Materials Science & Engineering A 580 (2013) 159–168

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Materials Science & Engineering A

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nchawla

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Temperature-dependent mechanical properties of an austenitic–ferriticstainless steel studied by in situ tensile loading in a scanningelectron microscope (SEM)

En-Yu Guo a,b, Ming-Yue Wang a, Tao Jing a, Nikhilesh Chawla b,n

a Key Laboratory for Advanced Materials Processing Technology, Department of Mechanical Engineering, Tsinghua University, Beijing 100084, Chinab Materials Science and Engineering, School for Engineering of Matter, Transport, and Energy, Arizona State University, Tempe, AZ 85287, USA

a r t i c l e i n f o

Article history:Received 3 January 2013Received in revised form15 April 2013Accepted 17 April 2013Available online 26 April 2013

Keywords:Dual phase stainless steelPlastic deformationScanning electron microscopyIn situ tension

93/$ - see front matter & 2013 Elsevier B.V. Ax.doi.org/10.1016/j.msea.2013.04.060

esponding author. Tel.: +1 480 965 2402; fax:ail addresses: jingtao@mail.tsinghua.edu.cn (T@asu.edu (N. Chawla).

a b s t r a c t

In situ tensile tests at various temperatures, ranging from 25 to 750 1C, were conducted on an austenitic–ferritic cast duplex stainless steel (CDSS) to investigate both the plastic deformation mechanisms and theeffect of temperature on mechanical properties. A continual reduction in the mechanical properties, suchas ultimate tensile strength (UTS) and yield strength (0.2% proof stress, s0.2), was found as thetemperature increased. Fractographic analysis demonstrated that tearing topography surface (TTS) wasmore likely to occur at elevated temperatures. In situ observations revealed that the plastic deformationoccurred within the soft austenite matrix at first and was followed by slip gliding in the ferrite phase asthe load increased. Voids tended to form at the ferrite–austenite interphase boundaries or around theinclusions and then merge and propagate in the austenite matrix. The present study also shows that theclustered distribution of the ferrite phase in the matrix can cause crack initiation easily at early stages ofdeformation.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Cast duplex stainless steels (CDSS) are widely used in majorcomponents such as primary coolant water pipes for pressurizedwater reactors (PWR) in nuclear power plants. The microstructureof CDSS consists of a duplex structure of γ-austenite matrix andresidual δ-ferrite. Due to the presence of the combined phases,CDSS has very good corrosion resistance and attractive mechanicalproperties [1,2]. Coolant water pipes made of CDSS in PWR aredesigned for a 40-year or even longer service life and they operatein a service temperature range of 288–327 1C [3].

The mechanical properties of two-phase materials, like duplexsteels, are dependent on the strength of the individual microcon-stituents [4,5]. Each phase in the material will have a differentresponse to the applied strain, leading to heterogeneous plasticdeformation in the dual-phase alloys. In ferritic-martensitic dualphases steels void initiation and growth can occur by martensitecracking, separation of adjacent martensite regions, or by decohe-sion at the ferrite–martensite interface [6,7]. Other studies havepointed out that the deformation behavior and the damagemechanisms for a particular dual phase stainless steel are also

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+1 480 727 9321.. Jing),

related to their chemical compositions, heat treatment history,inclusions, and morphology of the microstructure [8–12]. Foct andAkdut [13] found that a small change in the nitrogen content canincrease the yield stress of the austenitic phase, even change it frombeing the softer phase to becoming the harder phase. In situ analysisby X-ray diffraction [11,14,15] and by neutron diffraction [16] haverevealed that the mechanical properties of the duplex stainlesssteels depend strongly on micro- and macro-stress partitioningbetween austenite and ferrite phase-domains and grains as well asmorphological and crystallographic texture in ferrite and austenite,resulting in heterogeneous deformation behavior. As a consequence,in order to understand deformation and damage behavior of aparticular dual phase stainless steel, one needs to have an under-standing of constituent phases, complex deformation constraints,load partitioning between phases and grains, etc., during deforma-tion [14–18]. In addition, temperature is a very important factor tothe mechanical properties of materials. Several investigations ofaustenitic stainless steels have shown that the temperature-dependent fracture properties are weakened and that the dominantstrain hardening mechanisms change with temperature [19,20].

In situ observations using transmission electron microscopy(TEM) [21,22] or scanning electron microscopy (SEM) [23,24] havebeen conducted to study the behavior of stainless steels. But thework mostly focused on the initiation and continuous propagationof microcracks in single-phase austenitic stainless steels or on thebehavior of coatings on stainless steel under tensile loading.

