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Hot deformation behavior of an austenitic Fe–20Mn–3Si–3Al transformation induced plasticity steel
Materials and Design 34 (2012) 713–718
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Hot deformation behavior of an austenitic Fe–20Mn–3Si–3Al transformation induced plasticity steel Dejun Li a,b,⇑, Yaorong Feng a,b, Zhifu Yin c, Fengshou Shangguan a, Ke Wang c, Qiang Liu b, Feng Hu a a
School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China CNPC Tubular Goods Research Institute, Xi’an 710065, China c Petroleum Research Institute, Shaanxi Yanchang Petroleum Group Company, Xi’an 710075, China b
a r t i c l e
i n f o
Article history: Received 31 January 2011 Accepted 18 May 2011 Available online 23 May 2011 Keywords: A. Ferrous metals and alloys C. Forming E. Physical
a b s t r a c t Hot deformation behavior of an austenitic Fe–20Mn–3Si–3Al transformation induced plasticity (TRIP) steel was investigated by hot compression tests on Gleeble 3500D thermo-mechanical simulator in the temperature ranges of 900–1100 °C and the strain rate ranges of 0.01–10 s1. The results show that the flow stress is sensitively dependent on deformation temperature and strain rate, and the flow stress increases with strain rate and decreases with deformation temperature. The peak stress during hot deformation can be predicted by the Zener–Hollomon (Z) parameter in the hyperbolic sine equation with the hot deformation activation energy Q of 387.84 kJ/mol. The dynamic recrystallization (DRX) is the most important softening mechanism for the experimental steel during hot compression. Furthermore, DRX procedure is strongly affected by Z parameter, and decreasing of Z value lead to more adequate proceeding of DRX. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The development of steels for automotive applications is focused on an increase of strength combined with the improvement of its ductility. The high strength enables car manufacturers to reduce the weight of the car bodies and improve vehicle safety, whereas the increasing ductility admits complex component design. In order to meet the needs of automotive industry, high manganese austenitic steels have gradually become a focus of research due to their outstanding mechanical properties. Outstanding mechanical properties combining a satisfactory strength and excellent ductility result from the strain-induced microstructure transformations which are called the transformation induced plasticity (TRIP) effect and the twinning induced plasticity (TWIP) effect [1–6]. The microstructure transformations, responsible for the high working hardening rate, are correlated to the stacking fault energy (SFE) of the austenitic matrix. The SFE is a function of the alloy composition and deformation temperature. Low SFE (620 mJ/m2) favors cfcc ? ehcp transformation, however, when the SFE more than 20 mJ/m2, the cfcc ? ehcp transformation will be suppressed. High manganese austenitic steels with higher SFE tend to form mechanical twin instead of phase transformation, i.e., TWIP effect takes the place of ⇑ Corresponding author at: School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China. Tel.: +86 29 88726206; fax: +86 29 88223416. E-mail address: [email protected] (D. Li). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.05.031
TRIP effect. Both TRIP mechanism and TWIP mechanism have a beneficial effect on mechanical properties of steels [1,6]. Grässel et al. [1] found that Fe–20Mn–3Si–3Al TRIP steel exhibits a high tensile strength of about 900 MPa, and the extremely high uniform elongation of about 65% at room temperature. Because of their excellent comprehensive mechanical properties, high manganese austenitic steels have a great potential to manufacture components and car body parts in automotive industry [1–4,6]. Recently, most of investigations have been focused on relationships between deformation mechanism and SFE, mechanical properties of such steels and microstructure transformation during deformation process [1–5]. It is well known that hot working is commonly applied to metals manufacturing, determination of hot deformation behavior and deformation resistance is quite important for successful production of such materials. Studies [7–9] show that hot deformation parameters such as temperature, strain and strain rates have a major influence on hot deformation behavior of metals. In order to develop the manufacturing methods, it is also necessary to investigate the constitutive relation of such steels. Moreover, the deformation parameters must be controlled properly in order to obtain ideal microstructure and mechanical properties. In this study, the effect of deformation parameters including temperature and strain rate on the flow stress of an austenitic Fe–20Mn–3Si–3Al TRIP steel have been investigated by hot compression tests. And then, the constitutive constants of this steel have been determined, and constitutive equations relating
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D. Li et al. / Materials and Design 34 (2012) 713–718
Z parameter and hot deformation activation energy Q have been derived for this steel. This can give indispensable information for optimizing hot working parameters.
2. Experimental set up The experimental steel was melted in a vacuum induction furnace under argon atmosphere, using industrial pure iron, electrolytic manganese, commercial pure aluminum and low-carbon silicoferrite, and then cast into a low-carbon steel mould of 300 mm 100 mm 60 mm. The chemical composition of the experimental steel is given in Table 1. The ingot was homogenized at 1150 °C for 4 h to eliminate the segregation of the alloying elements. Subsequently, the ingot was hot rolled to 15 mm thick plate in the temperature range from 900 °C to 1100 °C. Finally, the plate was solution-treated at 1100 °C for 1 h followed by water-quenching in order to remove any particles precipitated in the course of rolling. Cylindrical specimens for hot compression test with a diameter of 8 mm and a height of 12 mm were machined from the solutiontreated plate along the rolling direction. The compression tests were carried out on a Gleeble 3500D thermo-mechanical simulator in the temperature range from 900 °C to 1100 °C and at the strain rate range from 0.01 s1 to 10 s1. Prior to the compression, specimens were heated in vacuum at the rate of 20 °C/s to 1150 °C for
Table 1 Chemical composition of the experimental steel. C
Mn
Si
Al
S
P
N
Fe
0.04
19.87
3.06
2.62
0.007
0.008
0.008
Bal.
