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Hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel
Materials and Design 32 (2011) 4443–4448
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel Dening Zou a,⇑, Ying Han b, Dongna Yan a, Duo Wang a, Wei Zhang c, Guangwei Fan b,c a
School of Metallurgy and Engineering, Xi’an University of Architecture and Technology, Xi’an, Shaanxi 710055, China State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an, Shaanxi 710049, China c Technology Center, Taiyuan Iron and Steel (Group) Co. Ltd., Taiyuan, Shanxi 030003, China b
a r t i c l e
i n f o
Article history: Received 4 January 2011 Accepted 26 March 2011 Available online 31 March 2011 Keywords: A. Ferrous metals and alloys C. Forging F. Plastic behavior
a b s t r a c t The hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel was investigated by hot compression and hot tension tests conducted over the temperature range of 950–1200 °C and strain rates varying between 0.1 and 50 s1. The processing map technique was applied on the basis of dynamic materials model and Prasad instability criterion. Microstructure evolutions, Zener–Hollomon parameter as well as hot tensile ductility were examined. The results show that, as for the hot working of 00Cr13Ni5Mo2 supermartensitic stainless steel in the industrial production, the large strain deformation should be carried out in the temperature range 1140–1200 °C and strain rate range 0.1–50 s1, where the corresponding Zener–Hollomon parameters exhibit low values. Moreover, when deformed under high strain rate range (above 15 s1), the deformation temperature can be reduced reasonably. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The supermartensitic stainless steels have attracted substantial attention recently due to their advantages over conventional martensitic stainless steel, such as the unique combination of high strength and high toughness, good weldability and excellent corrosion resistance. Because of these advantages, they have been increasingly used as pipelines in deeper/ultra-deeper oil and gas wells being usually operated at ever higher temperatures and pressures, and also containing hostile gases and ions such as H2S, CO2 and Cl [1,2]. In addition, they only need simple heat treatment and have advantages in reducing the thickness of pipe wall, pipe weight and cost. So the supermartensitic stainless steels can substitute more expensive duplex stainless steels in some certain conditions, especially in stripper wells [3,4]. Much work has been performed on the balance of chemical composition and thermal history during heat treatment, and corrosion resistance in diverse demanding environments [5–9]. These developments have made a significant contribution to maximize and evaluate the performance of supermartensitic stainless steels. However, in an industrial process, hot forging or rolling is also the main procedure for manufacturing pipelines. The optimum hot working parameters are necessary to improve the properties and product rate of the pipes. However, little literature is available on the theories of hot working technology of these steels in actual production. The hot processing map technique based on dynamic ⇑ Corresponding author. Tel.: +86 29 82201074; fax: +86 29 82202921. E-mail addresses: [email protected], [email protected] (D. Zou), [email protected], [email protected] (Y. Han). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.03.067
materials model (DMM) in the frame of deformation temperature and logarithm of strain rate is a useful approach to optimize the hot working parameters and control the microstructure and properties of the product [10]. Furthermore, the thermal deformation mechanisms can be reflected by the processing map which has been successfully applied in a wide range of metal materials, such as alloys of magnesium [11], aluminum [12], titanium [13] and Nibased super alloys [14] as well as stainless steels [15,16]. In the present paper, the hot processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel were constituted based on the results in hot compression tests. On this basis as well as microstructural analysis and Zener–Hollomon parameter evolution, the hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel was investigated. Besides, the hot tensile experiments were preformed to check the validity of those maps. 2. Experimental procedures The chemical composition of 00Cr13Ni5Mo2 supermartensitic stainless steel used in this work is as follows (mass%): 0.013C, 0.18Si, 0.59Mn, 12.97Cr, 4.92Ni, 2.04Mo, 0.06N, 0.024Nb, 0.03W, and balance Fe. The material had been manufactured by Taiyuan Iron and Steel (Group) Co. Ltd. and supplied as hot forged round bars of 130 mm in diameter. Hot compress specimens (U8 12 mm) and hot tensile specimens (U10 120 mm) were machined out of the as-received bars directly. The hot deformation tests were conducted on a Gleeble3800 thermal–mechanical simulator. The heating method involved first heating to 1250 °C and holding for 4 min and then cooling to the deformation temperature at about 5 °C/s was adopted to simulate the practical industrial process. Before
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D. Zou et al. / Materials and Design 32 (2011) 4443–4448
each testing, the specimen was soaked at the deformation temperature for 1 min to eliminate thermal gradients prior to deformation. Hot compression tests were preformed in the temperature range of 950–1200 °C and strain rate range of 0.1–50 s1 at an interval of an order of magnitude. All the specimens were deformed to a true strain of 0.9, and then instantly quenched into water to preserve the microstructure after high temperature deformation. Specimens for optical microscopy (OM) and scanning electron microscopy (SEM) were sectioned at mid plane parallel to the compression axis and the cut sections were mounted, mechanically polished according to the standard procedure and then etched with a solution of 5 g FeCl3 + 50 g HCl + 100 ml H2O. Hot tension tests were performed at temperatures of 950, 1000, 1050, 1100, 1150 and 1200 °C and at a strain rate of 5 s1. The reduction in area and resistance to deformation were measured after fracture. For each condition, two tests were carried out and the error obtained was acceptable (within 5%).
