Hot Working Characteristic of Superaustenitic Stainless Steel 254SMO

Transcription

1 Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), DOI /s Hot Working Characteristic of Superaustenitic Stainless Steel 254SMO Enxiang Pu Wenjie Zheng Jinzhong Xiang Zhigang Song Han Feng Yuliang Zhu Received: 30 August 2013 / Revised: 20 October 2013 / Published online: 4 April 2014 Ó The Chinese Society for Metals and Springer-Verlag Berlin Heidelberg 2014 Abstract Hot deformation characteristic of superaustenitic stainless steel 254SMO has been studied by isothermal compression testing in the temperature range of 950 1,200 C and strain rate range of s -1. The activation energy of 496 kj/mol was calculated by a hyperbolic-sine type equation over the entire range of strain rates and temperatures. In order to obtain optimum hot working conditions, processing maps consisting of power dissipation map and instability map were constructed at different strains. The power dissipation map exhibits two domains with relatively high efficiencies of power dissipation. The first domain occurs in the temperature range of 990 1,070 C and the strain rate range of s -1. Microstructure observation in this domain indicates the partial dynamic recrystallization (DRX) accompanied with precipitation of tetragonal sigma phase. The second domain occurs in the temperature range of 1,140 1,200 C and the strain rate range of s -1 with a peak efficiency of power dissipation of 39%, and in this domain, the microstructure observation reveals the full DRX. The instability map shows that flow instability occurs at the temperatures below 1,140 C and the strain rates above 0.1 s -1. KEY WORDS: Processing map; 254SMO; Superaustenitic stainless steel; Hot deformation 1 Introduction Superaustenitic stainless steel 254SMO with high contents of chromium, nickel, molybdenum, and nitrogen has high strength and excellent corrosion resistance. It is widely used in very aggressive environments such as chemical processing equipment, pharmaceutical plant, flue gas desulphurization, waste incineration plants, and sea water piping. The addition of many alloying elements results in the difficulty in the hot working process of the alloy. Available online at E. Pu J. Xiang School of Physics Science and Technology, Yunnan University, Kunming , China E. Pu W. Zheng (&) Z. Song H. Feng Y. Zhu Central Iron and Steel Research Institute, Beijing , China sxzwj@sohu.com However, with regard to the researches on 254SMO, much attention has been paid on the corrosion resistance, precipitation behavior, and the mechanical properties [1 3], while the information concerning hot working characteristic of this steel is limited because of its strategic fabrication technologies. Hot deformation is the most important method to achieve grain refinement of the austenitic stainless steel. Therefore, it is necessary to investigate the hot working behavior of this steel for providing theoretical basis for the design and optimization of the hot working process and for microstructural control. In the past two decades, processing map has been used to optimize the hot processing parameters and intrinsic workability of the metallic materials. Prasad and Seshacharyulu [4] evaluated hot deformation mechanism of titanium alloys, and pointed out that processing maps are very helpful in optimizing the process design without the need for expensive and time-consuming trial and error methods. Srinivasan and Prasad [5] constructed the

2 314 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), processing map of Inconel 718 superalloy and determined the instability regimes of hot deformation, and stated that strain does not have a significant influence on processing map. A large number of studies characterized the hot working behavior of the austenitic stainless steel using processing map technology, such as 304L, 316LN, and super304h austenitic heat-resistant stainless steel, and the best hot working conditions for these steels were obtained [6 8]. The studies on superaustenitic stainless steel 254SMO using processing map, however, remain lacking. In the present investigation, the processing map developed on the basis of the principles of dynamic material model (DMM) has been used to characterize the hot deformation mechanism and to optimize the hot working condition of 254SMO. According