5 min and then cooled to the test temperature with the cooling rate of 5 °C/s. All specimens were kept at the test temperature for 3 min before compression in order to homogenize the temperature. All specimens were compressed to a true strain of 0.7, and then water quenched immediately to room temperature in order to remain the hot deformation microstructure. The deformed specimens were sectioned through the longitudinal axis. The deformed specimens were polished and then etched in 4 vol% Nital. Finally, the microstructures were examined using an Olympus PMG3 optical microscopy.
3. Results and discussion 3.1. Flow stress behavior The true stress–true strain curves obtained during hot deformation of Fe–20Mn–3Si–3Al TRIP steel at different deformation temperatures and strain rates are shown in Fig. 1. The flow stress increased at a decreasing hardening rate with increase in strain till a peak stress (rp) was reached. Beyond the peak strain (ep) which corresponds to the peak stress, the flow stress either decreased with increase in strain or reached to a steady state flow. It is well known that the hot deformation process is a competing process of the work hardening and the dynamic softening. At the onset of deformation, the work hardening exceeds the dynamic softening due to the rapid multiplication of dislocation, leading a rapid increase of flow stress. With the increasing of strain, the dynamic softening such as dynamic recovery (DRV) and dynamic recrystallization (DRX) take place in the steel, which can offset or partially offset the effect of work hardening. If the softening rate is higher than work hardening rate, the flow stress decreases gradually. But the steady state flow is reached when the softening rate
Fig. 1. True stress–true strain curves of Fe–20Mn–3Si–3Al TRIP steel at strain rates of (a) 0.01 s1, (b) 0.1 s1, (c) 1 s1 and (d) 10 s1.
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D. Li et al. / Materials and Design 34 (2012) 713–718
is equal to the work hardening rate. Furthermore, the flow stress is sensitively dependent on deformation temperature and strain rate. From Fig. 1, it can be found that at a given strain rate the flow stress level decreased with increase in deformation temperature. At a given temperature, lower the strain rate, lower was the flow stress. It means that low strain rate and high deformation temperature are in favor of the dynamic flow softening. 3.2. Constitutive equations Constitutive equations are often used to modeling the flow behavior of metallic materials during hot deformation. The relationship between the flow stress r, the deformation temperature T and the strain rate e_ can be described by several constitutive equations expressed as [10–13]
e_ ¼ A1 rn1 exp
Q RT
ð1Þ
e_ ¼ A2 expðbrÞ exp
Q RT
e_ ¼ A½sinhðarÞn exp
Q RT
ð2Þ ð3Þ
where A1, A2, A, n1, b, a, and n are the material constants which are a function of chemical composition of alloy. Q is the hot deformation activation energy, R is the universal gas constant (8.314 J/mol K), T is the absolute temperature. The power law Eq. (1) breaks down at high stress level (ar > 1.2) whereas the exponential law Eq. (2) breaks down at low stress level (ar < 0.8). The hyperbolic-sine law Eq. (3) is a universal form suitable for a wide range. The hot deformation behavior of metallic materials is represented by the Zener–Hollomon (Z) parameter which correlates the strain rate e_ ; deformation temperature T and hot deformation activation energy Q by the expression [7,12–15]:
Z ¼ e_ exp
Q RT
ð4Þ
By combining Eqs. (3) and (4), the relationship between the Z parameter and flow stress can be described by the following equation:
Z ¼ A½sinhðarÞn
ð5Þ
In this study, the peak stress rp is taken for the r term in the above equations. The stress multiplier a is an adjustable constant that brings the values of arp into the suitable range, and the value of a can be defined by the following equation [10,16–19].
a b=n1
of solute atom. Previous researches found that solute atoms of Mn, Si and Al have very weak strengthening effects [12], so the Fe– 20Mn–3Si–3Al TRIP steel ought to has a low atomic-misfit parameter just like austenitic stainless steels. The value of n1 can be determined by the plot of ln e_ vs. ln rp for five different temperatures 1173, 1223, 1273, 1323 and 1373 K, as shown in Fig. 2. It can be found that the value of n1 increases with the decrease of deformation temperature, indicating that the power law breaks down due to the decrease of temperature resulting in an increasing stress level, and the dislocation climb is no longer the rate controlling with a rising stress level, which can be attributed to the exponential dependence of strain rate on the applied stress [20,21]. In order to determine n1 and b accurately, both low and high stress data should be present at the same temperature. However, because the present data is not showing a transition in the stress exponent on double logarithmic plot for the same temperature, the value of n1 can be determined by plotting lne_ –lnr at 1100 °C, while b can be determined from lne_ – r at 900 °C, then a can be calculated. Therefore, the peak stresses rp measured at 1100 °C and 0.01 s1, 0.1 s1 and 1 s1 were chosen to determine the value of n1, and then the value of n1 was determined as 4.5 by means of linear regression analysis. Meanwhile, the value of b can be determined from the plot of ln e_ vs. rp at 900 °C, and then the value of b was taken as 0.0391. Finally, the stress multiplier a in Eq. (3) was determined as 0.0087. Taking natural logarithm and then partial derivative on the both side of Eq. (3), we have
Q RT
ð7Þ
Q 1 @ R T
ð8Þ
ln e_ ¼ ln A þ n ln½sinhðarÞ @ðln e_ Þ ¼ n@½ln sinhðarÞ
If T is a constant, Eq. (8) can be written as
n¼
@ ln e_ @ ln½sinhðarÞT
ð9Þ
And if e_ is a constant, Eq. (8) can be written as following:
Q ¼R
@ ln e_ @ ln½sinhðarÞ ¼ RnS @ð1=TÞ @ ln½sinhðarÞ T e_
ð10Þ
It can be found from Eq. (7) that there exists a linear relationship between ln e_ and ln[sinh(ar)] when T remains constant, and ln[sinh(ar)] is linear with (1000/T) when e_ remains constant, as shown in Fig. 3a and b, respectively. Consequently, the values of n and S can be determined by means of linear regression analysis. The value of n in Eq. (3) was determined as 4.4, which is close to
ð6Þ
where the stress exponent n1 is an important constant in the power law Eq. (1), and the power law is successfully used to describe the creep process. Metals and alloys can be classified according to the value of n1. For solid solution alloys with high atomic-misfit parameter, such as Al–Mg alloys and Al–Cu alloys, the values of n1 are commonly close to 3. This kind of alloys has been termed as Class I alloys [20,21]. Solute drag or viscous glide resulting from elastic interaction between dislocation and solute atoms due to the size effect is rate controlling during deformation process. For pure metals and alloys with low atomic-misfit parameter, just like austenitic stainless steels and Fe–Cr–Ni alloys, the values of n1 are close to 5. The activation energy for deformation is approximately equal to that for self-diffusion, indicating that dislocation climb is rate controlling. This kind of metals and alloys has been termed as Class II alloy [20,21]. The atomic-misfit parameter between solute atoms and matrix can be reflected through the solute strengthening effect
Fig. 2. The stress exponent n1 in power law changes with deformation temperature.
