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Effect of carbon content on the recrystallization kinetics of Nb-steels
Scripta Materialia 56 (2007) 1059–1062 www.elsevier.com/locate/scriptamat
Effect of carbon content on the recrystallization kinetics of Nb-steels H. Beladi* and P.D. Hodgson Centre for Material and Fibre Innovation, Deakin University, Geelong, Vic. 3217, Australia Received 16 January 2007; revised 10 February 2007; accepted 21 February 2007 Available online 21 March 2007
A rapid method was used to study the effect of carbon content on the kinetics of post-deformation softening, t50, in Nb-steels. The hot deformation behaviour of austenite was not affected by carbon. However, the t50 was influenced by the carbon with different effects in different temperature regimes. At deformation temperatures above the non-recrystallization temperature, Tnr, carbon produced a small change in the softening behaviour. However, the t50 was significantly retarded with increasing carbon content at deformation temperatures lower than Tnr, due to Nb(C,N) precipitates. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Recrystallization kinetics; Nb-Steels; Carbon effect
Understanding the high-temperature deformation behaviour of steel is necessary to control the process and dimensional accuracy of hot-rolled products as well as the final properties of these materials. A number of studies have examined the effect of thermomechanical conditions, such as temperature, strain and strain rate, on the hot deformation behaviour of steels. However, there has been little investigation of the effect of steel composition. Although carbon has received wider attention, there is still disagreement regarding the effect of carbon on the hot deformation of austenite and the recrystallization kinetics of steel. Early work by Mead and Birchenall [1] showed that carbon increases the self-diffusion coefficient of iron. These authors also reported that the activation energy for iron self-diffusion decreases with an increase in carbon content. It would, therefore, be expected that higher-carbon steels would have a higher recovery rate. This finding was supported by other workers, who showed a reduction in the work-hardening rate and flow stress with increasing carbon content [2,3]. In contrast, other reports have showed an increase [4,5] or little effect [6,7]. In this work, very large changes in carbon content were used and most of the focus was on the hot deformation and dynamic recrystallization behaviour. The limited work to date on post-deformation recrystalliza-
* Corresponding author. Tel.: +61 3 5227 3321; fax: +61 3 5227 1103; e-mail: [email protected]
tion [7] has shown a very strong effect at low temperatures for large changes in carbon content, but little effect at high temperatures. To date, research has mainly focused on the effect of carbon on hot deformation of plain carbon steels, and no results have been reported for microalloyed steels. In microalloyed steels, carbon interacts with microalloying elements (e.g., Nb), affecting the hot deformation behaviour. The aim of the current study was to investigate the effect of carbon on the kinetics of recrystallization for Nb-microalloyed steels. In this case, the carbon levels have been restricted to those encountered in weldable microalloyed steels. Three Nb-microalloyed steels were prepared in a vacuum furnace using electrolytically pure iron with alloying additions. The alloying additions were kept constant apart from the carbon, which varied from 0.04% to 0.16%C (Table 1). A hot torsion deformation simulator was employed to deform the samples. The specimens had a gauge length of 20 mm and a gauge diameter of 6.7 mm. The entire length of the specimen was enclosed in a quartz tube, in which a positive pressure of argon gas was maintained to prevent oxidation of the specimen during testing. The maximum temperature difference between the two ends of the gauge length was 10 °C. The reheating temperature was chosen based on the Nb(C,N) solution temperature (Ts), which was estimated by a modified Irvine et al. [8] equation to take into account the Mn and Si levels [9] (Table 2). Therefore,
1359-6462/$ - see front matter Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2007.02.029
H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062
Table 1. Chemical composition of steels (wt.%) Steel
%C
%Mn
%Si
%Nb
%N
A B C
0.042 0.105 0.16
0.83 0.86 0.86
0.28 0.29 0.29
0.031 0.037 0.034
0.0039 0.0041 0.0033
a
dc (lm)
Ar3 (°C)
Tnr (°C)
Ts (°C)
A B C
41 33 30
826 793 786
932 970 987
1051 1181 1225
Ar3
808
844
882
150
778 °C
920 958
100 1020
50
Table 2. The temperature and deformation characteristics of steels Steel
200 Tnr
Stress (MPa)
1060
0 0.0
0.5
1.0
1.5
2.0
Strain
b 180
Ar3 =826 °C
the specimen was reheated at 1250 °C for 5 min, which is well above Ts for all steels. The specimen was then cooled to 1100 °C and held for 80 s, followed by two pass strains of 0.4 with an interpass time of 20 s at a strain rate of 1 s1. The deformed structure was fully recrystallized within the 20 s interpass time after each pass strain. This led to a homogeneous austenite grain distribution in the roughing condition. Afterwards, three different deformation schedules were carried out to determine different thermomechanical processing characteristics. A multipass torsion test was performed during cooling (1 °C s1) from 1100 °C to determine the characteristic temperatures for all alloys, the empirical austenite to ferrite transformation (Ar3) and the non-recrystallization temperature (Tnr). This follows the method developed by Bai et al. [10]. A pass strain of 0.15 and a strain rate of 1 s1 were applied. Fourteen passes were performed and the temperature decreased from pass to pass (20 °C). The Tnr and Ar3 temperatures for all alloys were estimated as