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Deformation-induced ferrite transformation in a low carbon Nb–Ti microalloyed steel
Materials & Design Materials and Design 28 (2007) 1021–1026 www.elsevier.com/locate/matdes
Short communication
Deformation-induced ferrite transformation in a low carbon Nb–Ti microalloyed steel B. Eghbali, A. Abdollah-Zadeh
*
Department of Materials Engineering, Tarbiat Modarres University, P.O. Box 14115-143, Tehran, Iran Received 15 April 2005; accepted 8 November 2005 Available online 27 December 2005
Abstract Laboratory thermomechanical processing (TMP) using single pass hot compression experiments on a low carbon Nb–Ti microalloyed steel have been carried out to study the effects of strain and strain rate on deformation-induced ferrite production through strain-induced transformation (SIT) at a temperature just above Ar3. The results show that the occurrence of SIT during deformation at 845 °C causes a dynamic softening behaviour in hot flow curves. With increasing strain from 0.25 to 0.8, higher volume fractions of strain-induced ferrite grains are obtained. The amount of strain rate plays an important role in the occurrence of SIT. Increasing the strain rate from 0.001 to 0.1 s 1 leads to an increase in the amount of minimum strain required for the initiation of SIT. Moreover, deformation to a strain of 0.8 with a strain rate of 0.1 s 1 produces equiaxed and homogeneous ultrafine ferrite grains of about 2 lm. Ó 2005 Elsevier Ltd. All rights reserved.
1. Introduction The benefits of grain refinement of steels have been well known for a long time. There are several routes to a very fine ferrite grain microstructure [1–4]. It has been demonstrated in a variety of steel compositions that ultrafine ferrite (UFF) can be produced using the so-called straininduced transformation (SIT) process [5–7]. Hodgson et al. [8] noted that a high level of undercooling as well as a large amount of strain was required to achieve strain-induced dynamic transformation, resulting in UFF (about 1 lm). They postulated three critical factors for the formation of UFF grains during SIT as follows: (i) a high shear strain; (ii) a heavy undercooling; (iii) an appropriate deformation temperature. All these factors are very important for obtaining equiaxed, homogeneous and UFF grains. Very fine ferrite grain sizes have been achieved via SIT in plain carbon steels, as confirmed by optical microstructural observation [9–13]. However, there is very
*
Corresponding author. Tel.: +98 21 8011 001x3347; fax: +98 21 8005 040. E-mail address: [email protected] (A. Abdollah-Zadeh). 0261-3069/$ - see front matter Ó 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2005.11.006
little information concerning ferrite grain refinement by SIT in the case of microalloyed steels. Furthermore, it is well known that hot deformation has an important effect not only on the microstructural changes but also on the hot flow behaviour of the material. Therefore, in addition to optical microstructural observation, hot flow stress– strain curves can be used as an evidence to determine whether SIT has been occurred during the deformation or not. In the present research, isothermal hot compression tests were conducted on a low carbon Nb–Ti microalloyed steel at constant deformation temperature, just above Ar3, to study the effects of strain and strain rate on the ferrite grain refinement. In particular, attention has been paid to study ferrite grain refinement through SIT. 2. Experimental procedures The material used in the experiments is a low carbon Nb–Ti microalloyed steel of the chemical composition shown in Table 1. The steel was prepared as 35 kg ingot in an induction furnace operating under argon atmosphere, and then refined by electro-slag remelting (ESR) in a
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B. Eghbali, A. Abdollah-Zadeh / Materials and Design 28 (2007) 1021–1026
Table 1 Chemical composition of steel (wt%) C
Si
Mn
P
S
Nb
Ti
Al
N
0.032
0.15
0.74
0.009
0.007
0.014
0.013
0.028
0.0031
laboratory unit. The ingot was reheated to 1250 °C for 1 h and hot rolled in six passes to 25 mm thick plate. Differential scanning calorimetry (DSC) was used to measure the critical transformation temperatures. The Ar1 and Ar3 temperatures were thus found to be 751 and 837 °C, respectively. Cylindrical compression samples were machined out from hot rolled plate. The deformation tests were carried out according to the schedule in Fig. 1. The samples were 18 mm in length and 12 mm in diameter, with the axis aligned in the rolling direction of the plate. The height-todiameter ratio of 1.5 was selected to ensure homogeneous deformation. Care was exercised to minimize friction between the test dies and the sample surface by machining flat-bottomed grooves on the end faces of samples. Graphite powders and thin pieces of mica sheet were used as lubricants in compression, resulting in fairly uniform deformation with negligibly small barreling. The uniaxial compression tests were performed on a servo-hydraulic 600 kN computerized Materials Testing System (MTS, Model 8500) equipped with a resistant furnace. During the test, the specimen temperature was measured and recorded using a Chromel–Alumel (Type K) thermocouple which was placed in contact with the surface of the compression specimen at the mid length position. Prior to deformation, the specimens were solutionized at 1150 °C for 5 min. The solution temperature was selected according to the solubility product of Ti and Nb precipitates [14,15]. At this temperature, Nb is completely dissolved and Ti is partially precipitated in the form of nitrides; very high temperatures, even in excess of melting point, being necessary for its total dissolution [16,17]. After solutionizing, the samples were cooled at a rate of 5 °C/s to