to DMM, the workpiece subjected to hot deformation is considered to be a nonlinear dissipator of power. The total input power (P) is transacted through a temperature rise (G content) and a microstructure change (J co-content). The power component that gets dissipated through microstructure changes may be normalized by assuming workpiece as an ideal linear dissipater [9, 10]. This will result in a dimensionless parameter called efficiency of power dissipation (g) through microstructural processes and is expressed as follows: g ¼ 2m m þ 1 ; ð1þ where m is the strain rate sensitivity of flow stress. The variation of g with deformation temperature and strain rate constitutes power dissipation map. The g parameter describes the constitutive response of the workpiece in terms of the various microstructural mechanisms that occur in a given regime of the temperature and strain rate [4]. The power dissipation map exhibits different domains, which may be directly correlated with specific microstructural mechanisms. Generally, dynamic recovery (DRV), dynamic recrystallization (DRX), and superplastic deformation are considered as safe hot deformation mechanism, while the damage mechanism includes the void formation, wedge cracking, and intercrystalline cracking [11]. A continuum criterion [12] obtained by utilizing the principle of maximum rate of entropy production in metallurgical system is employed to identify the regimes of flow instabilities and given by o ln½m=ðm þ 1ÞŠ nð _eþ ¼ þ m\0; ð2þ o ln _e where _e is the strain rate. The variation of n ( _e) with deformation temperature and strain rate constitutes instability map. Representative microstructural manifestations of flow instabilities are flow localization, adiabatic shear band, dynamic strain aging, mechanical twinning and kinking or flow rotations. The continuum criterion has been successfully applied to various materials, and the reasonability of predictions was demonstrated by microstructural examination of deformed specimens [13]. 2 Experimental Superaustenitic stainless steel 254SMO having the following composition (in wt%) was used in this investigation: C, P, S, 0.25 Si, 0.59 Mn, Cr, Ni, 0.70 Cu, 6.32 Mo, and 0.18 N, and Fe is the balance. The experimental steel was prepared by vacuum induction-melting method, and then the ingot was hot forged and rolled at 1,150 C into a sheet with a thickness of 12 mm. Cylindrical compression specimens with diameter of 8 mm and height of 15 mm were machined from the hot-rolled sheet. Compression tests were conducted on a Gleeble-3800 machine in the temperature range of 950 1,200 C at intervals of 50 C and the strain rate range of s -1 with intervals of an order of magnitude. Prior to the deformation tests, the specimens were heated by induction heating to 1,200 C at a heating rate of 30 C/s and were maintained at this temperature for 180 s. The specimens were subsequently cooled to the deformation temperature at a cooling rate of 10 C/s and maintained for 30 s before the compression tests. All specimens were deformed to a true strain of 1.2 and immediately water cooled to the room temperature. The load-stroke data obtained through hot compression were converted into true stress true strain values using standard equation, and true stress true strain curves were plotted. The flow stress data were obtained from the curves and used for constructing a processing map. The deformed specimens were sectioned parallel to the compression axis and prepared for microstructure analysis by conventional mechanical grounding and polishing. In order to reveal austenitic grain boundaries, the polished samples were etched in the potassium permanganate solution (1 g potassium permanganate and 10 ml concentrated sulfuric acid in 90 ml water) at 60 C for 2 h. The microstructural observation of deformed specimens was performed by light optical microscopy (OM) and scanning electron microscopy (SEM). The grain sizes were measured using the mean linear intercept method. The foils for transmission electron microscopy (TEM) were prepared first by grinding to a thickness of 40 lm and then thinned using a twin-jet electrolytic polishing apparatus, operated at 10 V in an electrolyte of 15% perchloric acid and methanol. The electrolyte was kept at -30 C. TEM examination was performed on an H-800 transmission electron microscopy.