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Studies using atomic force microscopy (AFM) were carried out tounderstand the interactive deformation behavior between theferrite and the austenite phases [25,26]. However, the effect oftemperature was not considered in those studies and the effect offerrite such as the distribution of the ferrite in the matrix needs tobe investigated in detail.

In our present study, we have conducted an in situ SEM study ofthe uniaxial tensile behavior of a Fe–Cr–Ni CDSS to understandmechanical properties of these steels both at room and elevatedtemperatures. An effort has been made to focus on the deforma-tion behavior by slip bands and interactions between austeniteand ferrite. The role of interphase boundaries and on damagenucleation and its development was studied and is discussed.Fracture behavior was also studied by investigation of fracturesurfaces.

2. Materials and experimental procedure

The materials were machined out of a centrifugally cast pipe forpressurized water reactor coolant system piping, with a thicknessof 105 mm. The pipes were heat treated at 1105710 1C for 4.5 hbefore water quenching to avoid precipitation of unwanted phases.The chemical composition of a Z3CN20-09M stainless steel is givenin Table 1. Fig. 1(a) shows an optical micrograph of the cross-section of the pipe after etching. An equal thickness of about20 mm from both the internal wall and outer wall was machinedaway. It also can be seen that solidification of the pipe resulted inradially orientated coarse columnar grains in the outer part and aninner zone with equiaxed grains.

Samples for tensile tests were prepared as follows. First, dog-bone-shaped tensile specimens were cut from the columnarregion (Fig. 1(a)) (about 1.5 mm thick) by using electro-dischargemachining followed by mechanical polishing. In this way, themicrostructure of the tensile specimen gage section was keptuniform. Then, the samples were further polished to a mirrorfinish with a final thickness of about 1.0 mm. The geometry of thefinal tensile sample had a gauge length of 3 mm and cross-sectional area of 1.2 mm�1.0 mm. The radius between the gagesection and the grip ends was 0.6 mm (see Fig. 2). Two to threetensile tests were performed on the tested samples, at eachtemperature. Before tensile testing, the microstructure of each

Table 1Chemical composition of a Z3CN20-09M cast duplex stainless steel (at%).

Elements C Cr Ni Mn Si S P FeConcentration 0.128 21.72 8.64 0.927 1.70 0.024 0.0389 Bal.

Fig. 1. (a) Optical micrograph and (b)

specimen was quantified by using an optical microscope (OLYM-PUS BX51M) after electro-etching using a solution mixture of100 mL H2O and 20 g KOH at an applied potential of 3–6 V. Thevolume fraction of each phase was determined by the pointcounting method which calculates the ratio of the pixels whichstand for a corresponding phase to the total pixels of the image byusing image analysis software (Image J, Bethesda. MD).

Tensile tests were performed on a servo-hydraulic testingsystem installed inside the chamber of a scanning electron micro-scope (Shimadzu Corporation, Japan) [27,28]. The experimentswere conducted at a constant displacement rate of1�10−3 mm s−1, at 20, 350, 550 and 750 1C. The schematic ofthe in situ testing stage is given in Fig. 2. For tests at elevatedtemperature, the specimen was resistively heated to the desiredtemperature. The temperature was measured by a thermocoupleattached to the sample's gage section and was maintained within1 1C of the desired testing temperature [27]. After a certain strainincrement, the test was interrupted and in situ observation of thedeformation behavior of the specimen surface was conducted.Then, the test process continued. The SEM process and the loadingprocess were conducted until final rupture of the specimen. Thecontinuous cross-head displacement of the mechanical testingmachine was recorded automatically during testing. Finally, thefracture surfaces of samples after tensile testing were examined ona scanning electron microscope (JEOL, JSM-6460LV) coupled withan energy dispersive spectroscopy (EDS) detector.

3. Results and discussion

3.1. Microstructural characterization and tensile stress–strainbehavior

The microstructures of all the samples present a morphology ofδ-ferrite dendrites in an austenite matrix, as shown in Fig. 1(b). Thevolume fraction of δ-ferrite phase was measured as 14.571.5% forall the samples. This was consistent for all specimens measured.Typical uniaxial tensile stress–strain curves of Z3CN20-09M stain-less steel along the tangential direction of the pipe at varioustemperatures are shown in Fig. 3. The strain values were calculatedfrom the displacement over the gage length. This is a reasonableapproximation for calculating the plastic strain of the material,since the elastic compliance of the load strain is much smaller thanthe total strain in the sample. The true stress–strain curves wereconverted from engineering strain–stress curves in order to revealthe hardening behavior, as shown in Fig. 3. It is seen that the flowcurve at 25 1C has a much longer hardening region than thosetested at elevated temperatures. Table 2 summarizes the tensile test

microstructure of the cast pipe.