1150°C/ 5 min
deformation temperature of 845 °C, and held for 20 s to homogenize the temperature within the samples. Then samples were isothermally deformed with single pass strain values of 0.25, 0.4, and 0.8 at different constant strain rates of 0.001, 0.01, and 0.1 s 1. All specimens were water quenched in 2 s after deformation. Optical microscopy was conducted on mid-plane sections containing the axis of compression, in order to study the microstructural changes. 3. Results and discussion 3.1. True stress–true strain curves and related dynamic softening processes Fig. 2 shows the true stress–true strain curves for the specimens deformed up to a strain of 0.8 at 845 °C, i.e. just above Ar3, with different constant strain rates. As it can be seen in the figure, for the strain rate of 0.001 s 1, at first, the stress increases rapidly with increasing the strain. Then the stress further increases but it tends to be slow and exhibit a transition from work hardening to work hardening plus dynamic restoration. This indicates that the deformation energy is partly released in this stage due to dynamic softening occurrence. This behaviour might be interpreted as indicating dynamic recovery (DRC) of austenite. However, deformation at a higher temperature, i.e. 950 °C, exhibited the work hardening behaviour for this steel [18]. Therefore, DRC of austenite is ruled out. There are three mechanisms to release the deformation energy in deformed austenite during deformation: (i) dynamic recrystallization (DRX), (ii) DRC, or (iii) SIT of deformed austenite to ferrite [6,9]. By ruling out the first two
ε = 0.25, 0.4, 0.8 ε° = 0.001, 0.01, 0.1 s-1
210 0.1 s-1
Stress (Mpa)
180
Temperature
5°C/Sec
845°C Ar3 = 837°C
150 120 0.01 s-1
90
0.001 s-1
60 30
Ar1 = 751°C
0 0
0.2
0.4
0.6
0.8
1
Strain
Time Fig. 1. Schematic representation of the thermomechanical processing conditions.
Fig. 2. Flow curves of low carbon Nb–Ti microalloyed steel obtained under various strain rate at deformation temperature 847 °C with strain of 0.8.
B. Eghbali, A. Abdollah-Zadeh / Materials and Design 28 (2007) 1021–1026
1023
possibilities, thus, the softening behaviour observed at a strain rate of 0.001 s 1 can be attributed to the occurrence of SIT of austenite to ferrite. This is also in agreement with the observations of other workers studying hot deformation of plain carbon steels [6,11,13]. Yang and Wang [19] showed that when the deformation was given at lower temperature, DRX of austenite would not happen in plain low carbon steel during deformation, and the deformation energy can be fully utilized to strain-induced ferrite transformation. Then the driving force for phase transformation of austenite to ferrite is further improved. As shown in Fig. 2, increasing the strain rate from 0.001 to 0.1 s 1 increases the level of flow stress. In addition, at a higher strain rate, i.e. 0.1 s 1, there is no evidence of dynamic softening during deformation and the stress– strain curve indicates only the work hardening behaviour. It means that the amount of strain rate plays a very important role in occurrence of SIT during deformation by controlling the holding time and activation energy required for SIT, when the specimen is given same strain value. Increasing of strain rate leads to an increase in deformation resistance, and this is due to the fact that kinetic of dynamic softening events are diminished. With increasing the strain rate, the time for dynamic softening is diminished and the nucleation rate decreases so that the dynamic softening is inhibited. Increasing flow stress level would be expected to increase the dislocation density and decrease the subgrain size and hence increase the stored energy.
3.2. Evolution of strain-induced ferrite 3.2.1. Effects of strain Optical micrograph illustrating the evolution of straininduced ferrite grains with the amount of strain at deformation temperature of 847 °C, just above Ar3, and strain rate of 0.001 s 1 is presented in Fig. 3. Three specimens were compressed at different strains to study the effects of strain on the ferrite grain refinement. As shown in this figure, the deformed and elongated grains are present for the strain of 0.25. The fine ferrite grains had begun to nucleate inside a few of the austenite grains, and undeformed austenite grains have been statically transformed to coarse ferrite during cooling stage after deformation. This has resulted a bimodal distribution of fine ferrite grains due to dynamic transformation, and coarse ferrite grains due to static transformation, in the final microstructure. It has been demonstrated that the principal reason for refining the ferrite structure is the nucleation rate of the new phase, which is a result of increasing dislocation density in austenite just before the start of phase transformation [20]. With increasing the strain from 0.25 to 0.4, very fine ferrite grains are produced (Fig. 3). These grains have nucleated on the austenite grain boundaries, the vicinity of austenite grain boundaries and within the austenite grains. With increasing strain from 0.25 to 0.4, the extent of intra-
Fig. 3. Effect of strain on the evolution of ferrite grains refinement at deformation temperature of 845 °C with strain rate of 0.001 s 1.