3. Results and discussion 3.1. True stresstrue strain curves of 00Cr13Ni5Mo2 supermartensitic stainless steel Fig. 1 shows the equilibrium phase diagram for the studied steel calculated by using the thermocalc software. It is confirmed that only austenite was formed in the temperature range between 950 and 1200 °C, where generously provides the advantage of obtaining an excellent combination of strength and ductility. When deformed above the recrystallization temperature the effect of strain on flow stress is insignificant, while strain rate and deforming temperature have strong effects on flow stress [1]. Fig. 2a shows the true stresstrue strain curves obtained at the strain rate of 0.1 s1 and various deformation temperatures ranging from 950 to 1200 °C. It is clear that the flow stress increases with the decrease in temperature. At temperatures above 1000 °C, the flow curves exhibit a peak stress followed by a slowly continuous decrease of flow stress to a steady level, which is the typical characterization of the occurrence of dynamic recrystallization [17,18]. The higher temperature, the lower peak strain corresponding to the peak stress and the longer steady state is. At the temperature of 950 °C, the flow curve exhibits steady flow after initial work hardening, with no signs of dynamic recrystallization characteristics. Fig. 2b shows the true stresstrue strain
Fig. 2. True stresstrue strain curves of 00Cr13Ni5Mo2 supermartensitic stainless steel under different deformation conditions: (a) 0.1 s1; (b) 1200 °C.
curves obtained at the temperature of 1200 °C and various strain rates ranging from 0.1 to 50 s1. The flow stress and peak strain increase with increasing strain rate. It is noted that, at the strain rate of 50 s1, the flow curve is characterized by waving shape, especially in the starting strain below 0.4. The reason for this may be the occurrence of discontinuous softening due to severe deformation or the oscillation of the experimental equipment in the initial deformation. With the increase in strain during hot compression, the flow stress becomes constant gradually. 3.2. Hot processing map The basis and principles of DMM for constructing the processing map have been described in detail earlier by Prasad and Sasidhara [10]. In this model, the work piece is considered to be a power dissipater, and the power is thought to be dissipated into two complementary parts: the power dissipated by plastic work and the power dissipated by microstructural changes. A non-dimensional efficiency index g (%), called the efficiency of power dissipation, is allowed to represent how efficiently the work piece dissipates energy under microstructural changes. It can be given as
Fig. 1. The equilibrium phase diagram calculated by using the thermocalc software for the steel.
g¼
2m mþ1
ð1Þ
D. Zou et al. / Materials and Design 32 (2011) 4443–4448
rÞ where, m is the strain rate sensitivity of flow stress given by @ðlog . @ðlog e_ Þ The variation of power dissipation efficiency with temperature, strain rate and strain constitutes a power dissipation map. However, during hot compression some materials may exhibit processes like dynamic strain aging (DSA), formation of new phases, adiabatic shear deformation, flow localization and void generation, all of which can make the materials instable and then decrease the ductility of the materials [16]. For the identification of these instabilities during plastic flow, the following continuum criterion is used [10]:
nðe_ Þ ¼
@ ln½m=ðm þ 1Þ þm60 @ ln e_
ð2Þ
Flow instabilities are predicted to occur when nðe_ Þ is negative. Consequently, the instability map can be obtained depending on the variation of nðe_ Þ with temperature and strain rate, and it is superimposed on the above power dissipation map to constitute a processing map. Fig. 3 shows the processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel at strains of 0.45 and 0.9. Contour numbers indicate percent efficiency of power dissipation and shaded region with different tonalities of gray represents flow instability occurring. The darker the color of gray tonality, the lower the calculated nðe_ Þ value is. Each domain on the map reflects microstructural mechanisms through the dissipation characteristics varied with dif-
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ferent operating under those conditions of the domain. It can be stated that the optimum hot working condition is when efficiency of power dissipation value is maximum in the un-shaded region [19]. As it is clearly seen, no significant differences are identified between the two processing maps except a limited increment of efficiency of power dissipation with increasing the strain. The evolution of power dissipation efficiency is very complicated over the entire range of test parameters. At strain rates below 1 s1, the efficiency of power dissipation increases with increasing temperature and decreasing strain rate. At strain rates between 1 s1 and 10 s1, contour lines present closed loops enlarged gradually in both directions of temperature and strain rate. However, at strain rates above 10 s1, the efficiency of power dissipation increases with the increase in strain rate and the maximum value is found at the mid of temperature range. For this reason, Fig. 3b can exhibit two domains (D1 and D2) with peak power dissipation efficiency of 60% (1150 °C, 0.1 s1) and 45% (1100 °C, 50 s1) respectively, which are higher than that in Fig. 3a. The domain (D3) has the lowest power dissipation efficiency, the same as the corresponding in Fig. 3a. Meanwhile, the unstable region at 950–1140 °C under a lower strain rate (about 1 s1) is observed on the map and little changes in shape with the increase in strain can be found. 