Fig. 2. (a) Flow chart of the in situ tensile testing stage, and (b) the shape and dimensions of the specimens (unit: mm). The apparatus for tensile tests, as inset in (a),is also shown.

Fig. 3. Stress vs. strain plots for a Z3CN20-09M stainless steel tested at various temperatures: (a) engineering stress–strain curves; (b) true stress–strain curves convertedfrom (a). The load drops indicate the stress relaxation occurred when the experiment was paused and SEM photomicrographs were acquired.

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data for steels tested at various temperatures in terms of ultimatetensile strength, yield strength, elongation at fracture (εf), uniformelongation (εu), strain hardening exponent (n), and strength coeffi-cient (K). The last two parameters were obtained by fitting the datato the Hollomon equation s¼Kεn [29]. Engineering stresses andstrains are shown in Table 2. The results show that there is a drasticreduction in tensile properties when the temperature increasesfrom 25 to 350 1C. The minimum in variation of elongation isobserved within the temperature range of 350–750 1C. A slightly

reduced value of both n and K was also noted. The reduced ductilitycan be confirmed by the shallow dimples at the fractured surface ofsamples tested at high temperatures. The phenomenon of reductionin strength and ductility was also revealed in stainless steels [19,30].For example, Byun [30] showed that the strength of the annealed316LN austenitic stainless steel was decreased from 630 MPa at20 1C to 405 MPa at 400 1C and the ductility was decreased from78% at 20 1C to 53% at 400 1C. Similar trend is also found in 304annealed and 316L annealed austenitic stainless steels [19,30].

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3.2. Deformation mechanism by in situ SEM

SEM micrographs showing plastic deformation in the micro-structures at room temperature under uniaxial loading withvarious strains and stresses are shown in Fig. 4. The tensile axisis vertical in the figure. It was observed that thin straight slip linesoriginated first in the austenite matrix. The slip lines tended tostretch along the long interphase boundaries depending on thegrain orientation and shape of the ferrite phases (see arrows inFig. 4(a)). Then, the slip lines started to coarsen to form slip bandsand steps and new multiple slip lines and cross-slip lines wereformed gradually as the tensile load was increased. Slip bands witha larger distance in the ferrite phases, compared with those in theaustenite matrix, were also observed. Moreover, it can also beobserved that slip lines or slip steps stopped at the interphaseboundaries. In Fig. 4(b), one can find that one type of slip bands inferrite can originate from the deformation of austenite (marked asregions B and C) while the second one results from the bulkbehavior in ferrite (marked as region A). These deformationphenomena were also reported and were dependent on thecrystallographic relationship between the ferrite phase and aus-tenite phase [25,26]. A magnified SEM micrograph (Fig. 4(c))revealed that the slip lines can pass through the ferrite phaseboundaries with or without changing the stretching directions.Necking deformation was observed on the lateral surface of thesamples upon further increase in loading, as presented in Fig. 4(d),

Table 2Mechanical properties of a Z3CN20-09M stainless steel at various temperatures.

Temperatures (1C) UTS (MPa) s0.2 (MPa) εf (%) εu (%) n K

25 552722 252715 56.971.7 46.072.0 0.523 3.14350 368710 177715 39.673.0 27.870.6 0.492 2.95550 317724 132710 35.270.2 26.072.6 0.458 2.89750 249732 112712 36.871.0 23.871.4 0.423 2.83

Fig. 4. Plastic deformation of the microstructures of a Z3CN20-09M stainless steel und(b) 479 MPa—37.3%; (c) 479 MPa—37.3%; (d) 488 MPa—56.9%.

where significant elongation of ferrite phases was observed alongthe tensile axis. Specimen surfaces were seriously distorted andgrain boundaries were difficult to distinguish, at this point, asshown in Fig. 4(d).