granular ferrite nucleation increases. However, some ferrite grains in this microstructure are still visibly larger. The large strain produces more deformation defects such as deformation bands. These defects are beneficial in increasing the preferential nucleation sites of ferrite [21]. The austenite to ferrite transformation depends on the stored energy produced by deformation leading to the destabilization of austenite and acceleration of the austenite to ferrite transformation. The nucleation probability largely increases with increasing of strain and this leads to the very fine strain-induced ferrite grains. It can be seen in Fig. 3 that with increasing the strain to 0.8, the equiaxed and homogeneous ferrite grains of about
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4 lm with little amount of proeutectoid-deformed ferrite are obtained. The existence of little amount of proeutectoid-deformed ferrite in this microstructure is due to this fact that the mechanical working increases internal energy and thus free energy of system. This leads to the facts that Ar3 temperature is increased, and then the transformation of austenite to ferrite could happen at a higher temperature, as suggested by other workers [19,21]. It has been proposed that the austenite decomposition kinetics is influenced by its deformation state [9]: (i) the higher internal energy of the deformed and thus less stable austenite; (ii) the larger number of nucleation sites provided by defects. These two facts, rapid nucleation followed by early impingement and random orientation distribution, are believed to be the main reasons for the grain refinement by dynamic transformation. The straininduced ferrite is precipitated at the original grain boundaries of austenite or at the interface between the remaining austenite and proeutectoid ferrite, and at deformation bands and then has separated out all over the matrix. Once ferrite nuclei are formed, they will grow very rapidly because the diffusion rate also greatly increases due to high dislocation density, and this will make the ferrite grains to impinge on each other at the very early stage of grain growth. The results show that increasing the strain value has a positive effect on strain-induced ferrite formation. Strain increases diffusivity, number of nucleation site, and driving force for transformation which individually have the effects of increasing nucleation rate [6]. This is because of the increase of deformation energy as well as the increase of transformation time resulted from increase of strain value. As can be seen in Fig. 3, small strain, 0.25, cannot considerably produce strain-induced ferrite grains. However, the larger strain, 0.8, has produced higher volume fraction of strain-induced fine ferrite grains. It has been reported [22] that the critical strain value for initiation of strain-induced transformation is generally 0.3.
3.2.2. Effect of strain rate Evolution of strain-induced ferrite grains in the specimens deformed at 845 °C up to strain of 0.8 with various strain rates are shown in Fig. 4. It can be seen in this figure that the dominant microstructure is equiaxed ferrite grains with small amounts of deformed pro-eutectoid ferrite for strain rates of 0.001 and 0.01 s 1. The presence of somewhat elongated ferrite after deformation explains that phase transformation has been occurred during deformation. Increasing of strain rate from 0.001 to 0.01 s 1 reduces the ferrite grain size from about 4 to 3 lm, respectively, and decreases the amount of deformed pro-eutectoid ferrite. In other words, decreasing the deformation rate has lead to the formation of warm-worked pro-eutectoid ferrite grains. It should be noticed that the ferrite grains for these two strain rates have been produced during the SIT as discussed previously in Section 3.1. In other words, these microstructures have not been formed from other dynamic
Fig. 4. Effect of strain rate on the evolution of strain-induced ferrite formation at deformation temperature of 845 °C with strain of 0.8.
restoration processes such as DRC. The dynamic recovered structure mainly consists of pancaked grains of which elongation direction are perpendicular to the compression axis while the SIT structure has polygonal or equiaxed grains for the most parts [6,19]. This means that austenite to ferrite transformation for strain rates of 0.001 and 0.01 s 1 is occurred dynamically during deformation similar to DRC, but by SIT. As shown in Fig. 4, the microstructure of the specimen deformed at strain rate of 0.1 s 1 consists of ultrafine ferrite grains of about 2 lm. As can be seen from the corresponding flow curve in Fig. 2, there is not any evidence of dynamic softening during deformation for this strain
B. Eghbali, A. Abdollah-Zadeh / Materials and Design 28 (2007) 1021–1026
rate. Therefore, the presence of very fine ferrite grains, 2 lm, and additionally, the absence of any elongated ferrite grains in the final microstructure could be an evidence of that the SIT at strain rate of 0.1 s 1 may be enhanced compared to that at lower strain rates as mechanical driving force, by more stored deformation energy, is expected to be higher at higher strain rates. It can be concluded that strain rate plays a very important role in SIT by controlling the holding time of SIT and activation energy required by SIT, when the specimen is given same deformation value. Low deformation temperature, near to Ar3, with high strain rate enables the austenite to store much more deformation energy in SIT processing than in conventional TMP. The high stored energy significantly increases the driving force for austenite to ferrite transformation and leads to the formation of very fine ferrite grains. As mentioned above, SIT is a diffusional transformation. Therefore, increasing the strain rate has a negative effect on SIT during deformation. On the other hand, decreasing the strain rate will increase the deformation time, and there would be enough time for the transformation to take place during deformation. If strain rate is very low, the strain-induced ferrite grains will have enough time to grow, and if the strain rate is very high, the strain-induced ferrite amount would be small. So, there would be an upper limit of strain rate for the SIT at a fixed temperature, and there would be an optimum strain rate range for the application of SIT in the ferrite refinement. Further investigation is needed to study the influence of precipitation of microalloying elements on SIT kinetics.