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D. Li et al. / Materials and Design 34 (2012) 713–718
Firstly, a mass of solute atoms can cause an obvious solute strengthening lead to a rising deformation resistance, and they interact with dislocations during hot deformation which makes the dislocation motion become more difficult. Secondly, because of the rapid increase of dislocation density during hot working before the peak strain ep, the interaction of dislocations can also make the dislocation motion become more difficult. Finally, the effect of solute atoms on diffusion coefficient cannot be neglected because of the high alloy content. In the regular hot working, cross slipping and dislocation climb are the main deformation mechanisms. However, under some special conditions, the grain boundary slipping (GBs) and diffusion creep can be considered as the main deformation mechanisms for hot deformation as well. For example, the ultra fine grained metal deformed at high temperature and low strain rate, the GBs and diffusion creep have a great contribution for thermoplastic deformation. At the present condition, the mechanisms of GBs and diffusion creep have limited effects on plastic deformation. In order to determine the value of A, Eq. (5) can be rewritten as,
ln Z ¼ ln A þ n ln½sinhðarÞ
ð11Þ
It can be found that lnA is the intercept of the ln Z vs. ln[sinh (ar)] plots. The values of Z parameter at various strain rates and deformation temperatures can be calculated by Eq. (4). Fig. 4 shows that there exists a good linear relationship between ln Z and ln[sinh(ar)], and the linear correlation coefficient is 0.997. The materials constants of this steel obtained from the experimental data are listed in Table 2. Taking a, A, n and Q into Eq. (3), the constitutive relationship between e_ , T and rp can be expressed as 4:4
14
e_ ¼ 1:066 10 ½sinhð0:0087 rp Þ
!
ð12Þ
Fig. 3. Plots of (a) ln e_ vs. ln[sinh (ar)] and (b) ln[sinh (ar)] vs. (1000/T) for Fe– 20Mn–3Si–3Al TRIP steel.
that of austenitic stainless steels. Previous researches indicate that the values of n for austenitic stainless steels embracing 301, 304, 316, and 317 lay in the range of 4.2–4.6 [12,13]. Q is an important material parameter serving as indicator of deformation difficulty degree in hot deformation. The value of Q for Fe–20Mn–3Si–3Al TRIP steel is 387.84 kJ/mol, which is somewhat lower than that of 304 austenitic stainless steel (about 400 kJ/mol) [14,22] and much higher than that of the common low-carbon steels [6,13]. In pervious studies, the hot deformation activation energies calculated by Hamada et al. for Fe–25Mn–3Al and Fe–25Mn–6Al high Mn austenitic steels are 397 kJ/mol and 405 kJ/mol, respectively [6]. Generally, Q is a function of the chemical composition of alloy and increases with alloy content. Comparing the value of Q obtained from this experiment with that calculated by Hamada et al., it can be found that alloy content has an important influence on hot deformation activation energy of high Mn austenitic steel. In creep, the activation energy for creep for Class II alloys equals to that for self diffusion, leading to the theory that dislocation climb is the rate-controlling mechanism, and the creep process is controlled by lattice self diffusion. But in hot working, especially for alloys undergoing DRX, the values of Q for hot working are considerably higher than those for self diffusion, and it has been difficult to associate the Q for rp with any specific mechanism [12]. In hot working, the climb of edge dislocations is still one of the most important deformation mechanisms, but the deformation rate is dependent on the applied stress. The activation energy for self diffusion of c-Fe is about 280 kJ/mol [12] which is significantly lower than the calculated Q value in this experiment. Some researchers had been found the same phenomenon in other metals [12,23,24]. This is presumably related to the following factors.
387:84 103 exp RT
And the peak stress can also be described by Z parameter as
8 1=2 ) n 1=4:4 2=4:4 > > Z Z > < rp ¼ 114:94 ln þ þ 1 1:0661014 1:0661014 > > > :
Z ¼ e_ exp
387:84103 RT
ð13Þ
Fig. 4. Relationship between Z parameters and peak stresses.
Table 2 Material constants of Fe–20Mn–3Si–3Al TRIP steel in the hyperbolic-sine equation.
a (MPa)
n
S
A (s1)
Q (kJ/mol)
0.0087
4.4
10.6
1.066 1014
387.84
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D. Li et al. / Materials and Design 34 (2012) 713–718
(b)
(a)
50µm
(c)
100µm
(d)
100µm
100µm
Fig. 5. Optical deformed microstructures under different conditions: (a) 900 °C, 10 s1, (b) 900 °C, 0.01 s1, (c) 1100 °C, 1 s1 and (d) 1100 °C, 0.01 s1.