Our observations show that dislocations begin to glide in thesofter austenite phase under low shear stress, which results inparallel slip because of a low rate of deformation. At the beginningof the tensile test, slip is observed in essentially only one directionin some grains indicating that only one slip system is dominant.As the strain increases, the slip lines coarsen and multiple sliplines are observed. Cross-slip becomes more predominant andinduces many tangles, homogenizing plastic deformation over thegage section of specimen. This results in a high uniform strain,fracture stress, and fracture strain. In addition locations of satu-rated slip bands crossing also results in crack initiation sites. Onespecimen tensile-tested with the strain ∼39% at 550 1C shows thisphenomenon, as shown in Fig. 5. Curves which are almost parallelto the slip bands, as depicted in Fig. 5(a), represent the crackpropagation path. Fig. 5(b) shows the locations of crack initiationwhere two slip bands interconnect each other.

After extensive plastic deformation, fracture of the dual phasesteel began by decohesion of the ferrite phase from the austenitematrix, as shown in Fig. 6. Extensive plastic deformation of theaustenite grain was observed around the separated ferrite grains.Crack initiation at the ferrite–austenite interphase boundaries is avery important mechanism and was observed in all tested sam-ples. Small cracks initiated at the interphase boundaries arearrested by the ductile austenite. As further deformation proceeds,the voids were observed to propagate in a direction perpendicularto the tensile direction by the tearing of surrounding austenitephase. Finally, the density of interphase cracks is higher atelevated temperatures than at room temperature.

In the dual phase alloy system, the second phase can play a partin many ways. In the present study, the initial plastic deformationalways occur in the austenite phase and is usually confined in localarea depending on the shape, volume fraction and distribution of

er uniaxial tensile loading at 25 1C (engineering stress–strain): (a) 321 MPa—7.1%;

Fig. 5. Effects of intercross slip bands on strain hardening and fracture behavior: (a) SEM micrograph at low magnification showing the cross-slip bands interconnecting inthe center line of the specimen surface; (b) initial fracture location as indicated in (a). The test was conducted at 550 1C, and the engineering strain was ∼ 39% when theimages were acquired.

Fig. 6. Cracks sites at the interphase boundaries at various temperatures and engineering strains: (a) 25 1C,; (b) 350 1C; (c) 550 1C, and (d) 750 1C. The images were takenafter tensile tests.

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the ferrite phases. Fig. 7(a) shows that slip bands crossing at anangle of 641 in the austenite matrix area were confined within anarea composed of ferrite in one of the samples tested at 350 1C.Strain was ∼50% when the images were taken. After yielding andwork hardening of the soft austenite phase, sufficient stress istransferred to the hard ferrite phase to cause deformation [31].It also results in the build-up of strain along the ferrite–austeniteinterphase boundaries. It becomes more obvious when the ferritephase is relatively large and the ferrite dendrites connect eachother. The fracture will initiate, by either the formation of voids atthe interphase boundaries due to the fact that internal stress builtup [32] in the ferrite–austenite boundary regions or cracking of thehard second phase particles such as inclusions or ferrite phases(this part of the fracture process is described in more detailbelow). Final fracture occurs by the propagation of cracks throughthe austenite matrix, resulting in a typical ductile dimple fracture,similarly to the results on ferrite–martensite dual phase steels [33].As shown in Fig. 7(b) and (c), a more serious distinct plasticdeformation in austenite near the fracture surface than that on the

other side of the ferrite can be seen, reflecting the effect of ferritephase on the fracture behavior.

The distribution of the ferrite phase also plays an importantrole on the fracture behavior. In one specimen tested at 350 1C, theferrite was found to be clustered in the austenite matrix, as shownin Fig. 8(a). Strain in Fig. 8(a) was only ∼14.7%. Crack initiationwhich was parallel to the slip bands in the austenite matrix wasfound at an early stage during tensile testing. In these areas,inclusions were found around some of the cracks but someindividual cracks were also found with no inclusions inside.Hardening would be faster in the ferrite/austenite region than inthe austenite domains, and heterogeneous deformation would beeasier in the austenitic local region. As the tensile processproceeds, the saturation of dislocation tangles can result inexhaustion of strain hardening capability of austenite morequickly. When the applied stress reaches the rupture strength ofaustenite, cracks initiate in those areas and result in final rupture.We believe that the distribution of the ferrite, together with theinclusions, causes the phenomenon as shown in Fig. 8.

Fig. 7. Effects of ferrite on obstructing the plastic deformation: (a) plastic deformation of austenite confined between the ferrite; (b) plastic deformation of austenite near thefractured surface; (c) magnified SEM micrograph as indicated in (b). The specimen was tested at 350 1C. Images were taken at engineering strain ∼50% in (a) and when thespecimen got fractured in (b).