4. Conclusions The isothermal single pass hot compression tests were conducted in a low carbon Nb–Ti microalloyed steel at a constant temperature of 845 °C, just above Ar3, to study the effects of strain and strain rate on the evolution of strain-induced ferrite grain refinement through straininduced transformation. The following main conclusions can be drawn. (1) The occurrence of the SIT during deformation at 845 °C, just above Ar3, causes a dynamic softening behaviour for strain rates of 0.001 and 0.01 s 1, as reflected in the corresponding hot flow curves. (2) With increasing the strain from 0.25 to 0.8, at strain rate of 0.001 s 1, high volume fraction of straininduced ferrite with small grain size of 4 lm is produced due to an increase in deformation energy as well as an increase in transformation time. (3) Very fine ferrite grains of about 4 and 3 lm with little amount of proeutectoid-deformed ferrite are obtained by deformation at strain rates of 0.001 and 0.01 s 1, respectively, for the strain of 0.8. This has been attributed to SIT mechanism.
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(4) Deformation at 845 °C with the strain of 0.8 and strain rate of 0.1 s 1 produces the ultrafine, equiaxed and homogeneous ferrite grains of about 2 lm with little amount of proeutectoid-deformed ferrite.
Acknowledgements The authors express their gratitude for the financial support of this research by SASAD-Iran. Experiments on MTS thermomechanical simulator were performed at Maham Science and Technology Research Center, which is gratefully acknowledged. Professor H. Assadi is gratefully acknowledged for reviewing the manuscript. References [1] Sunghak L, Dongil K, Young KL, Ohjoon K. Transformation strengthening by thermomechanical treatments in C–Mn–Ni–Nb steels. Metall Mater Trans A 1995;26A:1093. [2] Hurley PJ, Hodgson PD, Muddle BC. Analysis and characterisation of ultra-fine ferrite produced during a new steel strip rolling process. Scr Mater 1999;40:433. [3] Priestner R, Al-Horr YM, Ibraheem AK. Effect of strain on formation of ultrafine ferrite in surface of hot rolled microalloyed steel. Mater Sci Technol 2002;18:973. [4] Hodgson PD, Hickson MR, Gibbs RK. The production and mechanical properties of ultrafine ferrite. Mater Sci Forum 1998;284–286:63. [5] Hickson MR, Hodgson PD. Effect of preroll quenching and post-roll quenching on production and properties of ultrafine ferrite in steel. Mater Sci Technol 1999;15:85. [6] Chio JK, Seo DH, Lee JS, Um KK, Choo WY. Formation of ultrafine ferrite by strain-induced dynamic transformation in plain low carbon steel. ISIJ Int 2003;43:746. [7] Eghbali B, Abdollah-zadeh A. The influence of thermomechanical parameters in ferrite grain refinement in a low carbon Nb-microalloyed steel. Scr Mater 2005;53:41. [8] Hodgson PD, Hickson MR, Gibbs PK. Ultrafine ferrite in low carbon steel. Scr Metall Mater 1999;40:1179. [9] Santos DB, Bruzszek RK, Rodrigues PCM, Pereloma EV. Formation of ultra-fine ferrite microstructure in warm rolled and annealed C– Mn steel. Mater Sci Eng A 2003;A346:189. [10] Huang CJ, Li DZ, Li YY. A finite element analysis of strain induced transformation rolling and an experimental study on the grain refinement potential of severe undercooling thermo-mechanical treatment. Mater Sci Eng A 2003;A352:136. [11] Huang YD, Froyen L. Important factors to obtain homogeneous and ultrafine ferrite–pearlite microstructure in low carbon steel. Mater Proc Technol 2002;124:216. [12] Kelly GL, Beladi H, Hodgson PD. Ultrafine grained ferrite formed by interrupted hot torsion deformation of plain carbon steel. ISIJ Int 2002;42:1585. [13] Hong SC, Lim SH, Lee KJ, Shin DH, Lee KS. Effect of undercooling of austenite on strain induced ferrite transformation behaviour. ISIJ Int 2003;43:394. [14] Ghosh A, Das S, Chatterjee S, Mishra B, Rao RR. Influence of thermo-mechanical processing and different post-cooling techniques on structure and properties of an ultra low carbon Cu bearing HSLA forging. Mater Sci Eng A 2003;A348:299. [15] Cho SH, Kang KB, Jonas JJ. The dynamic, static and metadynamic recrystallization of a nb-microalloyed steel. ISIJ Int 2001;41:63. [16] Manohar PA, Dunne DP, Chandra T, Killmore CR. Grain growth predictions in microalloyed steels. ISIJ Int 1996;36:194.
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[17] Medina SF, Mancilla JE. Influence of alloying elements in solution on static recrystallization kinetics of hot deformed steels. ISIJ Int 1996;36:1063. [18] Eghbali B, Abdollah-zadeh A. Strain induced transformation in a low carbon microalloyed steel during hot compression testing. Scr Mater [in press]. [19] Yang Z, Wang R. Formation of ultra-fine grain structure of plain low carbon steel through deformation induced ferrite transformation. ISIJ Int 2003;43:761.