3.3. Microstructural analysis Typical optical micrographs of the microstructure for the specimens deformed are shown in Fig. 4. Full equiaxed dynamic recrystallization has occurred in Fig. 4a, corresponding to the region D1 of the processing map, but the larger grains are also observed because of recrystallization grain boundaries sliding and merging [20]. Fig. 4b shows the microstructure after hot deformation at 1100 °C and 50 s1, corresponding to the region D2. While the peak power dissipation efficiency in this region is lower than that of D1, the dynamic recrystallization grains are much finer. Fig. 4c corresponds to the region with the lowest power dissipation efficiency (D3). It can be seen that dynamic recovery is the main softening mechanism. Partial dynamic recrystallization nucleates at the prior austenite grain boundaries. Furthermore, at a lower strain rate (0.1 s1) and 1000 °C, much more new recrystallization grains are formed evenly and a small amount of delta ferrite exhibiting thin strip is distributed along grain boundaries due to the deformation (see Fig. 4d). The deformed microstructures in the instability region are shown in Fig. 5. Microstructure examination of the specimen deformed at 950 °C and 1 s1 reveals that the grains are elongated and intense flow localization occurs obviously nearby grain boundaries because of possible unevenly deformation, as shown in Fig. 5a. Fig. 5b shows the severe deformation zone in the microstructure deformed at 1000 °C and 1 s1, which exhibits wavy or irregular grain boundaries and low level of dynamic recovery. From above analyses for processing maps and microstructures, fully and uniform dynamic recrystallization will occur at temperatures higher than 1140 °C for the experimental strain rates. The relationship between average dynamic recrystallization grain size and power dissipation efficiency at 1150 °C and different strain rates is plotted in Fig. 6. The dynamic recrystallization grain size indeed decreases with the increment of strain rate up to 50 s1 though the power dissipation efficiency also decreases and there is a trough at the strain rate of 10 s1. 3.4. Activation energy and Zener–Hollomon parameter
Fig. 3. Processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel at different strains: (a) 0.45; (b) 0.9.
Zener–Hollomon parameter (Z–parameter) as an important hot processing index, proposed by Zener and Hollomon in 1944 [21], is widely used to characterize the combined effect of strain rate and deformation temperature on the hot deformation process,
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Fig. 4. Microstructure characteristics of the specimens deformed to the strain of 0.9: (a) 1150 °C, 0.1 s1; (b) 1100 °C, 50 s1; (c) 1000 °C, 10 s1; (d) 1000 °C, 0.1 s1.
especially on the deformation resistance. Z–parameter is identified by the following equation [22,23]:
Z ¼ e_ expðQ =RTÞ ¼ FðrÞ
ð3Þ
where, Q is the apparent activation energy for deformation, e_ is the strain rate, T is absolute temperature and R is the gas constant. F(r) is the stress function which can be expressed by the following two forms on the basis of stress values [23,24]:
F 1 ðrÞ ¼ A1 rn1 ðr < 70 MPaÞ F 2 ðrÞ ¼ A2 expðb rÞ ðr > 100 MPaÞ
ð4Þ ð5Þ
where, A1, A2 and b are constants, and n1 is the stress exponent. By combining the Eqs. (4) and (5), Sellars and Tegart [25] established a constitutive equation of hyperbolic sine function for all the range of stress:
FðrÞ ¼ A½sinhða rÞn
ð6Þ
where, n is the stress exponent, A and a are material constants, and a can be calculated by the following equation:
a¼
b n1
ð7Þ
It is proven that Eq. (3) can account for both the temperature and strain rate-dependence of the flow stress for any strain. And the stress function described by Eq. (6) is used in this investigation. By applying linear regression of empirical data, the a is determined as 0.0067. Thus the apparent activation energy Q, for hot compression can be defined by:
@fln½sinhð0:0067 rÞg Q ¼ nR @ð1=TÞ e_ ;e
ð8Þ
Fig. 7 shows the calculated apparent activation energy at different temperatures. It can be seen that the apparent activation energy decreases with increasing temperature, and noted that greater values concentrate in the temperature range of 950–1000 °C, corresponding power dissipation efficiency in the map is lower. The average values of Q is determined as 439 kJ/mol, and it can be use to calculate ln Z values under different deformation conditions by means of Eq. (3) as shown in Fig. 8. Obviously, the Z– parameter decreases with increasing temperature and decreasing strain rate. Fig. 9 shows the relation between flow stress functioned by hyperbolic sine and ln Z. A good linear relationship between them is observed and the reciprocal of slope is identified as the mean values of n in Eq. (8). The lower the Z–parameter, the smaller the flow stress is induced. In fact, the Z–parameter also reflects the microstructural and substructural changes occurring during deformation [26]. A lower Z–parameter favors the flow softening and the dynamic recrystallization occurring. 3.5. Hot tensile properties of 00Cr13Ni5Mo2 supermartensitic stainless steel Fig. 10 shows the measured results of the resistance to deformation and reduction in area under different deformation conditions (5 s1, 950–1200 °C). It can be seen that the resistance to deformation decreases lineally with the enhancing temperature. However, the plasticity increases with the increasing temperature when temperature is lower than 1100 °C, and decreases slowly with the increasing temperature in the range of 1150–1200 °C. The peak reduction in area is obtained at 1100–1150 °C and reach
D. Zou et al. / Materials and Design 32 (2011) 4443–4448
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Fig. 6. Relationship between dynamic recrystallization grain size and power dissipation efficiency at 1150 °C and different strain rates.