3.3. Microstructural evolution In hot deformation process, DRX is a common characteristic of metallic materials with low and medium SFE, such as the austenitic stainless steels, copper and c-iron [7,22]. The austenitic Fe–20Mn– 3Si–3Al TRIP steel owns a low SFE as well, so DRX ought to occur during hot compression. The microstructures of deformed specimens under different conditions are shown in Fig. 5. It can be seen from Fig. 5a that at the temperature of 900 °C and strain rate of 10 s1, the original grains were elongated along the deformation direction, and the initial grain boundaries and twinning boundaries were finely serrated, and the initial twin boundaries have lost their specific twin orientation and transformed to normal high-angle grain boundaries, moreover, a few new fine grains were nucleated at the serrated part of initial grain boundaries, demonstrating occurrence of DRX. As the strain rate decreased to 0.01 s1, more new fine DRX grains emerged at prior grain boundaries, and the necklace structure started to form around the initial grain boundaries, as seen in Fig. 5b. As the deformation temperature increased to 1100 °C, the grain size as well as the volume fraction of the recrystallized grains increased dramatically, and the original grains were replaced completely by new fine DRX grains, the DRX proceeding has been completed, as seen in Fig. 5c. Fig. 5d shows the microstructure of the specimen deformed at 1100 °C and 0.01 s1. It can be seen that as the deformation temperature increased to 1100 °C and the strain rate decreased to 0.01 s1, the DRX grain size coarsened remarkably. It is well known that DRX is dependent sensitively on the deformation temperature and strain rate. Both effects of deformation temperature and strain rate on DRX procedure can be summed up in Z parameter. Decreasing of Z value, that is increasing deformation temperature and/or decreasing strain rate, lead to more adequate proceeding of DRX. 4. Conclusions Hot compression tests of Fe–20Mn–3Si–3Al TRIP steel were carried out on Gleeble 3500D thermo-mechanical simulator in the temperature range from 900 °C to 1100 °C and strain rate range from 0.01 s1 to 10 s1.
(1) The flow stress increases at a decreasing hardening rate with increase in strain till a peak stress (rp) is reached. Beyond the peak strain (ep) which corresponds to the peak stress, the flow stress either decreases with increases in strain or reached to a steady state flow. The peak stress decreases with increasing deformation temperature and decreasing strain rate. (2) The peak stress can be predicted by Z parameter in the hyperbolic sine equation with the hot deformation activation energy of 387.84 kJ/mol. (3) DRX is the main reason for dynamic flow softening of Fe– 20Mn–3Si–3Al TRIP steel. DRX is operative over a wide range of deformation temperatures and strain rates even at low temperature and high strain rate. Decreasing of Z value lead to more adequate proceeding of DRX. Acknowledgements This work is supported by National High-Technology Research and Development Project of China (No. 2006AA06A). In addition, the authors would appreciate Prof. Fengzhang Ren of Henan University of Science and Technology for preparation of the experimental steel. References [1] Grässel O, Kruger L, Frommeyer G, Meyer LW. High strength Fe–Mn–(Al, Si) TRIP/TWIP steels development–properties–application. Int J Plast 2000;16:1391–409. [2] Barbier D, Gey N, Allian S, Bozzolo N, Humbert M. Analysis of the tensile behavior of a TWIP steel based on the texture and microstructure evolutions. Mater Sci Eng A 2009;500:196–206. [3] Vercammen S, Blanpain B, De Cooman BC, Wollants P. Cold rolling behavior of an austenitic Fe–30Mn–3Si–3Al TWIP-steel: the importance of deformation twinning. Acta Mater 2004;52:2005–12. [4] Yang P, Xie Q, Meng L, Ding H, Tang Z. Dependence of deformation twinning on grain orientation in high manganese steel. Scripta Mater 2006;55:629–31. [5] Allain S, Chateau JP, Bouaziz O. Correlations between the calculated stacking fault energy and the plasticity mechanism in Fe–Mn–C alloy. Mater Sci Eng A 2004;387–389:158–62. [6] Hamada AS, Karjalainen LP, Somani MC. The influence of aluminum on hot deformation behavior and tensile properties of high-Mn TWIP steels. Mater Sci Eng A 2007;467:114–24.
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[7] Abbasi SM, Shokuhfar A. Prediction of hot deformation behaviour of 10Cr– 10Ni–5Mo–2Cu steel. Mater Lett 2007;61:2523–6. [8] Cho JR, Bae WB, Hwang WJ, Hartly P. A study on the hot deformation behavior and dynamic recrystallization of Al–5 wt.%Mg alloy. J Mater Process Technol 2010;50:227–33. [9] Zhang H, Li LX, Yuan D, Peng DS. Hot deformation behavior of the new Al–Mg– Si–Cu aluminum alloy during compression at elevated temperatures. Mater Charact 2007;58:168–73. [10] Banerjee S, Robi PS, Srinivasan A, Kumar LP. High temperature deformation behavior of Al–Cu–Mg alloys micro-alloyed with Sn. Mater Sci Eng A 2010;527:2498–503. [11] Farnoush H, Momeni A, Dehghani K, Aghazadeh Mohandesi J, Keshmiri H. Hot deformation characteristics of 2205 duplex stainless steel based on the behavior of constituent phases. Mater Des 2010;31:220–6. [12] McQueen HJ, Yue S, Ryan ND, Fry E. Hot working characteristics of steel in austenitic state. J Mater Process Technol 1995;53:293–310. [13] McQueen HJ, Ryan ND. Constitutive analysis in hot working. Mater Sci Eng A 2002;322:43–63. [14] Dehghan-Manshadi A, Barnett MR, Hodgson PD. Hot deformation and recrystallization of austenitic stainless steel: part I. Dynamic recrystallization. Metall Mater Trans 2008;39A:1359–70. [15] Ji GL, Li FG, Li QH, Li HQ, Li Z. Research on dynamic recrystallization kinetics of Aermet 100 steel. Mater Sci Eng A 2010;527:2350–5.