3 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), Fig. 1 Microstructures of 254SMO: a as-received state, b after soaking at 1,200 C for 180 s, and subsequent water quenching 3 Results and Discussion 3.1 Initial Microstructure Figure 1 shows the initial microstructure (after hot rolling) and the microstructure before hot compression, respectively, both of which exhibit homogeneous and equiaxed grains. The average grain size of the former is about lm, while the latter is lm. As expected, the grains had grown considerably in the pre-heating stage before deformation. The driving force for the grain growth is the reduction in interface energy through the increase of the grain size [14]. 3.2 Stress Strain Behavior Typical true stress true strain curves of the superaustenitic stainless steel 254SMO deformed at temperatures of 1,000, 1,050, 1,100, and 1,150 C under different strain rates are shown in Fig. 2. The flow behavior indicated by Fig. 2 is in agreement with the result obtained by the previous investigation [15]. It is observed from the figure that there is an evident work hardening at the initial stage of the deformation, and the rate of work hardening decreases with strain increasing before flow stress reaches the peak. The maximum stress (peak stress) that is sensitive to both the strain rate and the deformation temperature increases with decreasing of temperature and increasing of strain rate. At lower temperatures (\1,050 C), the stress strain curves exhibit continuous flow softening no matter the strain rate; while at the temperatures above 1,050 C and strain rates below 10 s -1, the flow stresses increase to a peak followed by a slow strain softening and then a steady state is reached. Such a flow behavior is believed to be a typical phenomenon that the dynamic recrystallization (DRX) has taken place under hot deformation [16]. However, it is observed that at the temperature of 1,100 C and the strain rate of 0.01 s -1, the flow stress increases significantly with strain increasing when the true strain exceeds 0.8, as indicated in Fig. 2c. According to the model proposed by Prasad and Ravichandran [17], dynamic recrystallization may consist of two competing processes: generation of the interfaces (nucleation) and migration of the interfaces (growth). The rate of the interface generation will compete with the rate of migration, and the relative values of these two rates will determine the shape of stress strain curves. If the rate of the interface migration is larger than the rate of generation, then flow softening results, otherwise flow stress, will abnormally increase even steady state has been reached. Similarly, if a balance is established between these two rates, steady state can be observed. 3.3 Constitutive Equation of 254SMO The relationship among flow stress, strain rate, and deformation temperature can be represented by constitutive equations proposed by Sellars and Tegart [18]: Z ¼ _e expðq=rtþ; ð3þ Z ¼ _e expðq=rtþ ¼ A 1 r n 1 p ; ð4þ Z ¼ _e expðq=rtþ ¼ A 2 expðbr p Þ; ð5þ Z ¼ _e expðq=rtþ ¼ A½sinhðar p ÞŠ n ; ð6þ where Z is the Zener Hollomon parameter which is the temperature-compensated strain rate; r p is the peak stress; Q is the apparent activation energy of deformation; R is the gas constant (8.31 J/(mol K)); both n 1 and n are the stress exponents; and A 1, A 2, A, b, and a (&b/n 1 ) are the material constants. The power law description of the stress (Eq. 4) and the exponential law (Eq. 5) are preferred for relatively low stress and high stress, respectively, while the hyperbolic sine law (Eq. 6) is suitable for a wide range of the strain

4 316 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), (a) (b) (c) (d) Fig. 2 True stress true strain curves for 254SMO at different strain rates with the temperatures of 1,000 C a, 1,050 C b, 1,100 C c, 1,150 C d (a) (b) Fig. 3 Plots of lnð _eþ versus ln[sinh(ar p )] a and ln[sinh(ar p )] versus (10,000/T) b rates and temperatures [19]. Therefore, in the present investigation, Eq. (5) is used to describe the dependence of the stress on deformation temperature and strain rate. Taking natural logarithm from each side of Eq. (6) yields to the following equation: ln½sinhðar p ÞŠ ¼ 1 Q n RT þ 1 n ln _e 1 ln A: n ð7þ Partial differentiation of Eq. (7) yields to the following equation:

5 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), those of AISI 304 (380 kj/mol) [20] and 316LN (459 kj/ mol) [6]. This may be because 254SMO has higher contents of Cr and Mo alloying elements than conventional austenitic stainless steel, and solution strengthening caused by these alloying elements may result in higher activation energy. In addition, the value of hot deformation activation energy is well above those for the self-diffusion of c-fe (268 kj/mol), Cr (292 kj/mol), Ni (280 kj/mol), and Mo (240 kj/mol) [21, 22], which is attributed to the occurrence of DRX [23]. 3.4 Power Dissipation Maps Fig. 4 Variation of the Zener Hollomon parameter with flow stress o ln _e Q ¼ R o ln½sinhðar p ÞŠ o ln½sinhðar p ÞŠ T oð1=tþ : _e ð8þ On the basis of Eqs. (4) and (5), the linear regression of the peak stress data results in the average values of 8.56 and 0.04 for n 1 and b, respectively. This gives the value of a = b/n 1 = As shown in the Fig. 3a and b, the average slope (k 1 ) of plot of ln( _e) versus ln[sinh(ar p )] and the average slope (k 2 ) of plot of ln[sinh(ar p )] versus (10,000/T) are 6.2 and 0.96, respectively. Thus, according to Eq. (8), the value of the apparent activation energy (Q = 10000Rk 1 k 2 ) calculated is 496 kj/mol. It follows from the expression of Eqs. (3) and (6) that the plot of ln(z) versus ln[sinh(ar p )] was obtained, by which the values of n and A were determined to be 6.13 and , respectively, as shown in Fig. 4. The plots show a good linear correlation between the peak stress and Z value with regression coefficient of 0.99, indicating that the hyperbolic-sine function can fit well the data obtained under various hot deformation conditions. Consequently, substituting the values a, n, Q, and A into Eq. (7) yields to the peak stress constitutive equation of hot deformation for 254SMO, as expressed in the following Z ¼ _e exp RT ¼ 1: ½sinhð0:0047r p ÞŠ 6:13 ; ð9þ or r p ¼ 212:766 lnfðz 1: Þ 1=6:13 þ½ðz 1: Þ 2=6:13 þ 1Š 1 2 g: ð10þ In the present investigation, the average value of the apparent activation energy for hot deformation of 254SMO (496 kj/mol) is consistent with the reported value (508 kj/mol) obtained by Ren et al. [15]. This value is much larger than Typical power dissipation maps for superaustenitic stainless steel 254SMO obtained at strains of 0.2, 0.4, 0.6, 0.8, 1.0, and 1.1 are presented in Fig. 5. The strains of 0.2 and 0.4 correspond to the strains close to stress peak. The strains of 0.6 and 0.8 represent the situation at flow softening state, while the strains of 1.0 and 1.1 correspond to the steady state. The contour numbers in the maps represent the constant efficiencies of power dissipation expressed as percentages. Although exact values of the power dissipation efficiency at these strains are slightly different, the features of the processing maps are essentially similar, which indicate that the strain does not have a significant influence on processing maps. The power dissipation maps developed at the strain of 1.1 exhibit two different domains with relatively high peak efficiencies of power dissipation, as shown in Fig. 5f. The first domain occurs in the temperature range of 990 1,070 C and the strain rate range of s -1 with a peak efficiency of power dissipation of about 45% occurring at 1,050 C and 0.01 s -1. The second domain occurs in the temperature range of 1,140 1,200 C and the strain rate range of s -1 with a peak efficiency of power dissipation of about 39% occurring at 1,175 C and 0.01 s -1. These two domains are located at the low temperature and high temperature range, respectively. It should be noted that the curvature of the contour changes in the temperature range of 1,075 1,100 C, which is close to the dissolution temperature of sigma phase in the superaustenitic stainless steel 254SMO; and this phenomenon was also observed in other superalloy system, which is associated with the occurrence of phase transformation and dissolution [24]. The various domains of the processing maps could be illustrated based on the characteristic efficiency variation associated with the microstructural process [25]. As mentioned before, deformation mechanisms of the safe regimes include DRV, DRX, and superplasticity. During hot forging or rolling, several metallurgical phenomena such as work hardening, DRV, and DRX occur simultaneously. Especially, the occurrence of DRX can give rise to grain