Fig. 7. Variations of apparent activation energy with increasing temperature.
Fig. 5. Microstructure characteristics of the specimens deformed at the unstable region: (a) 950 °C, 1 s1; (b) 1000 °C, 1 s1.
about 95%, corresponding to resistance to deformation is from 101.9 to 112 MPa. The optimum plasticity region obtained here is much agreement with that in the processing map (see Fig. 3).
4. Conclusions (1) For 00Cr13Ni5Mo2 supermartensitic stainless steel deformed at 950–1200 °C and 0.1–50 s1, the processing maps were constructed. In the processing map with the true strain of 0.9, there exhibit two domains with peak power dissipation efficiency, D1 (1150 °C, 0.1 s1) and D2 (1100 °C, 50 s1), where deformation mechanisms are both complete dynamic recrystallization. Flow localization and low level of dynamic recovery are the main reasons for instability occurring.
Fig. 8. ln Z values of 00Cr13Ni5Mo2 supermartensitic stainless steel under different conditions.
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Acknowledgements This work is supported by the project of Developing Plan for Science and Technology Research of Shaanxi Province (2010K10– 13), and the Scientific Research Program funded by Shaanxi Provincial Education Department (2010JC10). References
Fig. 9. Linear relationship between flow stress functioned by hyperbolic sine and ln Z.
Fig. 10. Hot tensile properties of 00Cr13Ni5Mo2 supermartensitic stainless steel at different deformation temperatures (950–1200 °C) and 5 s1.
(2) The classical hyperbolic sine equation is applied to estimate the apparent activation energy and Zparameter at different working conditions. The apparent activation energy (mean value 439 kJ/mol) has a large increase at low temperatures. The lower the Z–parameter, the lower flow stress and the larger the extent of flow softening is. (3) The peak hot plasticity of 00Cr13Ni5Mo2 supermartensitic stainless steel is obtained at 1100–1150 °C and 5 s1, corresponding to resistance to deformation is 101.9–112 MPa. (4) During the industrial production, it is suggested that large strain deformation should be carried out in the temperature range 1140–1200 °C and strain rate range 0.1–50 s1 in the hot working of 00Cr13Ni5Mo2 supermartensitic stainless steel. Moreover, when deformed under high strain rate range (above 15 s1), the deformation temperature can be reduced reasonably.
[1] Nakamichi H, Sato K, Miyata Y, Kimura M, Masamura K. Quantitative analysis of Cr-depleted zone morphology in low carbon martensitic stainless steel using FE–(S) TEM. Corros Sci 2008;50(2):309–15. [2] Griffiths A, Nimmo W, Roebuck B, Hinds G, Turnbull A. A novel approach to characterizing the mechanical properties of supermartensitic 13Cr stainless steel welds. Mater Sci Eng A 2004;384:83–91. [3] Carrouge D, Bhadeshia HKDH, Woollin P. Effect of d-ferrite on impact properties of supermartensitic stainless steel heat affected zones. Sci Technol Weld Join 2004;9(5):377–89. [4] Zou DN, Han Y, Zhang W, Fang XD. Influence of tempering process on mechanical properties of 00Cr13Ni4Mo supermartensitic stainless steel. J Iron Steel Res Int 2010;17(8):50–4. [5] Leem DS, Lee YD, Jun JH, Choi CS. Amount of retained austenite at room temperature after reverse transformation of martensite to austenite in an Fe– 13%Cr–7%Ni–3%Si martensitic stainless steel. Scripta Mater 2001;45(7): 767–72. [6] Rodrigues CAD, Lorenzo PLD, Sokolowski A, Barbosa CA, Rollo JMDA. Titanium and molybdenum content in supermartensitic stainless steel. Mater Sci Eng A 2007;460–461:149–52. [7] Rodrigues CAD, Lorenzo PLD, Sokolowski A, Barbosa CA, Rollo JMDA. Development of a supermartensitic stainless steel microalloyed with niobium. J ASTM Int 2006;3(5):1–5. [8] Qin B, Wang ZY, Sun QS. Effect of tempering temperature on properties of 00Cr16Ni5Mo stainless steel. Mater Charact 2008;59(8): 1096–100. [9] Anselmo N, May JE, Mariano NA, Nascente PAP, Kuri SE. Corrosion behavior of supermartensitic stainless steel in aerated and CO2-saturated synthetic seawater. Mater Sci Eng A 2006;428:73–9. [10] Prasad YVRK, Sasidhara S. Hot working guide: a compendium of processing maps. America: ASM; 1997. [11] Srinivasan N, Prasad YVRK, Rama Rao P. Hot deformation behavior of Mg–3Al alloy – A study using processing map. Mater Sci Eng A 2008;476:146–56. [12] Meng G, Li BL, Li HM, Huang H, Nie ZR. Hot deformation and processing maps of an Al-5.7 wt% Mg alloy with erbium. Mater Sci Eng A 2009;517: 132–7. [13] Prasad YVRK, Seshacharyulu T. Processing maps for hot working of titanium alloys. Mater Sci Eng A 1998;243:82–8. [14] Cai DY, Xiong LY, Liu WC, Sun GD, Yao M. Characterization of hot deformation behavior of a Ni–base superalloy using processing map. Mater Des 2009;30:921–5. [15] Wang ZH, Fu WT, Wang BZ, Zhang WH, Lv ZQ, Jiang P. Study on hot deformation characteristics of 12%Cr ultra-super-critical rotor steel using processing maps and Zener–Hollomon parameter. Mater Charact 2010;61:25–30. [16] Venugopal S, Mannan SL, Prasad YVRK. Optimization of cold and warm workability in stainless steel type AISI 316L using instability maps. J Nucl Mater 1995;227:1–10. [17] Łyszkowski R, Bystrzycki J. Hot