. 3c shows the reconstructed ferrite grains of the same area. the corresponding micrograph. austenite layer in the light 2.3. Results for austenite In the case of austenite, the step size of the EBSD mapping (55 mm) is not negligible compared with the austenite grain size, at least in the plane of the sample of Fig. 3. Therefore, the grain definition parameters for the austenitic phase may have a strong influence on the final orientation and morphology obtained with data processing. The selected values correspond to those giving a stable analysis. The Kurdjumov–Sachs relationships were first checked using the inverse pole figure ([001] axis) of the austenite in the standard triangle of ferrite (Fig. 5a and b). Experimental results are compared with the theoretical pole figure. Each data point corresponds to a zone where one ferrite and one austenite grain overlap. These zones will be referred to as ‘bicrystals’ in the following paras. About 41 bicrystals were determined for the microstructure shown in Fig. 3. For each bicrystal, the misorientation angle between the nearest theoretical orientation of the g-grain with respect to the d-grain and the experimental one was determined. Results are shown in Fig. 5c. About 54% of the zones have a misorientation less than 5° and 78% a misorientation 2.2. Results for ferrite About 12 ferrite grains were obtained in the case of the sample shown in Fig. 2, so that the average grain area was determined to be about 2 mm2. The grains are highly textured, their [001] direction corresponds to the radial direction (R) of the pipe and the other axes ([100] and [010]) are weakly misoriented with respect to the longitudinal (L) and tangential (T) axes. This is due to the centrifugation process. However, two grains (marked by * in Fig. 3c) had a Ž111 direction parallel to the R direction of the tube. Within all grains, the misorientation did not exceed a few degrees. Comparison between light micrograph and EBSD map of reconstructed ferritic grains shows that the ferritic grains are surrounded by a thin continuous layer of austenite. This can be explained because ferrite boundaries are preferential sites for austenite nucleation during solid-phase transformation. This allows a correlation between the phase morphology and the ferritic grains. An example is given in Fig. 4, where a ferritic grain boundary obtained by EBSD is compared with Fig. 1. (a) Austenite (light gray) and ferrite (dark gray) laths, (b) Cleavage crack cluster in a highly damaged zone. Fig. 2. Optical micrograph of the sample.
  • 3. 34 S. Bugat et al. / Materials Science and Engineering A317 (2001) 32–36 3b and c, it can be seen that the austenitic grain boundaries do not systematically correspond to the ferritic grain boundaries. This is consistent with the fact that ferritic grains are only slightly misoriented: a given austenitic grain can easily grow in two different ferritic grains and keep the K–S relationship with both. 3. Characterization of strains and damage nucleation 3.1. In situ mechanical tests In order to study the behavior and the damage nucleation process, SEM in situ tensile tests were performed on plane samples. Two kinds of geometries were used; smooth and notched samples. A gold grid was vacuum deposited on one side of the specimens by means of a microelectrolithographic technique [8]. The grid step size was 38 mm. It allows to compute the local strain using image analysis. The macroscopic stress and strain were also monitored. 3.2. Quantification of strain Fig. 3. (a) Raw data indicating data points corresponding to the ferrite (white) and the austenite (black), (b) reconstructed austenite grains, (c) reconstructed ferrite grains. Thick lines represent grain boundaries (misorientation angle (q) 15°, and thin lines represent subgrains (5°B q B15°) Nine grids were deposited on the plane sample shown in Fig. 3. They did not cover the whole surface. Table 2 compares the stress–strain curve obtained for the in situ tensile test with results obtained on a standard tensile specimen (i.e. a much larger specimen). Both sets of data differ for average strains larger than 2.5%; this can be attributed to a scale effect as the ferrite grain size is of the same order of magnitude as the dimensions of the in situ sample. The 1024× 1024 pixels images were used to analyze the grid deformation. By comparison with the initial Fig. 4. Ferritic grain boundary obtained by EBSD (left) versus corresponding ferrite grain which is surrounded by an austenite layer (light micrograph) (right). less than 10°, showing a satisfactory agreement with the K –S relationship. This low misorientation suggests that the austenitic lattice is rotated due to plastic straining during cooling down after solid-phase transformation. This is consistent with the fact that within each austenitic grain some sub-grain boundaries can be found. The austenite grains have a complex, non-convex morphology and are highly intricate. Comparing Fig. Fig. 5. (a) Theoretical inverse pole figure of the austenite ([001] axis) in the reference axis of the ferrite, (b) experimental inverse pole figure of the austenite ([001] axis) in the reference axis of the ferrite, (c) histogram of the misorientation angle between the nearest theoretical orientation of the g-grain with respect to the d-grain and the experimental measurement.