6 318 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), Fig. 5 Power dissipation maps for 254SMO at true strains of 0.2 a, 0.4 b, 0.6 c, 0.8 d, 1.0 (e), 1.1 f (The contour number represents the percentage efficiency of power dissipation) refinement and significantly reduce the deformation resistance in practical steel forging or rolling. For austenitic stainless steel with higher deformation resistance than plain carbon steel, controlled rolling, which was based on the DRX, can provide a high quality of product [26]. Therefore, the condition of optimum deformation process and the control of the microstructure should be related to the DRX mechanism, which is thought to be a preferred choice for hot working. According to Raj maps [27], both the first and second domains correspond to either DRX or superplasticity. Generally, the efficiency values associated with DRX are about 30% 50%, while the values related to superplasticity are higher than 60%. In the present investigation, the peak efficiency in the maps is only approximately 49%. In view of this observation, the obtained domains do not represent superplastic deformation. The superaustenitic stainless steel used in this study seems to undergo dynamic recrystallization. The microstructural changes that occur in the deformed specimens can give evidence regarding the specified mechanism dominating in the domain [28]. In the first domain of power dissipation map obtained at the strain of 1.1, the values of power dissipation at 1,000 and 1,050 C with strain rate of 0.01 s -1 are about 33 and 49%, respectively. From these two values, it can be deduced that the microstructure formed under these conditions are fine grained. However, as can be seen in Fig. 6a, b, the resultant microstructure is inhomogeneous and contains partially recrystallized fine grains and some initial grains that had not undergone DRX. The precipitates are not visually identified by means of optical microscopy. The precipitates that take place intergranularly and also within the grains of the specimen deformed at 1,050 C at strain rate of 0.01 s -1 are observed by SEM, as shown in Fig. 7a, and EDS analysis (Fig. 7b) indicates that the precipitates are riched in Cr, Ni, and Mo. Fig. 8a presents the TEM micrograph of the precipitates with the size range of nm, which are determined to be tetragonal sigma phase through examination of selected area diffraction pattern (indicated in Fig. 8b). The equilibrium phase diagrams of the alloy calculated by Thermo-calc based on TCFE3 thermodynamic database are shown in the Fig. 9. According to the diagrams, the sigma phase is the dominant second phase in the temperatures between 1,000 and 1,100 C, which show a good agreement with the result obtained by TEM examination. It has been reported that the dissolution or growth of the particles or the phase under dynamic conditions and deformation-induced phase transformation or precipitation under dynamic conditions may contribute to the changes in the value of power dissipation [10]. Therefore, based on above analysis, it can be concluded that high

7 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), values of power dissipation in the first domain are associated with the combined occurrence of partial DRX and sigma-phase precipitation. It is important to note that lattice defects formed during hot deformation can significantly accelerate the precipitation rate of the sigma phase, and dynamic precipitation occurring at high-angle Fig. 6 Optical micrographs of 254SMO deformed to a true strain of 1.2 at the strain rate of 0.01 s -1 and the temperatures of 1,000 C a and 1,050 C b Fig. 7 a SEM micrograph of 254SMO deformed to a true strain of 1.2 at the strain rate of 0.01 s -1 and temperature of 1,050 C, b EDS analysis of chemical composition of precipitates in Fig. 7a Fig. 8 a TEM micrograph of precipitates in the specimen deformed to a true strain of 1.2 at the strain rate of 0.01 s -1 and temperature of 1,050 C, b selected area diffraction pattern of precipitate in Fig. 8a

8 320 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), grain boundaries can retard the DRX by consuming the stored energy of deformation. In addition, the hard and brittle sigma phase may seriously deteriorate hot plasticity of material, and make the steel difficult to coil and even lead to formation of crack during hot working. Therefore, the final forging or rolling should be carried out at the temperatures higher than the susceptible temperature to the precipitation of the sigma phase. Fig. 9 Phase diagrams of 254SMO simulated by themo-calc Fig. 10 Optical micrographs of 254SMO deformed to a true strain of 1.2 at 1,150 C and 0.1 s -1 a, 1,150 C and 1 s -1 b, 1,200 C and 0.1 s -1 c, 1,200 C 1s -1 d