[16] Jin NP, Zhang H, Han Y, Wu WX, Chen JH. Hot deformation behavior of 7150 aluminum alloy during compression at elevated temperature. Mater Charact 2009;60:530–6. [17] Huang XD, Zhang H, Han Y, Wu WX, Chen JH. Hot deformation behavior of 20,026 aluminum alloy during compression at elevated temperature. Mater Sci Eng A 2010;527:485–90. [18] Mirzadeh H, Najafizadeh A, Moazeny M. Flow curve analysis of 17–4 PH stainless steel under hot compression test. Metall Mater Trans 2009;40A:2950–8. [19] McQueen HJ. Elevated-temperature deformation at forming rates of 102 to 102 s1. Metall Mater Trans 2002;33A:345–62. [20] Ei-Danaf EA, Almajid AA, Soliman MS. High-temperature deformation and ductility of a modified 5083 Al alloy. J Mater Eng Perform 2008;17:572–9. [21] Kassner ME, Perez-Prado M-T. Five-power-law creep in single phase metals and alloys. Prog Mater Sci 2000;45:1–102. [22] Dehghan-Manshadi A, Barnett MR, Hodgson PD. Recrystallization in AISI 304 austenitic stainless steel during and after hot deformation. Mater Sci Eng A 2008;485:664–72. [23] Medina SF, Hernandez CA. General expression of the Zener–Hollomon parameter as a function of the chemical composition of low alloy and microalloyed steels. Acta Mater 1996;44:137–48. [24] Avramovic-Cingara G, Perovic DD, McQueen HJ. Hot deformation mechanisms of a solution-treated Al–Li–Cu–Mg–Zr Alloy. Metall Mater Trans 1996;27A:3478–90.
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Hot deformation behavior of an austenitic Fe–20Mn–3Si–3Al transformation induced plasticity steel Dejun Li a,b,⇑, Yaorong Feng a,b, Zhifu Yin c, Fengshou Shangguan a, Ke Wang c, Qiang Liu b, Feng Hu a a
School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China CNPC Tubular Goods Research Institute, Xi’an 710065, China c Petroleum Research Institute, Shaanxi Yanchang Petroleum Group Company, Xi’an 710075, China b
a r t i c l e
i n f o
Article history: Received 31 January 2011 Accepted 18 May 2011 Available online 23 May 2011 Keywords: A. Ferrous metals and alloys C. Forming E. Physical
a b s t r a c t Hot deformation behavior of an austenitic Fe–20Mn–3Si–3Al transformation induced plasticity (TRIP) steel was investigated by hot compression tests on Gleeble 3500D thermo-mechanical simulator in the temperature ranges of 900–1100 °C and the strain rate ranges of 0.01–10 s1. The results show that the flow stress is sensitively dependent on deformation temperature and strain rate, and the flow stress increases with strain rate and decreases with deformation temperature. The peak stress during hot deformation can be predicted by the Zener–Hollomon (Z) parameter in the hyperbolic sine equation with the hot deformation activation energy Q of 387.84 kJ/mol. The dynamic recrystallization (DRX) is the most important softening mechanism for the experimental steel during hot compression. Furthermore, DRX procedure is strongly affected by Z parameter, and decreasing of Z value lead to more adequate proceeding of DRX. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The development of steels for automotive applications is focused on an increase of strength combined with the improvement of its ductility. The high strength enables car manufacturers to reduce the weight of the car bodies and improve vehicle safety, whereas the increasing ductility admits complex component design. In order to meet the needs of automotive industry, high manganese austenitic steels have gradually become a focus of research due to their outstanding mechanical properties. Outstanding mechanical properties combining a satisfactory strength and excellent ductility result from the strain-induced microstructure transformations which are called the transformation induced plasticity (TRIP) effect and the twinning induced plasticity (TWIP) effect [1–6]. The microstructure transformations, responsible for the high working hardening rate, are correlated to the stacking fault energy (SFE) of the austenitic matrix. The SFE is a function of the alloy composition and deformation temperature. Low SFE (620 mJ/m2) favors cfcc ? ehcp transformation, however, when the SFE more than 20 mJ/m2, the cfcc ? ehcp transformation will be suppressed. High manganese austenitic steels with higher SFE tend to form mechanical twin instead of phase transformation, i.e., TWIP effect takes the place of ⇑ Corresponding author at: School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China. Tel.: +86 29 88726206; fax: +86 29 88223416. E-mail address: [email protected] (D. Li). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.05.031
TRIP effect. Both TRIP mechanism and TWIP mechanism have a beneficial effect on mechanical properties of steels [1,6]. Grässel et al. [1] found that Fe–20Mn–3Si–3Al TRIP steel exhibits a high tensile strength of about 900 MPa, and the extremely high uniform elongation of about 65% at room temperature. Because of their excellent comprehensive mechanical properties, high manganese austenitic steels have a great potential to manufacture components and car body parts in automotive industry [1–4,6]. Recently, most of investigations have been focused on relationships between deformation mechanism and SFE, mechanical properties of such steels and microstructure transformation during deformation process [1–5]. It is well known that hot working is commonly applied to metals manufacturing, determination of hot deformation behavior and deformation resistance is quite important for successful production of such materials. Studies [7–9] show that hot deformation parameters such as temperature, strain and strain rates have a major influence on hot deformation behavior of metals. In order to develop the manufacturing methods, it is also necessary to investigate the constitutive relation of such steels. Moreover, the deformation parameters must be controlled properly in order to obtain ideal microstructure and mechanical properties. In this study, the effect of deformation parameters including temperature and strain rate on the flow stress of an austenitic Fe–20Mn–3Si–3Al TRIP steel have been investigated by hot compression tests. And then, the constitutive constants of this steel have been determined, and constitutive equations relating
714
D. Li et al. / Materials and Design 34 (2012) 713–718
Z parameter and hot deformation activation energy Q have been derived for this steel. This can give indispensable information for optimizing hot working parameters.