Fig. 8. Effects of band-distributed ferrite phase on deformation behavior: (a) initial microstructure; (b) and (c) magnified micrographs of regions as indicated in (a), showingthe crack initiation in the austenite matrix. The tensile test was conducted at 350 1C and the total engineering strain was ∼14.7%. The marked arrow in (a) represents thetensile direction.

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Fig. 9. Effects of impurities: (a) fracture of impurity; (b) separation of impurity from the matrix; (c) cracks propagated along the impurities.

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In addition to ferritic and austenitic regions, large regularspherical or ellipsoidal impurities were observed in the austenitematrix, at the interphase boundaries and even within the ferritephase-domains. The composition of the typical impurities wasanalyzed by EDS, as shown in Fig. 9(a) and (b). Further analysisshows that most of the impurities are MnS, FeO or a complexcombination of a sulfur compound and oxide. The diameters of thelargest FeO and MnS particles were measured as around 11.3 μmand 6.5 μm, respectively. These impurities tended to split or breakoff as the matrix material elongated under tensile loading, asshown in Fig. 9(a) and (c). Microcracks can easily form adjacent tothese locations as a result of stress concentrations. The separatedregions, between the impurity particles and the austenite or the

ferrite phase and the microcracks resulting from the fracture of theparticles, are very weak and can serve as crack initiation sites.

Finally, some comments on the temperature dependency ofmechanical properties would appear to be in order. The tempera-ture dependency of fracture properties are associated with achange in the dominant strain hardening mechanism with testtemperature [19]. For the present material, phase transformationdue to plasticity is not likely to take place. Low stacking faultenergy (SFE) and extended dislocations are identified as being ofcrucial importance with regard to the deformation behavior ofaustenitic steels [34]. Low stacking fault energy (SFE) at lowtemperature results in the easy activation of twins and planarslips that control the deformation of austenite. Austenite has been

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shown to exhibit planar slip in austenitic stainless steel and duplexstainless steel as well during early plastic deformation[11,26,35,36]. As the test temperature increases, the critical stressfor dislocation glide decreases, while the critical stress formechanical twinning increases slowly. The increased SFE atelevated temperatures results in more cross-slipping, even atlow strain, so that the strain hardening by dislocation tangles issaturated at relatively early stages of plastic deformation [37].Thus, planar slip becomes more important than mechanical twin-ning and governs plastic flow at elevated temperatures [38].Besides, the hardening effect of the ferrite phase begins to weakenat elevated temperature.

3.3. Fractographic analysis

All the tensile-tested specimens showed an obvious neckingfeature before the failure, as an indication of the ductile fracturemechanism. Fig. 10 represents the typical SEM micrographs for thedeformation and fracture morphologies of the tensile-testedZ3CN20-09M cast duplex stainless steel after tensile testing.Depending on temperature, the tensile-fracture mode of theZ3CN20-09M alloy varied considerably. The fracture surface ofthe specimen after tensile testing at 25 1C and 350 1C, as shown inFig. 10(a) and (b), respectively, exhibited many large and deepelongated voids with non-metallic inclusions sitting at the bottom,due to the presence of nucleation and growth of cavities. Theapproximate average diameter of the deepest (grown) voids hasbeen measured to be of the order of 24 μm and 21 μm for speci-mens at 25 and 350 1C, respectively. For the specimens tested at25 1C, a relatively uniform distribution of fine dimples wasobserved. Different dimple sizes in ridges were also observed.Shear-type dimples around the large voids and along the sides ofthe specimen were observed suggesting that a substantial amountof plastic deformation occurred before complete separation of thefracture surface by localized shear. Although small dimples werealso observed in the specimen tested at 350 1C, the plasticityassociated with each dimple appeared to be substantially reduced,

Fig. 10. SEM fractographs of a Z3CN20-09M cast duplex stainless steel teste

considering the depth of individual dimples, compared to those ofdimples tested at 25 1C and the dimples were much elongated inthe shear direction. As the test temperature increased to 550 1Cand 750 1C, SEM images in Fig. 10(c) and (d) shows that thedimples are even more shallow. One might notice that there arewaves on the smooth regions as shown in Fig. 10(b) and (c), whichdemonstrate that plastic deformation took place on those areas.The fracture morphology of some regions for the samples tested athigh temperatures indicates that fracture occurred through micro-void nucleation and ductile tearing. A similar microplastic mode ofductile tearing can also be found in particle reinforced aluminummatrix composites [39], steels [40,41] and aluminum alloys [40],and is termed Tearing Topography Surface (TTS) [39]. It has beenproposed in TTS that microvoid nucleation occurs at very closelyspaced nuclei, and that strain localization prevents any significantamount of subsequent void growth, preventing the formation ofwell-developed voids. This is indeed the case for the samplestested at high temperatures. Although, the TTS was also observedin the samples tested at room temperature, it is more likely toappear at high temperatures in our case as ductility is decreasedwith increase in temperature, and therefore decreasing the prob-ability of development of microvoids.