9 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), The typical microstructures of the specimens deformed at the temperatures of 1,150 and 1,200 C at the strain rates of 0.1 and 1 s -1 in the second domain of power dissipation map are shown in Fig. 10. It can be seen that the microstructures are composed of homogeneous grains with some wavy and bulging boundaries and grain interiors nearly free of annealing twins, which exhibit the typical features of DRX. The corresponding flow curves under conditions within this domain also show the typical DRX features which include a peak in the flow stress followed by steady state at large strains. Table 1 lists the grain sizes of specimens deformed under different conditions in the second domain. As shown in Table 1, the size of new DRX grains are much finer than the initial grain size of the specimens soaked at 1,200 C for 180 s, and increase with decreasing of strain rate and increasing of the temperature. Due to the lower strain rate as well as the higher temperature in this domain, the second phase seems to be almost completely dissolved, and the full DRX and defect-free microstructures can be achieved without suffering from the risk of the precipitation of sigma phase. Therefore, the second domain is considered an optimum window for hot working of superaustenitic stainless steel 254SMO. It is worth mentioning that there exists a domain in the temperature range of 1,075 1,130 C and the strain rate range of s -1. This domain starts appearing at strain of 0.6, and the efficiency contours are closely spaced, indicating a sharp change in the efficiency values with increasing temperature and strain rate. The minimum efficiency of power dissipation in this domain decreases with increasing of strain, and the negative value of efficiency occurs at the strain higher than 0.8. From the Fig. 5f, it can be seen that the values of efficiencies are negative in the temperature range of 1,085 1,115 C and the strain rate range of s -1, although this regime is located in the stable domain. Not coincidentally, the flow curve obtained at 1,100 C and 0.01 s -1 presents secondary work hardening as strain exceeds 0.8, as indicated in the Fig. 2c. Under this deformation condition, the softening due to DRV and DRX cannot overcome the hardening caused by accumulation of dislocation. As expected, the secondary work hardening could be eliminated at higher temperatures due to the enhanced DRV and DRX, as shown in the Fig. 2d. Therefore, secondary work hardening could be responsible for the occurrence of negative values of efficiencies. 3.5 Instability Maps In order to delineate the regions of flow instability, the instability map for 254SMO was developed on the basis of the instability criterion corresponding to a strain of 1.1, as shown in Fig. 11. The map predicts a wide region of flow instability in the temperature range of 950 1,140 C and the strain rate ranging from 0.1 to 10 s -1, and in this regime, the values of n ( _e) are negative. The lower flow instability value of n ( _e) at higher strain rate indicates the higher possibility of unstable flow. The dissipation efficiency values obtained in this region is relatively lower than that in other stable domains, indicating that the most of the plastic power input was converted to heat and dissipated in the form of temperature rise in the material [24]. The corresponding flow curves in the instable domain exhibit a high rate of working hardening up to a peak at larger strain. The typical micrographs of the specimens deformed under the conditions in the instability region are shown in Fig. 12. As indicated in the Fig. 12a, the microstructure shows local deformation bands and demonstrates the existence of instability in a form of unstable flow due to flow localization, and fine grains are formed due to recrystallization within the bands. In addition, the elongated grains with large quantities of mechanical twin inside are observed in the Fig. 12b d. Prasad and Seshacharyulu [29] reported that flow localization and mechanical twin are detrimental to the processing of the Table 1 Results of grain size measurement of the 254SMO was treated with different treatments Treatment Grain size (lm) Grade Soaking at 1,200 C for 180 s Compression at 1,150 C and 0.01 s Compression at 1,150 C and 0.1 s Compression at 1,150 C and 1 s Compression at 1,200 C and 0.1 s Compression at 1,200 C and 1 s Fig. 11 Instability map for 254SMO at the true strain of 1.1

10 322 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), Fig. 12 Optical micrographs of 254SMO deformed to a true strain of 1.2 at 950 C and 1 s -1 a, 950 C and 10 s -1 b, 1,000 C and 10 s -1 c, and 1,050 C 10s -1 d (The compression direction is vertical) material. It should be noted that both deformation bands and mechanical twin boundaries tend to form about 45 with respect to the direction of compression. It is well known that during compression, planes of maximum shear stress are 45 to the axis of compression, and plastic deformation occurs on this plane. Since the time for hot deformation is short at lower temperatures and higher strain rates, the heat generated by plastic deformation is not conducted away to the colder parts of the specimen, and the higher temperature regions with lower flow stress are preferred deforming further