deformation and processing maps of an Fe3Al intermetallic alloy. Intermetallics 2006;14:1231–7. [18] Wu K, Liu GQ, Hu BF, Li F, Zhang YW, Tao Y, et al. Hot compressive deformation behavior of a new hot isostatically pressed Ni–Cr–Co based powder metallurgy superalloy. Mater Design 2011;32:1872–9. [19] Han Y, Zou DN, Chen ZY, Fan GW, Zhang W. Investigation on hot deformation behavior of 00Cr23Ni4N duplex stainless steel under medium-high strain rates. Mater Charact 2011;62:198–203. [20] Anbuselvan S, Ramanathan S. Hot deformation and processing maps of extruded ZE41A magnesium alloy. Mater Design 2010;31:2319–23. [21] Zener C, Hollomon JH. Effect of strain-rate upon the plastic flow of steel. J Appl Phys 1944;15(1):22–7. [22] Khamei AA, Dehghani K. Modeling the hot-deformation behavior of Ni60 wt%– Ti40 wt% intermetallic alloy. J Alloy Compd 2010;490:377–81. [23] Liu JT, Chang HB, Wu RH, Hsu TY, Ruan XY. Investigation on hot deformation behavior of AISI T1 high-speed steel. Mater Charact 2000;45:175–86. [24] Garofalo F. Fundamentals of creep and creep rupture in metals. New York: Macmillan; 1965. [25] Sellars CM, Mc Tegart WJ. On the mechanism of hot deformation. Acta Metall 1966;14:1136–8. [26] Tan SP, Wang ZH, Cheng SC, Liu ZD, Han JC, Fu WT. Processing maps and hot workability of Super304H austenitic heat-resistant stainless steel. Mater Sci Eng A 2009;517:312–5.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel Dening Zou a,⇑, Ying Han b, Dongna Yan a, Duo Wang a, Wei Zhang c, Guangwei Fan b,c a
School of Metallurgy and Engineering, Xi’an University of Architecture and Technology, Xi’an, Shaanxi 710055, China State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an, Shaanxi 710049, China c Technology Center, Taiyuan Iron and Steel (Group) Co. Ltd., Taiyuan, Shanxi 030003, China b
a r t i c l e
i n f o
Article history: Received 4 January 2011 Accepted 26 March 2011 Available online 31 March 2011 Keywords: A. Ferrous metals and alloys C. Forging F. Plastic behavior
a b s t r a c t The hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel was investigated by hot compression and hot tension tests conducted over the temperature range of 950–1200 °C and strain rates varying between 0.1 and 50 s1. The processing map technique was applied on the basis of dynamic materials model and Prasad instability criterion. Microstructure evolutions, Zener–Hollomon parameter as well as hot tensile ductility were examined. The results show that, as for the hot working of 00Cr13Ni5Mo2 supermartensitic stainless steel in the industrial production, the large strain deformation should be carried out in the temperature range 1140–1200 °C and strain rate range 0.1–50 s1, where the corresponding Zener–Hollomon parameters exhibit low values. Moreover, when deformed under high strain rate range (above 15 s1), the deformation temperature can be reduced reasonably. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The supermartensitic stainless steels have attracted substantial attention recently due to their advantages over conventional martensitic stainless steel, such as the unique combination of high strength and high toughness, good weldability and excellent corrosion resistance. Because of these advantages, they have been increasingly used as pipelines in deeper/ultra-deeper oil and gas wells being usually operated at ever higher temperatures and pressures, and also containing hostile gases and ions such as H2S, CO2 and Cl [1,2]. In addition, they only need simple heat treatment and have advantages in reducing the thickness of pipe wall, pipe weight and cost. So the supermartensitic stainless steels can substitute more expensive duplex stainless steels in some certain conditions, especially in stripper wells [3,4]. Much work has been performed on the balance of chemical composition and thermal history during heat treatment, and corrosion resistance in diverse demanding environments [5–9]. These developments have made a significant contribution to maximize and evaluate the performance of supermartensitic stainless steels. However, in an industrial process, hot forging or rolling is also the main procedure for manufacturing pipelines. The optimum hot working parameters are necessary to improve the properties and product rate of the pipes. However, little literature is available on the theories of hot working technology of these steels in actual production. The hot processing map technique based on dynamic ⇑ Corresponding author. Tel.: +86 29 82201074; fax: +86 29 82202921. E-mail addresses: [email protected], [email protected] (D. Zou), [email protected], [email protected] (Y. Han). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.03.067