2. Experimental set up The experimental steel was melted in a vacuum induction furnace under argon atmosphere, using industrial pure iron, electrolytic manganese, commercial pure aluminum and low-carbon silicoferrite, and then cast into a low-carbon steel mould of 300 mm 100 mm 60 mm. The chemical composition of the experimental steel is given in Table 1. The ingot was homogenized at 1150 °C for 4 h to eliminate the segregation of the alloying elements. Subsequently, the ingot was hot rolled to 15 mm thick plate in the temperature range from 900 °C to 1100 °C. Finally, the plate was solution-treated at 1100 °C for 1 h followed by water-quenching in order to remove any particles precipitated in the course of rolling. Cylindrical specimens for hot compression test with a diameter of 8 mm and a height of 12 mm were machined from the solutiontreated plate along the rolling direction. The compression tests were carried out on a Gleeble 3500D thermo-mechanical simulator in the temperature range from 900 °C to 1100 °C and at the strain rate range from 0.01 s1 to 10 s1. Prior to the compression, specimens were heated in vacuum at the rate of 20 °C/s to 1150 °C for
Table 1 Chemical composition of the experimental steel. C
Mn
Si
Al
S
P
N
Fe
0.04
19.87
3.06
2.62
0.007
0.008
0.008
Bal.
5 min and then cooled to the test temperature with the cooling rate of 5 °C/s. All specimens were kept at the test temperature for 3 min before compression in order to homogenize the temperature. All specimens were compressed to a true strain of 0.7, and then water quenched immediately to room temperature in order to remain the hot deformation microstructure. The deformed specimens were sectioned through the longitudinal axis. The deformed specimens were polished and then etched in 4 vol% Nital. Finally, the microstructures were examined using an Olympus PMG3 optical microscopy.
3. Results and discussion 3.1. Flow stress behavior The true stress–true strain curves obtained during hot deformation of Fe–20Mn–3Si–3Al TRIP steel at different deformation temperatures and strain rates are shown in Fig. 1. The flow stress increased at a decreasing hardening rate with increase in strain till a peak stress (rp) was reached. Beyond the peak strain (ep) which corresponds to the peak stress, the flow stress either decreased with increase in strain or reached to a steady state flow. It is well known that the hot deformation process is a competing process of the work hardening and the dynamic softening. At the onset of deformation, the work hardening exceeds the dynamic softening due to the rapid multiplication of dislocation, leading a rapid increase of flow stress. With the increasing of strain, the dynamic softening such as dynamic recovery (DRV) and dynamic recrystallization (DRX) take place in the steel, which can offset or partially offset the effect of work hardening. If the softening rate is higher than work hardening rate, the flow stress decreases gradually. But the steady state flow is reached when the softening rate
Fig. 1. True stress–true strain curves of Fe–20Mn–3Si–3Al TRIP steel at strain rates of (a) 0.01 s1, (b) 0.1 s1, (c) 1 s1 and (d) 10 s1.
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is equal to the work hardening rate. Furthermore, the flow stress is sensitively dependent on deformation temperature and strain rate. From Fig. 1, it can be found that at a given strain rate the flow stress level decreased with increase in deformation temperature. At a given temperature, lower the strain rate, lower was the flow stress. It means that low strain rate and high deformation temperature are in favor of the dynamic flow softening. 3.2. Constitutive equations Constitutive equations are often used to modeling the flow behavior of metallic materials during hot deformation. The relationship between the flow stress r, the deformation temperature T and the strain rate e_ can be described by several constitutive equations expressed as [10–13]
e_ ¼ A1 rn1 exp
Q RT
ð1Þ
e_ ¼ A2 expðbrÞ exp
Q RT
e_ ¼ A½sinhðarÞn exp
Q RT
ð2Þ ð3Þ
where A1, A2, A, n1, b, a, and n are the material constants which are a function of chemical composition of alloy. Q is the hot deformation activation energy, R is the universal gas constant (8.314 J/mol K), T is the absolute temperature. The power law Eq. (1) breaks down at high stress level (ar > 1.2) whereas the exponential law Eq. (2) breaks down at low stress level (ar < 0.8). The hyperbolic-sine law Eq. (3) is a universal form suitable for a wide range. The hot deformation behavior of metallic materials is represented by the Zener–Hollomon (Z) parameter which correlates the strain rate e_ ; deformation temperature T and hot deformation activation energy Q by the expression [7,12–15]:
Z ¼ e_ exp
Q RT
ð4Þ
By combining Eqs. (3) and (4), the relationship between the Z parameter and flow stress can be described by the following equation:
Z ¼ A½sinhðarÞn
ð5Þ
In this study, the peak stress rp is taken for the r term in the above equations. The stress multiplier a is an adjustable constant that brings the values of arp into the suitable range, and the value of a can be defined by the following equation [10,16–19].