  • 4. S. Bugat et al. / Materials Science and Engineering A317 (2001) 32–36 Table 2 Comparison for a given strain of the engineering stress obtained during in situ testing and the stress obtained on a macroscopic tensile bar E (%) 0.556 1.581 2.706 3.806 5.850 8.140 10.20 12.20 in situ (MPa) 348 427 473 497 537 570 593 613 macro (MPa) 350 436 490 531 575 610 640 660 35 The laths are obviously too small to be accounted for. A homogenization procedure for the bicrystal, adjusted on unit cell calculations, was developed, in which the mechanical behavior of each phase was modeled using single crystal plasticity [9]. In that case, the size of the representative volume element is (50 mm)3. This allows the computation of the local values of stresses and strains for each phase. Calculations were performed using 3D meshes and constitutive equations for FCC and BCC crystals [9]. In addition, the test geometry is neither plane stress nor plane strain. Preliminary results are presented here in the case of a notched specimen. A localization band, located at the interface between two bicrystals (Fig. 7a), was experimentally observed (Fig. 7b and c). This band appeared at the onset of plastic deformation. The FE simulation correctly represented this effect (Fig. 7d). It is limited to the early stages of deformation as small deformation behaviors are used. 3.4. Damage initiation Sites of damage initiation are shown (circles) in Fig. 7 for a notched sample and in Fig. 8 for a smooth tensile sample. Fig. 7 shows that cleavage cracks did Fig. 6. Average local Green –Lagrange strains and standard deviation versus macroscopic Green –Lagrange strain (tensile direction 1). grid, the local displacement field is evaluated. This field is then derived in order to obtain the local strains. The average value and the standard deviation (S.D.) of strains are also calculated. As the initial undeformed grid is not fully regular, an initial S.D. is measured ( 92.5%). The actual S.D. was corrected assuming that the initial fluctuations and the actual displacement field are uncorrelated. Results are shown in Fig. 6 as a function of the macroscopic strain. The average local strain in the tensile direction (E11) does not exactly match the macroscopic strain: this is due to strain heterogeneity, as the grids did not cover the whole sample. This is confirmed by the increase in the S.D. with strain. Due to the grid step size (38 mm), which is larger than the lath size, the S.D. is representative of the strain heterogeneity between the different bicrystals. 3.3. Modeling of the stress– strain beha6ior In order to simulate the in situ tests using the Finite Element (FE) method, the bicrystal had to be modeled. Fig. 7. Notched specimen; (a) EBSD map of austenite grains (only one ferrite grain was observed), (b) deformed grid, (c) experimental strain field, (d) computed strain field. The initial notch width was 1 mm.
  • 5. 36 S. Bugat et al. / Materials Science and Engineering A317 (2001) 32–36 predict the nucleation rates obtained experimentally by [1,2]. 4. Conclusions Fig. 8. (a) First observed damage zones, (b) damaged area with a low CSF exhibiting discontinuous slip, (c) undamaged area with a high CSF exhibiting continuous slip. (CSF, Schmid factor of the slip system common to g and d phases). not initiate in the highly deformed area, but close to the root of the notch. The macro-crack leading to the final fracture of the sample initiated in this area, where the elastic stress concentration factor is equal to 1.10. This seems to indicate that stress plays a more important role in damage than strain. However, it can be seen that both sides of the notch are not equivalent in terms of local deformations and damage initiation. The present modeling, which accounts for the morphology of the grains and the anisotropic plasticity of the phases is, therefore, necessary as an isotropic behavior would not give this effect. Two crack clusters were observed on the surface of the tensile smooth specimen. Ferrite and austenite lattices related by the K–S relationship share one common slip system. The Schmid factor corresponding to this system (CSF) was computed for different bicrystals. It leads to low values in the damaged zones (CSFB 0.2), and higher values for the fully undamaged areas (0.3BCSF B 0.5). This result is corroborated by the observation of slip traces in damaged and undamaged areas (Fig. 8). Damage is, therefore, initiated in regions where the strain incompatibility between both phases is important, thus generating high local stresses. In the case of the notched sample, the highly deformed region had a common Schmid factor equal to 0.4, whereas it was equal to 0.05 in the neighboring, less deformed, zone. Based on these observations, the origin of heterogeneous damage nucleation can be interpreted. A predictive model for damage nucleation should incorporate a random distribution of cleavage stresses and the values of the local stresses in the ferrite (obtained using the homogenization procedure). The aim of this model is to . In this study, the microstructure and damage process of an austenite/ferrite duplex stainless steel were investigated. Concerning the microstructure, it was shown that ferrite grains are textured; their average size is equal to 2 mm, lattice misorientation within a grain remains limited. Ferrite grains are surrounded by a thin continuous layer of austenite. Austenite grains have an irregular morphology. They contain sub-grain boundaries. In areas corresponding to constant austenite and ferrite orientations (bicrystals), both lattices are related by the K– S relationship. A single austenite grain can grow in several ferrite grains. In situ tensile tests were used to monitor the local strain fields and to detect the sites of damage initiation. Heterogeneity of deformation increases with increasing average strain. Damage preferentially initiates in areas where the common slip system of the bicrystal has a low Schmid factor. Acknowledgements This work was supported by Electricite de France. ´ References [1] P. Joly, A. Pineau, Defect Assessment in Components — Fundamentals and Applications, Mechanical Engineering Publications, London, 1991, pp. 381 – 414. [2] P. Joly, Y. Meyzaud, A. Pineau, in: J. Giovanola (Ed.), Advances in Fracture/Damage Models for the Analysis of Engineering Problems, ASME, New York, 1992, pp. 151 – 180. [3] L. Devillers-Guerville, J. Besson, A. Pineau, Nucl. Eng. Design 168 (1997) 211 – 225. [4] G. Kurdjumov, G. Sachs, Zeitschrift Physik. 64 (1930) 325 –343. [5] V. Randle, The Institute of Materials, London, 1992. [6] B.L. Adams, S.I. Wright, K. Kunze, Met. Trans. 24A (1993) 819 – 831. [7] S. Bugat, Technical report, ENSMP, 1998. [8] L. Allais, M. Bornert, T. Bretheau, D. Caldemaison, Acta Metall. Mater. 42 (11) (1994) 3865 – 3880. [9] S. Bugat, J. Besson, A. Pineau, Computational Mater. Sci. 16 (1999) 158 – 166.