[20] Mesplont C, DeCooman BC. Effect of austenite deformation on crystallographic texture during transformations in microalloyed bainitic steel. Mater Sci Technol 2003;19:875. [21] Huang YD, Yang WY, Sun ZQ. Formation of ultrafine grained ferrite in low carbon steel by heavy deformation in ferrite or dual phase region. Mater Proc Technol 2003;134:19. [22] DU L, Zhang C, Ding H, Liu X, Wang G. Determination of upper limit temperature of strain-induced transformation of low carbon steels. ISIJ Int 2002;42:1119.
Short communication
Deformation-induced ferrite transformation in a low carbon Nb–Ti microalloyed steel B. Eghbali, A. Abdollah-Zadeh
*
Department of Materials Engineering, Tarbiat Modarres University, P.O. Box 14115-143, Tehran, Iran Received 15 April 2005; accepted 8 November 2005 Available online 27 December 2005
Abstract Laboratory thermomechanical processing (TMP) using single pass hot compression experiments on a low carbon Nb–Ti microalloyed steel have been carried out to study the effects of strain and strain rate on deformation-induced ferrite production through strain-induced transformation (SIT) at a temperature just above Ar3. The results show that the occurrence of SIT during deformation at 845 °C causes a dynamic softening behaviour in hot flow curves. With increasing strain from 0.25 to 0.8, higher volume fractions of strain-induced ferrite grains are obtained. The amount of strain rate plays an important role in the occurrence of SIT. Increasing the strain rate from 0.001 to 0.1 s 1 leads to an increase in the amount of minimum strain required for the initiation of SIT. Moreover, deformation to a strain of 0.8 with a strain rate of 0.1 s 1 produces equiaxed and homogeneous ultrafine ferrite grains of about 2 lm. Ó 2005 Elsevier Ltd. All rights reserved.
1. Introduction The benefits of grain refinement of steels have been well known for a long time. There are several routes to a very fine ferrite grain microstructure [1–4]. It has been demonstrated in a variety of steel compositions that ultrafine ferrite (UFF) can be produced using the so-called straininduced transformation (SIT) process [5–7]. Hodgson et al. [8] noted that a high level of undercooling as well as a large amount of strain was required to achieve strain-induced dynamic transformation, resulting in UFF (about 1 lm). They postulated three critical factors for the formation of UFF grains during SIT as follows: (i) a high shear strain; (ii) a heavy undercooling; (iii) an appropriate deformation temperature. All these factors are very important for obtaining equiaxed, homogeneous and UFF grains. Very fine ferrite grain sizes have been achieved via SIT in plain carbon steels, as confirmed by optical microstructural observation [9–13]. However, there is very
*
Corresponding author. Tel.: +98 21 8011 001x3347; fax: +98 21 8005 040. E-mail address: [email protected] (A. Abdollah-Zadeh). 0261-3069/$ - see front matter Ó 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2005.11.006
little information concerning ferrite grain refinement by SIT in the case of microalloyed steels. Furthermore, it is well known that hot deformation has an important effect not only on the microstructural changes but also on the hot flow behaviour of the material. Therefore, in addition to optical microstructural observation, hot flow stress– strain curves can be used as an evidence to determine whether SIT has been occurred during the deformation or not. In the present research, isothermal hot compression tests were conducted on a low carbon Nb–Ti microalloyed steel at constant deformation temperature, just above Ar3, to study the effects of strain and strain rate on the ferrite grain refinement. In particular, attention has been paid to study ferrite grain refinement through SIT. 2. Experimental procedures The material used in the experiments is a low carbon Nb–Ti microalloyed steel of the chemical composition shown in Table 1. The steel was prepared as 35 kg ingot in an induction furnace operating under argon atmosphere, and then refined by electro-slag remelting (ESR) in a
1022
B. Eghbali, A. Abdollah-Zadeh / Materials and Design 28 (2007) 1021–1026
Table 1 Chemical composition of steel (wt%) C
Si
Mn
P
S
Nb
Ti
Al
N
0.032
0.15
0.74
0.009
0.007
0.014
0.013
0.028
0.0031
laboratory unit. The ingot was reheated to 1250 °C for 1 h and hot rolled in six passes to 25 mm thick plate. Differential scanning calorimetry (DSC) was used to measure the critical transformation temperatures. The Ar1 and Ar3 temperatures were thus found to be 751 and 837 °C, respectively. Cylindrical compression samples were machined out from hot rolled plate. The deformation tests were carried out according to the schedule in Fig. 1. The samples were 18 mm in length and 12 mm in diameter, with the axis aligned in the rolling direction of the plate. The height-todiameter ratio of 1.5 was selected to ensure homogeneous deformation. Care was exercised to minimize friction between the test dies and the sample surface by machining flat-bottomed grooves on the end faces of samples. Graphite powders and thin pieces of mica sheet were used as lubricants in compression, resulting in fairly uniform deformation with negligibly small barreling. The uniaxial compression tests were performed on a servo-hydraulic 600 kN computerized Materials Testing System (MTS, Model 8500) equipped with a resistant furnace. During the test, the specimen temperature was measured and recorded using a Chromel–Alumel (Type K) thermocouple which was placed in contact with the surface of the compression specimen at the mid length position. Prior to deformation, the specimens were solutionized at 1150 °C for 5 min. The solution temperature was selected according to the solubility product of Ti and Nb precipitates [14,15]. At this temperature, Nb is completely dissolved and Ti is partially precipitated in the form of nitrides; very high temperatures, even in excess of melting point, being necessary for its total dissolution [16,17]. After solutionizing, the samples were cooled at a rate of 5 °C/s to