materials model (DMM) in the frame of deformation temperature and logarithm of strain rate is a useful approach to optimize the hot working parameters and control the microstructure and properties of the product [10]. Furthermore, the thermal deformation mechanisms can be reflected by the processing map which has been successfully applied in a wide range of metal materials, such as alloys of magnesium [11], aluminum [12], titanium [13] and Nibased super alloys [14] as well as stainless steels [15,16]. In the present paper, the hot processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel were constituted based on the results in hot compression tests. On this basis as well as microstructural analysis and Zener–Hollomon parameter evolution, the hot workability of 00Cr13Ni5Mo2 supermartensitic stainless steel was investigated. Besides, the hot tensile experiments were preformed to check the validity of those maps. 2. Experimental procedures The chemical composition of 00Cr13Ni5Mo2 supermartensitic stainless steel used in this work is as follows (mass%): 0.013C, 0.18Si, 0.59Mn, 12.97Cr, 4.92Ni, 2.04Mo, 0.06N, 0.024Nb, 0.03W, and balance Fe. The material had been manufactured by Taiyuan Iron and Steel (Group) Co. Ltd. and supplied as hot forged round bars of 130 mm in diameter. Hot compress specimens (U8 12 mm) and hot tensile specimens (U10 120 mm) were machined out of the as-received bars directly. The hot deformation tests were conducted on a Gleeble3800 thermal–mechanical simulator. The heating method involved first heating to 1250 °C and holding for 4 min and then cooling to the deformation temperature at about 5 °C/s was adopted to simulate the practical industrial process. Before
4444
D. Zou et al. / Materials and Design 32 (2011) 4443–4448
each testing, the specimen was soaked at the deformation temperature for 1 min to eliminate thermal gradients prior to deformation. Hot compression tests were preformed in the temperature range of 950–1200 °C and strain rate range of 0.1–50 s1 at an interval of an order of magnitude. All the specimens were deformed to a true strain of 0.9, and then instantly quenched into water to preserve the microstructure after high temperature deformation. Specimens for optical microscopy (OM) and scanning electron microscopy (SEM) were sectioned at mid plane parallel to the compression axis and the cut sections were mounted, mechanically polished according to the standard procedure and then etched with a solution of 5 g FeCl3 + 50 g HCl + 100 ml H2O. Hot tension tests were performed at temperatures of 950, 1000, 1050, 1100, 1150 and 1200 °C and at a strain rate of 5 s1. The reduction in area and resistance to deformation were measured after fracture. For each condition, two tests were carried out and the error obtained was acceptable (within 5%).
3. Results and discussion 3.1. True stresstrue strain curves of 00Cr13Ni5Mo2 supermartensitic stainless steel Fig. 1 shows the equilibrium phase diagram for the studied steel calculated by using the thermocalc software. It is confirmed that only austenite was formed in the temperature range between 950 and 1200 °C, where generously provides the advantage of obtaining an excellent combination of strength and ductility. When deformed above the recrystallization temperature the effect of strain on flow stress is insignificant, while strain rate and deforming temperature have strong effects on flow stress [1]. Fig. 2a shows the true stresstrue strain curves obtained at the strain rate of 0.1 s1 and various deformation temperatures ranging from 950 to 1200 °C. It is clear that the flow stress increases with the decrease in temperature. At temperatures above 1000 °C, the flow curves exhibit a peak stress followed by a slowly continuous decrease of flow stress to a steady level, which is the typical characterization of the occurrence of dynamic recrystallization [17,18]. The higher temperature, the lower peak strain corresponding to the peak stress and the longer steady state is. At the temperature of 950 °C, the flow curve exhibits steady flow after initial work hardening, with no signs of dynamic recrystallization characteristics. Fig. 2b shows the true stresstrue strain
Fig. 2. True stresstrue strain curves of 00Cr13Ni5Mo2 supermartensitic stainless steel under different deformation conditions: (a) 0.1 s1; (b) 1200 °C.
curves obtained at the temperature of 1200 °C and various strain rates ranging from 0.1 to 50 s1. The flow stress and peak strain increase with increasing strain rate. It is noted that, at the strain rate of 50 s1, the flow curve is characterized by waving shape, especially in the starting strain below 0.4. The reason for this may be the occurrence of discontinuous softening due to severe deformation or the oscillation of the experimental equipment in the initial deformation. With the increase in strain during hot compression, the flow stress becomes constant gradually. 3.2. Hot processing map The basis and principles of DMM for constructing the processing map have been described in detail earlier by Prasad and Sasidhara [10]. In this model, the work piece is considered to be a power dissipater, and the power is thought to be dissipated into two complementary parts: the power dissipated by plastic work and the power dissipated by microstructural changes. A non-dimensional efficiency index g (%), called the efficiency of power dissipation, is allowed to represent how efficiently the work piece dissipates energy under microstructural changes. It can be given as
Fig. 1. The equilibrium phase diagram calculated by using the thermocalc software for the steel.