a b=n1
of solute atom. Previous researches found that solute atoms of Mn, Si and Al have very weak strengthening effects [12], so the Fe– 20Mn–3Si–3Al TRIP steel ought to has a low atomic-misfit parameter just like austenitic stainless steels. The value of n1 can be determined by the plot of ln e_ vs. ln rp for five different temperatures 1173, 1223, 1273, 1323 and 1373 K, as shown in Fig. 2. It can be found that the value of n1 increases with the decrease of deformation temperature, indicating that the power law breaks down due to the decrease of temperature resulting in an increasing stress level, and the dislocation climb is no longer the rate controlling with a rising stress level, which can be attributed to the exponential dependence of strain rate on the applied stress [20,21]. In order to determine n1 and b accurately, both low and high stress data should be present at the same temperature. However, because the present data is not showing a transition in the stress exponent on double logarithmic plot for the same temperature, the value of n1 can be determined by plotting lne_ –lnr at 1100 °C, while b can be determined from lne_ – r at 900 °C, then a can be calculated. Therefore, the peak stresses rp measured at 1100 °C and 0.01 s1, 0.1 s1 and 1 s1 were chosen to determine the value of n1, and then the value of n1 was determined as 4.5 by means of linear regression analysis. Meanwhile, the value of b can be determined from the plot of ln e_ vs. rp at 900 °C, and then the value of b was taken as 0.0391. Finally, the stress multiplier a in Eq. (3) was determined as 0.0087. Taking natural logarithm and then partial derivative on the both side of Eq. (3), we have
Q RT
ð7Þ
Q 1 @ R T
ð8Þ
ln e_ ¼ ln A þ n ln½sinhðarÞ @ðln e_ Þ ¼ n@½ln sinhðarÞ
If T is a constant, Eq. (8) can be written as
n¼
@ ln e_ @ ln½sinhðarÞT
ð9Þ
And if e_ is a constant, Eq. (8) can be written as following:
Q ¼R
@ ln e_ @ ln½sinhðarÞ ¼ RnS @ð1=TÞ @ ln½sinhðarÞ T e_
ð10Þ
It can be found from Eq. (7) that there exists a linear relationship between ln e_ and ln[sinh(ar)] when T remains constant, and ln[sinh(ar)] is linear with (1000/T) when e_ remains constant, as shown in Fig. 3a and b, respectively. Consequently, the values of n and S can be determined by means of linear regression analysis. The value of n in Eq. (3) was determined as 4.4, which is close to
ð6Þ
where the stress exponent n1 is an important constant in the power law Eq. (1), and the power law is successfully used to describe the creep process. Metals and alloys can be classified according to the value of n1. For solid solution alloys with high atomic-misfit parameter, such as Al–Mg alloys and Al–Cu alloys, the values of n1 are commonly close to 3. This kind of alloys has been termed as Class I alloys [20,21]. Solute drag or viscous glide resulting from elastic interaction between dislocation and solute atoms due to the size effect is rate controlling during deformation process. For pure metals and alloys with low atomic-misfit parameter, just like austenitic stainless steels and Fe–Cr–Ni alloys, the values of n1 are close to 5. The activation energy for deformation is approximately equal to that for self-diffusion, indicating that dislocation climb is rate controlling. This kind of metals and alloys has been termed as Class II alloy [20,21]. The atomic-misfit parameter between solute atoms and matrix can be reflected through the solute strengthening effect
Fig. 2. The stress exponent n1 in power law changes with deformation temperature.
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Firstly, a mass of solute atoms can cause an obvious solute strengthening lead to a rising deformation resistance, and they interact with dislocations during hot deformation which makes the dislocation motion become more difficult. Secondly, because of the rapid increase of dislocation density during hot working before the peak strain ep, the interaction of dislocations can also make the dislocation motion become more difficult. Finally, the effect of solute atoms on diffusion coefficient cannot be neglected because of the high alloy content. In the regular hot working, cross slipping and dislocation climb are the main deformation mechanisms. However, under some special conditions, the grain boundary slipping (GBs) and diffusion creep can be considered as the main deformation mechanisms for hot deformation as well. For example, the ultra fine grained metal deformed at high temperature and low strain rate, the GBs and diffusion creep have a great contribution for thermoplastic deformation. At the present condition, the mechanisms of GBs and diffusion creep have limited effects on plastic deformation. In order to determine the value of A, Eq. (5) can be rewritten as,
ln Z ¼ ln A þ n ln½sinhðarÞ
ð11Þ
It can be found that lnA is the intercept of the ln Z vs. ln[sinh (ar)] plots. The values of Z parameter at various strain rates and deformation temperatures can be calculated by Eq. (4). Fig. 4 shows that there exists a good linear relationship between ln Z and ln[sinh(ar)], and the linear correlation coefficient is 0.997. The materials constants of this steel obtained from the experimental data are listed in Table 2. Taking a, A, n and Q into Eq. (3), the constitutive relationship between e_ , T and rp can be expressed as 4:4
14
e_ ¼ 1:066 10 ½sinhð0:0087 rp Þ
!
ð12Þ
Fig. 3. Plots of (a) ln e_ vs. ln[sinh (ar)] and (b) ln[sinh (ar)] vs. (1000/T) for Fe– 20Mn–3Si–3Al TRIP steel.
that of austenitic stainless steels. Previous researches indicate that the values of n for austenitic stainless steels embracing 301, 304, 316, and 317 lay in the range of 4.2–4.6 [12,13]. Q is an important material parameter serving as indicator of deformation difficulty degree in hot deformation. The value of Q for Fe–20Mn–3Si–3Al TRIP steel is 387.84 kJ/mol, which is somewhat lower than that of 304 austenitic stainless steel (about 400 kJ/mol) [14,22] and much higher than that of the common low-carbon steels [6,13]. In pervious studies, the hot deformation activation energies calculated by Hamada et al. for Fe–25Mn–3Al and Fe–25Mn–6Al high Mn austenitic steels are 397 kJ/mol and 405 kJ/mol, respectively [6]. Generally, Q is a function of the chemical composition of alloy and increases with alloy content. Comparing the value of Q obtained from this experiment with that calculated by Hamada et al., it can be found that alloy content has an important influence on hot deformation activation energy of high Mn austenitic steel. In creep, the activation energy for creep for Class II alloys equals to that for self diffusion, leading to the theory that dislocation climb is the rate-controlling mechanism, and the creep process is controlled by lattice self diffusion. But in hot working, especially for alloys undergoing DRX, the values of Q for hot working are considerably higher than those for self diffusion, and it has been difficult to associate the Q for rp with any specific mechanism [12]. In hot working, the climb of edge dislocations is still one of the most important deformation mechanisms, but the deformation rate is dependent on the applied stress. The activation energy for self diffusion of c-Fe is about 280 kJ/mol [12] which is significantly lower than the calculated Q value in this experiment. Some researchers had been found the same phenomenon in other metals [12,23,24]. This is presumably related to the following factors.