1150°C/ 5 min
deformation temperature of 845 °C, and held for 20 s to homogenize the temperature within the samples. Then samples were isothermally deformed with single pass strain values of 0.25, 0.4, and 0.8 at different constant strain rates of 0.001, 0.01, and 0.1 s 1. All specimens were water quenched in 2 s after deformation. Optical microscopy was conducted on mid-plane sections containing the axis of compression, in order to study the microstructural changes. 3. Results and discussion 3.1. True stress–true strain curves and related dynamic softening processes Fig. 2 shows the true stress–true strain curves for the specimens deformed up to a strain of 0.8 at 845 °C, i.e. just above Ar3, with different constant strain rates. As it can be seen in the figure, for the strain rate of 0.001 s 1, at first, the stress increases rapidly with increasing the strain. Then the stress further increases but it tends to be slow and exhibit a transition from work hardening to work hardening plus dynamic restoration. This indicates that the deformation energy is partly released in this stage due to dynamic softening occurrence. This behaviour might be interpreted as indicating dynamic recovery (DRC) of austenite. However, deformation at a higher temperature, i.e. 950 °C, exhibited the work hardening behaviour for this steel [18]. Therefore, DRC of austenite is ruled out. There are three mechanisms to release the deformation energy in deformed austenite during deformation: (i) dynamic recrystallization (DRX), (ii) DRC, or (iii) SIT of deformed austenite to ferrite [6,9]. By ruling out the first two
ε = 0.25, 0.4, 0.8 ε° = 0.001, 0.01, 0.1 s-1
210 0.1 s-1
Stress (Mpa)
180
Temperature
5°C/Sec
845°C Ar3 = 837°C
150 120 0.01 s-1
90
0.001 s-1
60 30
Ar1 = 751°C
0 0
0.2
0.4
0.6
0.8
1
Strain
Time Fig. 1. Schematic representation of the thermomechanical processing conditions.
Fig. 2. Flow curves of low carbon Nb–Ti microalloyed steel obtained under various strain rate at deformation temperature 847 °C with strain of 0.8.
B. Eghbali, A. Abdollah-Zadeh / Materials and Design 28 (2007) 1021–1026
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possibilities, thus, the softening behaviour observed at a strain rate of 0.001 s 1 can be attributed to the occurrence of SIT of austenite to ferrite. This is also in agreement with the observations of other workers studying hot deformation of plain carbon steels [6,11,13]. Yang and Wang [19] showed that when the deformation was given at lower temperature, DRX of austenite would not happen in plain low carbon steel during deformation, and the deformation energy can be fully utilized to strain-induced ferrite transformation. Then the driving force for phase transformation of austenite to ferrite is further improved. As shown in Fig. 2, increasing the strain rate from 0.001 to 0.1 s 1 increases the level of flow stress. In addition, at a higher strain rate, i.e. 0.1 s 1, there is no evidence of dynamic softening during deformation and the stress– strain curve indicates only the work hardening behaviour. It means that the amount of strain rate plays a very important role in occurrence of SIT during deformation by controlling the holding time and activation energy required for SIT, when the specimen is given same strain value. Increasing of strain rate leads to an increase in deformation resistance, and this is due to the fact that kinetic of dynamic softening events are diminished. With increasing the strain rate, the time for dynamic softening is diminished and the nucleation rate decreases so that the dynamic softening is inhibited. Increasing flow stress level would be expected to increase the dislocation density and decrease the subgrain size and hence increase the stored energy.
3.2. Evolution of strain-induced ferrite 3.2.1. Effects of strain Optical micrograph illustrating the evolution of straininduced ferrite grains with the amount of strain at deformation temperature of 847 °C, just above Ar3, and strain rate of 0.001 s 1 is presented in Fig. 3. Three specimens were compressed at different strains to study the effects of strain on the ferrite grain refinement. As shown in this figure, the deformed and elongated grains are present for the strain of 0.25. The fine ferrite grains had begun to nucleate inside a few of the austenite grains, and undeformed austenite grains have been statically transformed to coarse ferrite during cooling stage after deformation. This has resulted a bimodal distribution of fine ferrite grains due to dynamic transformation, and coarse ferrite grains due to static transformation, in the final microstructure. It has been demonstrated that the principal reason for refining the ferrite structure is the nucleation rate of the new phase, which is a result of increasing dislocation density in austenite just before the start of phase transformation [20]. With increasing the strain from 0.25 to 0.4, very fine ferrite grains are produced (Fig. 3). These grains have nucleated on the austenite grain boundaries, the vicinity of austenite grain boundaries and within the austenite grains. With increasing strain from 0.25 to 0.4, the extent of intra-
Fig. 3. Effect of strain on the evolution of ferrite grains refinement at deformation temperature of 845 °C with strain rate of 0.001 s 1.