g¼
2m mþ1
ð1Þ
D. Zou et al. / Materials and Design 32 (2011) 4443–4448
rÞ where, m is the strain rate sensitivity of flow stress given by @ðlog . @ðlog e_ Þ The variation of power dissipation efficiency with temperature, strain rate and strain constitutes a power dissipation map. However, during hot compression some materials may exhibit processes like dynamic strain aging (DSA), formation of new phases, adiabatic shear deformation, flow localization and void generation, all of which can make the materials instable and then decrease the ductility of the materials [16]. For the identification of these instabilities during plastic flow, the following continuum criterion is used [10]:
nðe_ Þ ¼
@ ln½m=ðm þ 1Þ þm60 @ ln e_
ð2Þ
Flow instabilities are predicted to occur when nðe_ Þ is negative. Consequently, the instability map can be obtained depending on the variation of nðe_ Þ with temperature and strain rate, and it is superimposed on the above power dissipation map to constitute a processing map. Fig. 3 shows the processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel at strains of 0.45 and 0.9. Contour numbers indicate percent efficiency of power dissipation and shaded region with different tonalities of gray represents flow instability occurring. The darker the color of gray tonality, the lower the calculated nðe_ Þ value is. Each domain on the map reflects microstructural mechanisms through the dissipation characteristics varied with dif-
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ferent operating under those conditions of the domain. It can be stated that the optimum hot working condition is when efficiency of power dissipation value is maximum in the un-shaded region [19]. As it is clearly seen, no significant differences are identified between the two processing maps except a limited increment of efficiency of power dissipation with increasing the strain. The evolution of power dissipation efficiency is very complicated over the entire range of test parameters. At strain rates below 1 s1, the efficiency of power dissipation increases with increasing temperature and decreasing strain rate. At strain rates between 1 s1 and 10 s1, contour lines present closed loops enlarged gradually in both directions of temperature and strain rate. However, at strain rates above 10 s1, the efficiency of power dissipation increases with the increase in strain rate and the maximum value is found at the mid of temperature range. For this reason, Fig. 3b can exhibit two domains (D1 and D2) with peak power dissipation efficiency of 60% (1150 °C, 0.1 s1) and 45% (1100 °C, 50 s1) respectively, which are higher than that in Fig. 3a. The domain (D3) has the lowest power dissipation efficiency, the same as the corresponding in Fig. 3a. Meanwhile, the unstable region at 950–1140 °C under a lower strain rate (about 1 s1) is observed on the map and little changes in shape with the increase in strain can be found. 3.3. Microstructural analysis Typical optical micrographs of the microstructure for the specimens deformed are shown in Fig. 4. Full equiaxed dynamic recrystallization has occurred in Fig. 4a, corresponding to the region D1 of the processing map, but the larger grains are also observed because of recrystallization grain boundaries sliding and merging [20]. Fig. 4b shows the microstructure after hot deformation at 1100 °C and 50 s1, corresponding to the region D2. While the peak power dissipation efficiency in this region is lower than that of D1, the dynamic recrystallization grains are much finer. Fig. 4c corresponds to the region with the lowest power dissipation efficiency (D3). It can be seen that dynamic recovery is the main softening mechanism. Partial dynamic recrystallization nucleates at the prior austenite grain boundaries. Furthermore, at a lower strain rate (0.1 s1) and 1000 °C, much more new recrystallization grains are formed evenly and a small amount of delta ferrite exhibiting thin strip is distributed along grain boundaries due to the deformation (see Fig. 4d). The deformed microstructures in the instability region are shown in Fig. 5. Microstructure examination of the specimen deformed at 950 °C and 1 s1 reveals that the grains are elongated and intense flow localization occurs obviously nearby grain boundaries because of possible unevenly deformation, as shown in Fig. 5a. Fig. 5b shows the severe deformation zone in the microstructure deformed at 1000 °C and 1 s1, which exhibits wavy or irregular grain boundaries and low level of dynamic recovery. From above analyses for processing maps and microstructures, fully and uniform dynamic recrystallization will occur at temperatures higher than 1140 °C for the experimental strain rates. The relationship between average dynamic recrystallization grain size and power dissipation efficiency at 1150 °C and different strain rates is plotted in Fig. 6. The dynamic recrystallization grain size indeed decreases with the increment of strain rate up to 50 s1 though the power dissipation efficiency also decreases and there is a trough at the strain rate of 10 s1. 3.4. Activation energy and Zener–Hollomon parameter
Fig. 3. Processing maps of 00Cr13Ni5Mo2 supermartensitic stainless steel at different strains: (a) 0.45; (b) 0.9.