3.4. Summary of deformation mechanisms and temperature effects

The schematic of deformation and fracture mechanism of thepresent steel is given in Fig. 11. Straight slip lines originate fromsoft austenite matrix and stretch along the ferrite phase. As thestress is increased, multiple slip lines form and the previous thinslip lines merge to form coarse slip bands. Due to the deformationin the austenite which causes large stress concentration at theaustenite/ferrite phase boundaries, slip lines start to form in theferrite grains. At the same time, the deformation caused by theferrite itself can also give rise to slip bands. Further deformationcan result in decohesion at the interphase boundaries and inclu-sions, as well, which finally leads to fracture of the bulk material.Fracture phenomenon in Figs. 5 and 7 are shown clearly in Fig. 11(d),

d at different temperatures: (a) 25 1C; (b) 350 1C; (c) 550 1C; (d) 750 1C.

Fig. 11. Schematic of deformation and fracture mechanism of austenitic–ferritic stainless steel: (a) original microstructure; (b) initial slip lines in austenite matrix; (c) sliplines coarsen to form slip bands and multiple direction slip lines emerge with increasing stress; (d) decohesion phenomenon in the interphase boundaries and cracksinitiation(in red curves); (e) crack expanding (in red curves). The tensile direction in (a) applies to all the images. (For interpretation of the references to color in this figurelegend, the reader is referred to the web version of this article.)

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and is denoted by regions B and A, respectively. It should bepointed out that the previous mentioned deformation behavior forthis type of material is similar at all temperatures within the rangefrom 20 to 750 1C according to our in situ SEM observations.The effects of temperature can be summarized as follows. Firstly,the mechanical properties of the present material are weakened.The material becomes soft as temperature increases which isconfirmed by the reduction in both strength and elongation. Butone should also notice that the reduced properties do not show alinear relationship with the increased temperature. Secondly,tearing topography surface seems to be more likely to occur atelevated temperatures than at room temperature as microvoidsare less able to develop due to lower ductility which is confirmedby the shallow dimples. At last, qualitatively, the slip deformationand decohesion phenomena are supposed to happen at all thetested temperatures, but to what degree the density of the sliplines or the decohesion sites is affected due to the temperaturechange is still under discussion. Therefore, it is expected that thequantitative analysis in the future study can result in advancedunderstanding related to the change of the mechanical propertiesand deformation mechanism of dual phase materials.

4. Conclusions

The results of our investigation on the in situ tensile behavior ofdual phase stainless steels can be summarized as follows:

(1)

Within the test temperature ranging from 25 to 750 1C, themechanical properties of the Z3CN20-09M stainless steel werereduce drastically from 25 to 350 1C while there was amoderate reduction in mechanical properties at the tempera-ture range from 350 to 750 1C.

(2)

Slip lines occur in the soft austenite matrix at first and arefollowed by the slip glide in the ferrite phase. Slip lines inferrite can be caused by deformation in austenite or the bulkdeformation of the ferrite itself.

(3)

Voids tend to form around the inclusions or at the ferrite/austenite interphase boundaries, prior to the austenite matrix,

and then, merge and propagate in a direction perpendicular tothe tensile direction by the tearing of surrounding austenite.

(4)

The fracture morphologies reveal that tearing topographysurface is more likely to occur in the samples tested at 350,550 and 750 1C than at room temperature as ductility isdecreased with increasing of test temperature.

(5)

The distribution of the hard phase, here referred to as theferrite phase, has a significant effect on the early crackinitiation in the soft austenite matrix.

Acknowledgments

This research is financially supported by the National Scienceand Technology Major Project of the Ministry of Science andTechnology of China under Project Nos. 2011ZX04014-052 and2012ZX04012011. The authors are also grateful for partial fundsupport from the laboratory fund of Tsinghua University.

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