granular ferrite nucleation increases. However, some ferrite grains in this microstructure are still visibly larger. The large strain produces more deformation defects such as deformation bands. These defects are beneficial in increasing the preferential nucleation sites of ferrite [21]. The austenite to ferrite transformation depends on the stored energy produced by deformation leading to the destabilization of austenite and acceleration of the austenite to ferrite transformation. The nucleation probability largely increases with increasing of strain and this leads to the very fine strain-induced ferrite grains. It can be seen in Fig. 3 that with increasing the strain to 0.8, the equiaxed and homogeneous ferrite grains of about
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4 lm with little amount of proeutectoid-deformed ferrite are obtained. The existence of little amount of proeutectoid-deformed ferrite in this microstructure is due to this fact that the mechanical working increases internal energy and thus free energy of system. This leads to the facts that Ar3 temperature is increased, and then the transformation of austenite to ferrite could happen at a higher temperature, as suggested by other workers [19,21]. It has been proposed that the austenite decomposition kinetics is influenced by its deformation state [9]: (i) the higher internal energy of the deformed and thus less stable austenite; (ii) the larger number of nucleation sites provided by defects. These two facts, rapid nucleation followed by early impingement and random orientation distribution, are believed to be the main reasons for the grain refinement by dynamic transformation. The straininduced ferrite is precipitated at the original grain boundaries of austenite or at the interface between the remaining austenite and proeutectoid ferrite, and at deformation bands and then has separated out all over the matrix. Once ferrite nuclei are formed, they will grow very rapidly because the diffusion rate also greatly increases due to high dislocation density, and this will make the ferrite grains to impinge on each other at the very early stage of grain growth. The results show that increasing the strain value has a positive effect on strain-induced ferrite formation. Strain increases diffusivity, number of nucleation site, and driving force for transformation which individually have the effects of increasing nucleation rate [6]. This is because of the increase of deformation energy as well as the increase of transformation time resulted from increase of strain value. As can be seen in Fig. 3, small strain, 0.25, cannot considerably produce strain-induced ferrite grains. However, the larger strain, 0.8, has produced higher volume fraction of strain-induced fine ferrite grains. It has been reported [22] that the critical strain value for initiation of strain-induced transformation is generally 0.3.
3.2.2. Effect of strain rate Evolution of strain-induced ferrite grains in the specimens deformed at 845 °C up to strain of 0.8 with various strain rates are shown in Fig. 4. It can be seen in this figure that the dominant microstructure is equiaxed ferrite grains with small amounts of deformed pro-eutectoid ferrite for strain rates of 0.001 and 0.01 s 1. The presence of somewhat elongated ferrite after deformation explains that phase transformation has been occurred during deformation. Increasing of strain rate from 0.001 to 0.01 s 1 reduces the ferrite grain size from about 4 to 3 lm, respectively, and decreases the amount of deformed pro-eutectoid ferrite. In other words, decreasing the deformation rate has lead to the formation of warm-worked pro-eutectoid ferrite grains. It should be noticed that the ferrite grains for these two strain rates have been produced during the SIT as discussed previously in Section 3.1. In other words, these microstructures have not been formed from other dynamic
Fig. 4. Effect of strain rate on the evolution of strain-induced ferrite formation at deformation temperature of 845 °C with strain of 0.8.
restoration processes such as DRC. The dynamic recovered structure mainly consists of pancaked grains of which elongation direction are perpendicular to the compression axis while the SIT structure has polygonal or equiaxed grains for the most parts [6,19]. This means that austenite to ferrite transformation for strain rates of 0.001 and 0.01 s 1 is occurred dynamically during deformation similar to DRC, but by SIT. As shown in Fig. 4, the microstructure of the specimen deformed at strain rate of 0.1 s 1 consists of ultrafine ferrite grains of about 2 lm. As can be seen from the corresponding flow curve in Fig. 2, there is not any evidence of dynamic softening during deformation for this strain
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rate. Therefore, the presence of very fine ferrite grains, 2 lm, and additionally, the absence of any elongated ferrite grains in the final microstructure could be an evidence of that the SIT at strain rate of 0.1 s 1 may be enhanced compared to that at lower strain rates as mechanical driving force, by more stored deformation energy, is expected to be higher at higher strain rates. It can be concluded that strain rate plays a very important role in SIT by controlling the holding time of SIT and activation energy required by SIT, when the specimen is given same deformation value. Low deformation temperature, near to Ar3, with high strain rate enables the austenite to store much more deformation energy in SIT processing than in conventional TMP. The high stored energy significantly increases the driving force for austenite to ferrite transformation and leads to the formation of very fine ferrite grains. As mentioned above, SIT is a diffusional transformation. Therefore, increasing the strain rate has a negative effect on SIT during deformation. On the other hand, decreasing the strain rate will increase the deformation time, and there would be enough time for the transformation to take place during deformation. If strain rate is very low, the strain-induced ferrite grains will have enough time to grow, and if the strain rate is very high, the strain-induced ferrite amount would be small. So, there would be an upper limit of strain rate for the SIT at a fixed temperature, and there would be an optimum strain rate range for the application of SIT in the ferrite refinement. Further investigation is needed to study the influence of precipitation of microalloying elements on SIT kinetics.