Zener–Hollomon parameter (Z–parameter) as an important hot processing index, proposed by Zener and Hollomon in 1944 [21], is widely used to characterize the combined effect of strain rate and deformation temperature on the hot deformation process,
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Fig. 4. Microstructure characteristics of the specimens deformed to the strain of 0.9: (a) 1150 °C, 0.1 s1; (b) 1100 °C, 50 s1; (c) 1000 °C, 10 s1; (d) 1000 °C, 0.1 s1.
especially on the deformation resistance. Z–parameter is identified by the following equation [22,23]:
Z ¼ e_ expðQ =RTÞ ¼ FðrÞ
ð3Þ
where, Q is the apparent activation energy for deformation, e_ is the strain rate, T is absolute temperature and R is the gas constant. F(r) is the stress function which can be expressed by the following two forms on the basis of stress values [23,24]:
F 1 ðrÞ ¼ A1 rn1 ðr < 70 MPaÞ F 2 ðrÞ ¼ A2 expðb rÞ ðr > 100 MPaÞ
ð4Þ ð5Þ
where, A1, A2 and b are constants, and n1 is the stress exponent. By combining the Eqs. (4) and (5), Sellars and Tegart [25] established a constitutive equation of hyperbolic sine function for all the range of stress:
FðrÞ ¼ A½sinhða rÞn
ð6Þ
where, n is the stress exponent, A and a are material constants, and a can be calculated by the following equation:
a¼
b n1
ð7Þ
It is proven that Eq. (3) can account for both the temperature and strain rate-dependence of the flow stress for any strain. And the stress function described by Eq. (6) is used in this investigation. By applying linear regression of empirical data, the a is determined as 0.0067. Thus the apparent activation energy Q, for hot compression can be defined by:
@fln½sinhð0:0067 rÞg Q ¼ nR @ð1=TÞ e_ ;e
ð8Þ
Fig. 7 shows the calculated apparent activation energy at different temperatures. It can be seen that the apparent activation energy decreases with increasing temperature, and noted that greater values concentrate in the temperature range of 950–1000 °C, corresponding power dissipation efficiency in the map is lower. The average values of Q is determined as 439 kJ/mol, and it can be use to calculate ln Z values under different deformation conditions by means of Eq. (3) as shown in Fig. 8. Obviously, the Z– parameter decreases with increasing temperature and decreasing strain rate. Fig. 9 shows the relation between flow stress functioned by hyperbolic sine and ln Z. A good linear relationship between them is observed and the reciprocal of slope is identified as the mean values of n in Eq. (8). The lower the Z–parameter, the smaller the flow stress is induced. In fact, the Z–parameter also reflects the microstructural and substructural changes occurring during deformation [26]. A lower Z–parameter favors the flow softening and the dynamic recrystallization occurring. 3.5. Hot tensile properties of 00Cr13Ni5Mo2 supermartensitic stainless steel Fig. 10 shows the measured results of the resistance to deformation and reduction in area under different deformation conditions (5 s1, 950–1200 °C). It can be seen that the resistance to deformation decreases lineally with the enhancing temperature. However, the plasticity increases with the increasing temperature when temperature is lower than 1100 °C, and decreases slowly with the increasing temperature in the range of 1150–1200 °C. The peak reduction in area is obtained at 1100–1150 °C and reach
D. Zou et al. / Materials and Design 32 (2011) 4443–4448
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Fig. 6. Relationship between dynamic recrystallization grain size and power dissipation efficiency at 1150 °C and different strain rates.
Fig. 7. Variations of apparent activation energy with increasing temperature.
Fig. 5. Microstructure characteristics of the specimens deformed at the unstable region: (a) 950 °C, 1 s1; (b) 1000 °C, 1 s1.
about 95%, corresponding to resistance to deformation is from 101.9 to 112 MPa. The optimum plasticity region obtained here is much agreement with that in the processing map (see Fig. 3).
4. Conclusions (1) For 00Cr13Ni5Mo2 supermartensitic stainless steel deformed at 950–1200 °C and 0.1–50 s1, the processing maps were constructed. In the processing map with the true strain of 0.9, there exhibit two domains with peak power dissipation efficiency, D1 (1150 °C, 0.1 s1) and D2 (1100 °C, 50 s1), where deformation mechanisms are both complete dynamic recrystallization. Flow localization and low level of dynamic recovery are the main reasons for instability occurring.
Fig. 8. ln Z values of 00Cr13Ni5Mo2 supermartensitic stainless steel under different conditions.
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Acknowledgements This work is supported by the project of Developing Plan for Science and Technology Research of Shaanxi Province (2010K10– 13), and the Scientific Research Program funded by Shaanxi Provincial Education Department (2010JC10). References
Fig. 9. Linear relationship between flow stress functioned by hyperbolic sine and ln Z.
Fig. 10. Hot tensile properties of 00Cr13Ni5Mo2 supermartensitic stainless steel at different deformation temperatures (950–1200 °C) and 5 s1.
(2) The classical hyperbolic sine equation is applied to estimate the apparent activation energy and Zparameter at different working conditions. The apparent activation energy (mean value 439 kJ/mol) has a large increase at low temperatures. The lower the Z–parameter, the lower flow stress and the larger the extent of flow softening is. (3) The peak hot plasticity of 00Cr13Ni5Mo2 supermartensitic stainless steel is obtained at 1100–1150 °C and 5 s1, corresponding to resistance to deformation is 101.9–112 MPa. (4) During the industrial production, it is suggested that large strain deformation should be carried out in the temperature range 1140–1200 °C and strain rate range 0.1–50 s1 in the hot working of 00Cr13Ni5Mo2 supermartensitic stainless steel. Moreover, when deformed under high strain rate range (above 15 s1), the deformation temperature can be reduced reasonably.
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