causing localization [30]. In view of the analysis above, the hot working of superaustenitic stainless steel 254SMO should be avoided in the instability region. As shown in the Fig. 12a, c, d, DRX prefers to occur in high dislocation density and strain gradient, such as grain boundaries, deformation bands, and twinning boundaries, and recrystallized grain formed a typical necklace structure at the prior austenite grain boundaries. Asli and Hanzaki [31] pointed out that DRX nucleation occurs by bulging and successive strain induced boundary migration by investigating the DRX behavior of a superaustenitic stainless steel, and Ren et al. [15] also obtained the same result through the study on hot plasticity of 254SMO. However, the mechanism of DRX nucleation is still ambiguous, and further investigations are necessary for providing the evidences of illustration for the DRX behavior of 254SMO. 4 Conclusions (1) At temperatures below 1,050 C, the stress strain curves exhibit continuous flow softening no matter the strain rate, while at the temperatures higher than 1,050 C and the strain rates lower than 10 s -1, the flow stresses increase to a peak followed by a slow strain softening and then a steady state is reached. (2) Constitutive equation for hot deformation is proposed to describe the dependence of maximum stress on the temperature and the strain rate, and hot deformation apparent activation energy of the 254SMO is about 496 kj/mol. (3) The optimum window for hot working of superaustenitic stainless steel 254SMO is in the temperature range of 1,140 1,200 C and the strain rate range of

11 Enxiang Pu et al.: Acta Metall. Sin. (Engl. Lett.), 2014, 27(2), s -1 with a peak efficiency of power dissipation of 39% occurring at 1,175 C and 0.01 s -1, and in this domain, the microstructures with full DRX can be achieved. (4) The instability map shows that flow instability occurs at the temperatures below 1,140 C and the strain rates above 0.1 s -1, and local deformation bands and deformation twins are observed in this region. (5) The negative values of efficiencies occurs in the temperature range of 1,085 1,115 C and the strain rate range of s -1 due to secondary work hardening at the strains higher than 0.8. References [1] T. Koutsoukis, S. Zormalia, P. Kokkonidis, Adv. Mater. Res , 301 (2011) [2] C.C. Wu, S.H. Wang, C.Y. Chen, Scr. Mater. 56, 717 (2007) [3] E.A. Abd El Meguid, A.A. Abd El Latif, Corros. Sci. 46, 2431 (2004) [4] Y.V.R.K. Prasad, T. Seshacharyulu, Mater. Sci. Eng. A 243, 82 (1998) [5] N. Srinivasan, Y.V.R.K. Prasad, Metall. Mater. Trans. A 25, 2275 (1994) [6] B.F. Guo, H.P. Ji, X.G. Liu, L. Gao, J. Mater. Eng. Perform. 21, 1455 (2012) [7] S. Venugopal, S.L. Mannan, Y.V.R.K. Prasad, Metall. Trans. A 23, 3093 (1992) [8] S.P. Tan, Z.H. Wang, S.C. Cheng, Mater. Sci. Eng. A 517, 312 (2009) [9] Y.V.R.K. Prasad, Metall. Mater. Trans. A 27, 235 (1996) [10] Y.V.R.K. Prasad, H.L. Gegel, S.M. Draivelu, J.C. Malas, J.T. Morgan, K.A. Lark, D.A. Barker, Metall. Trans. A 15, 1883 (1984) [11] Y.V.R.K. Prasad, Mater. Eng. Perform. 12, 638 (2003) [12] H. Ziegler, in Progress in solid mechanics, vol. 4, ed. by I.N. Sneedon, R. Hill (Wiley, New York, 1963), p. 63 [13] S.C. Medeoros, W.G. Frazier, Y.V.R.K. Prasad, Metall. Mater. Trans. A 31, 2317 (2000) [14] S.J. Lee, Y.K. Lee, Mater. Des. 29, 1840 (2008) [15] J.B. Ren, Z.G. Song, W.J. Zheng, J.Z. Xiang, J. Iron Steel Res. 24, 41 (2012). (in Chinese) [16] Y. Wang, W.Z. Shao, L. Zhen, L. Yang, X.M. Zhang, Mater. Sci. Eng. A 497, 479 (2008) [17] Y.V.R.K. Prasad, N. Ravichandran, Bull. Mater. Sci. 14, 1241 (1991) [18] C.M. Sellars, W.J.M. Tegart, Int. Metall. Rev. 17, 1 (1972) [19] H. Mirzadeh, J.M. Cabrera, J.M. Prado, A. Najafizadeh, Mater. Sci. Eng. A 528, 3876 (2011) [20] S.I. Kim, Y.C. Yoo, Mater. Sci. Eng. A 311, 108 (2001) [21] Q.L. Yong, Secondary phases in steels (Metallurgy Industry press, Beijing, 2006), pp in Chinese [22] W.F. Smith, J. Hashemi, Foundations of materials science and engineering (China Machine Press, Beijing, 2011), pp [23] L. Briottet, J.J. Jonas, F. Montheillet, Acta Mater. 44, 1665 (1996) [24] Y.Q. Ning, Z.K. Yao, H. Li, Mater. Sci. Eng. A 527, 961 (2010) [25] S.C. Medeiros, Y.V.R.K. Prasad, W.G. Frazier, R. Srinivasan, Mater. Sci. Eng. A 193, 198 (2000) [26] S.H. Cho, S.I. Kim, Y.C. Yoo, J. Mater. Sci. Lett. 16, 1836 (1997) [27] R. Raj, Metall. Trans. A 12, 1089 (1981) [28] Y. Wang, L. Zhen, W.Z. Shao, L. Yang, X.M. Zhang, J. Alloys Compd. 474, 341 (2009) [29] Y.V.R.K. Prasad, T. Seshacharyulu, Int. Mater. Rev. 43, 245 (1998) [30] S.L. Guo, D.F. Li, H.J. Pen, Q.M. Guo, J. Hu, J. Nucl. Mater. 410, 52 (2011) [31] A.H. Asli, A.Z. Hanzaki, J. Mater. Sci. Technol. 25, 603 (2009)