387:84 103 exp RT
And the peak stress can also be described by Z parameter as
8 1=2 ) n 1=4:4 2=4:4 > > Z Z > < rp ¼ 114:94 ln þ þ 1 1:0661014 1:0661014 > > > :
Z ¼ e_ exp
387:84103 RT
ð13Þ
Fig. 4. Relationship between Z parameters and peak stresses.
Table 2 Material constants of Fe–20Mn–3Si–3Al TRIP steel in the hyperbolic-sine equation.
a (MPa)
n
S
A (s1)
Q (kJ/mol)
0.0087
4.4
10.6
1.066 1014
387.84
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(b)
(a)
50µm
(c)
100µm
(d)
100µm
100µm
Fig. 5. Optical deformed microstructures under different conditions: (a) 900 °C, 10 s1, (b) 900 °C, 0.01 s1, (c) 1100 °C, 1 s1 and (d) 1100 °C, 0.01 s1.
3.3. Microstructural evolution In hot deformation process, DRX is a common characteristic of metallic materials with low and medium SFE, such as the austenitic stainless steels, copper and c-iron [7,22]. The austenitic Fe–20Mn– 3Si–3Al TRIP steel owns a low SFE as well, so DRX ought to occur during hot compression. The microstructures of deformed specimens under different conditions are shown in Fig. 5. It can be seen from Fig. 5a that at the temperature of 900 °C and strain rate of 10 s1, the original grains were elongated along the deformation direction, and the initial grain boundaries and twinning boundaries were finely serrated, and the initial twin boundaries have lost their specific twin orientation and transformed to normal high-angle grain boundaries, moreover, a few new fine grains were nucleated at the serrated part of initial grain boundaries, demonstrating occurrence of DRX. As the strain rate decreased to 0.01 s1, more new fine DRX grains emerged at prior grain boundaries, and the necklace structure started to form around the initial grain boundaries, as seen in Fig. 5b. As the deformation temperature increased to 1100 °C, the grain size as well as the volume fraction of the recrystallized grains increased dramatically, and the original grains were replaced completely by new fine DRX grains, the DRX proceeding has been completed, as seen in Fig. 5c. Fig. 5d shows the microstructure of the specimen deformed at 1100 °C and 0.01 s1. It can be seen that as the deformation temperature increased to 1100 °C and the strain rate decreased to 0.01 s1, the DRX grain size coarsened remarkably. It is well known that DRX is dependent sensitively on the deformation temperature and strain rate. Both effects of deformation temperature and strain rate on DRX procedure can be summed up in Z parameter. Decreasing of Z value, that is increasing deformation temperature and/or decreasing strain rate, lead to more adequate proceeding of DRX. 4. Conclusions Hot compression tests of Fe–20Mn–3Si–3Al TRIP steel were carried out on Gleeble 3500D thermo-mechanical simulator in the temperature range from 900 °C to 1100 °C and strain rate range from 0.01 s1 to 10 s1.
(1) The flow stress increases at a decreasing hardening rate with increase in strain till a peak stress (rp) is reached. Beyond the peak strain (ep) which corresponds to the peak stress, the flow stress either decreases with increases in strain or reached to a steady state flow. The peak stress decreases with increasing deformation temperature and decreasing strain rate. (2) The peak stress can be predicted by Z parameter in the hyperbolic sine equation with the hot deformation activation energy of 387.84 kJ/mol. (3) DRX is the main reason for dynamic flow softening of Fe– 20Mn–3Si–3Al TRIP steel. DRX is operative over a wide range of deformation temperatures and strain rates even at low temperature and high strain rate. Decreasing of Z value lead to more adequate proceeding of DRX. Acknowledgements This work is supported by National High-Technology Research and Development Project of China (No. 2006AA06A). In addition, the authors would appreciate Prof. Fengzhang Ren of Henan University of Science and Technology for preparation of the experimental steel. References [1] Grässel O, Kruger L, Frommeyer G, Meyer LW. High strength Fe–Mn–(Al, Si) TRIP/TWIP steels development–properties–application. Int J Plast 2000;16:1391–409. [2] Barbier D, Gey N, Allian S, Bozzolo N, Humbert M. Analysis of the tensile behavior of a TWIP steel based on the texture and microstructure evolutions. Mater Sci Eng A 2009;500:196–206. [3] Vercammen S, Blanpain B, De Cooman BC, Wollants P. Cold rolling behavior of an austenitic Fe–30Mn–3Si–3Al TWIP-steel: the importance of deformation twinning. Acta Mater 2004;52:2005–12. [4] Yang P, Xie Q, Meng L, Ding H, Tang Z. Dependence of deformation twinning on grain orientation in high manganese steel. Scripta Mater 2006;55:629–31. [5] Allain S, Chateau JP, Bouaziz O. Correlations between the calculated stacking fault energy and the plasticity mechanism in Fe–Mn–C alloy. Mater Sci Eng A 2004;387–389:158–62. [6] Hamada AS, Karjalainen LP, Somani MC. The influence of aluminum on hot deformation behavior and tensile properties of high-Mn TWIP steels. Mater Sci Eng A 2007;467:114–24.
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