4. Conclusions The isothermal single pass hot compression tests were conducted in a low carbon Nb–Ti microalloyed steel at a constant temperature of 845 °C, just above Ar3, to study the effects of strain and strain rate on the evolution of strain-induced ferrite grain refinement through straininduced transformation. The following main conclusions can be drawn. (1) The occurrence of the SIT during deformation at 845 °C, just above Ar3, causes a dynamic softening behaviour for strain rates of 0.001 and 0.01 s 1, as reflected in the corresponding hot flow curves. (2) With increasing the strain from 0.25 to 0.8, at strain rate of 0.001 s 1, high volume fraction of straininduced ferrite with small grain size of 4 lm is produced due to an increase in deformation energy as well as an increase in transformation time. (3) Very fine ferrite grains of about 4 and 3 lm with little amount of proeutectoid-deformed ferrite are obtained by deformation at strain rates of 0.001 and 0.01 s 1, respectively, for the strain of 0.8. This has been attributed to SIT mechanism.
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(4) Deformation at 845 °C with the strain of 0.8 and strain rate of 0.1 s 1 produces the ultrafine, equiaxed and homogeneous ferrite grains of about 2 lm with little amount of proeutectoid-deformed ferrite.
Acknowledgements The authors express their gratitude for the financial support of this research by SASAD-Iran. Experiments on MTS thermomechanical simulator were performed at Maham Science and Technology Research Center, which is gratefully acknowledged. Professor H. Assadi is gratefully acknowledged for reviewing the manuscript. References [1] Sunghak L, Dongil K, Young KL, Ohjoon K. Transformation strengthening by thermomechanical treatments in C–Mn–Ni–Nb steels. Metall Mater Trans A 1995;26A:1093. [2] Hurley PJ, Hodgson PD, Muddle BC. Analysis and characterisation of ultra-fine ferrite produced during a new steel strip rolling process. Scr Mater 1999;40:433. [3] Priestner R, Al-Horr YM, Ibraheem AK. Effect of strain on formation of ultrafine ferrite in surface of hot rolled microalloyed steel. Mater Sci Technol 2002;18:973. [4] Hodgson PD, Hickson MR, Gibbs RK. The production and mechanical properties of ultrafine ferrite. Mater Sci Forum 1998;284–286:63. [5] Hickson MR, Hodgson PD. Effect of preroll quenching and post-roll quenching on production and properties of ultrafine ferrite in steel. Mater Sci Technol 1999;15:85. [6] Chio JK, Seo DH, Lee JS, Um KK, Choo WY. Formation of ultrafine ferrite by strain-induced dynamic transformation in plain low carbon steel. ISIJ Int 2003;43:746. [7] Eghbali B, Abdollah-zadeh A. The influence of thermomechanical parameters in ferrite grain refinement in a low carbon Nb-microalloyed steel. Scr Mater 2005;53:41. [8] Hodgson PD, Hickson MR, Gibbs PK. Ultrafine ferrite in low carbon steel. Scr Metall Mater 1999;40:1179. [9] Santos DB, Bruzszek RK, Rodrigues PCM, Pereloma EV. Formation of ultra-fine ferrite microstructure in warm rolled and annealed C– Mn steel. Mater Sci Eng A 2003;A346:189. [10] Huang CJ, Li DZ, Li YY. A finite element analysis of strain induced transformation rolling and an experimental study on the grain refinement potential of severe undercooling thermo-mechanical treatment. Mater Sci Eng A 2003;A352:136. [11] Huang YD, Froyen L. Important factors to obtain homogeneous and ultrafine ferrite–pearlite microstructure in low carbon steel. Mater Proc Technol 2002;124:216. [12] Kelly GL, Beladi H, Hodgson PD. Ultrafine grained ferrite formed by interrupted hot torsion deformation of plain carbon steel. ISIJ Int 2002;42:1585. [13] Hong SC, Lim SH, Lee KJ, Shin DH, Lee KS. Effect of undercooling of austenite on strain induced ferrite transformation behaviour. ISIJ Int 2003;43:394. [14] Ghosh A, Das S, Chatterjee S, Mishra B, Rao RR. Influence of thermo-mechanical processing and different post-cooling techniques on structure and properties of an ultra low carbon Cu bearing HSLA forging. Mater Sci Eng A 2003;A348:299. [15] Cho SH, Kang KB, Jonas JJ. The dynamic, static and metadynamic recrystallization of a nb-microalloyed steel. ISIJ Int 2001;41:63. [16] Manohar PA, Dunne DP, Chandra T, Killmore CR. Grain growth predictions in microalloyed steels. ISIJ Int 1996;36:194.
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