discussed later (Table 2). The second hot torsion test was performed after the roughing process to establish the continuous flow curve for full dynamic recrystallization (DRX) under different thermomechanical conditions. This was used to identify the characteristic strains at a given thermomechanical condition: the strain to the initiation of DRX (eC) and peak stress (eP). Finally, interrupted torsion tests were carried out to study the kinetics of post-deformation softening. After the roughing condition, the specimen was cooled at 1 °C s1 to a given deformation temperature and held for 80 s before a double-hit deformation was applied. Each experiment involved deforming to the required strain followed by holding for a given time and then re-deforming to a strain of 0.2. The deformation temperature was varied between 900 and 1100 °C. The specimen was isothermally deformed at strain rates of 1 1 100, at and 0.1 s1. The softening fraction, X s ¼ rr33 r r2 a given thermomechanical condition was then estimated. r1 is the yield stress of the first deformation, r2 is the yield stress of the second deformation and r3 is the stress at the end of first deformation [11]. The multipass torsion test showed that the flow stress increased from pass to pass with decreasing temperature (Fig. 1a). There was, however, little interdeformation softening at a temperature below 940 °C. The mean flow stress (MFS) as a function of inverse temperature
MFS (MPa)
160 140 120
Tnr=932 °C
100 80 60 7.5
Region I 8.0
Region III
Region II 8.5
9.0 -4
9.5
10.0
-1
1/T X 10 (K )
Figure 1. (a) Stress–strain curves corresponding to multipass torsion tests for steel A (the temperatures of the passes are shown). (b) Dependence of the MFS on the inverse absolute temperature during a multipass torsion test of steel A.
showed two slope changes, which allowed the graph to be divided into three regions (Fig. 1b). In region I, which corresponds to high-temperature deformation, full recrystallization took place and there was no passto-pass strain accumulation. The increase in MFS was solely due to the decrease in temperature. In region II, which corresponds to deformation below the Tnr, there was only partial recrystallization or no recrystallization at all. Here, the strain accumulated from pass to pass, so that the MFS increased more rapidly with decreasing temperature. The MFS dropped significantly with temperature in region III due to the austenite to ferrite transformation (Ar3). The Ar3 temperature decreased with an increase in carbon, whereas the Tnr temperature increased with carbon (Table 2) as reported elsewhere [12]. In some ways the shift in Tnr was stronger than expected as this continuous cooling test often means the first indications are due to solute drag, which would be largely determined by the Nb content. This is somewhat reinforced by Figure 1a, which shows that for the low-carbon steel there is more softening after the 920 °C pass than in the interval before. Below 900 °C there is clearly no recrystallization between passes. Figure 2 represents some flow curves at different thermomechanical conditions for all alloys. The curves exhibit work hardening to a peak stress followed by softening to a steady state, typical for the occurrence of DRX. The characteristic strains, eC and eP, were identified along the continuous stress–strain curves. eC was estimated using the approach developed by Poliak and Jonas [13]. In this method, eC is associated with an inflection point the strain-hardening rate as a function of flow stress plot.
H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062 10 900°C, 1s
A B C B
-1
Stress (MPa)
140 120 100
A 1100 °C, 1s
80 40 0 0.0
900°C, 0.1s
-1
A
60 20
-1
1100°C, 0.1s
0.5
-1
B C A C B
1.0
1.5
Time to 50% softening (s)
180 160
1061
*
ε
strain independent softening
0.1 0.1
1
Strain
Figure 3. Time to 50% softening as a function of strain at different alloys at a strain rate of 1 s1 and a deformation temperature of 1050 °C.
Time to 50% softening (s)
a
ε•=1/s
Steel A Steel B Steel C
100
1050°C
1000°C 900°C
10 1100°C
1
0.1
1 Strain
b Time to 50% softening (s)
1000
100
Steel A Steel B Steel C 1000°C
•
ε =0.1/s
900°C
900°C
1050°C 1100 °C
10
1 0.1
1
Strain
Figure 4. Time for 50% softening for different alloys under different thermomechanical conditions.
91.4 s
100 Time to 50% Softening (s)
The peak stresses and strains increased with a decrease in temperature and increasing strain rate for all alloys. Carbon appears to not have any specific effect on the stress–strain behaviour (Fig. 2). This is not consistent with the early reports that showed a reduction in work-hardening rate and flow stress with carbon [1–3]. However, Collinson et al. [7] also reported very little effect of carbon on the stress–strain behaviour in plain carbon steels for Zener–Hollomon (Z) conditions similar to the current study. They have shown that the carbon effect only appears in steels with >0.4 carbon and under high Z conditions. As noted below, one interesting change in behaviour was for the 0.1C steel at 0.1 s1 and 900 °C. Here there was no clear peak in the stress–strain curve. A rapid method of determining post-deformation softening kinetics through the study of the time to 50% softening (t50) was used for the current analysis (the details of which are described elsewhere [14]). The graph obviously consisted of two regions: strain-dependent softening and strain-independent softening that changed at the transition strain (e*). t50 depends strongly on thermomechanical parameters as well as steel composition. The softening kinetics were accelerated by increasing temperature in all alloys (Fig. 4), although the strain-independent softening is significantly less affected by temperature than straindependent softening (Fig. 4). However, the strain rate showed a strong influence over strain-independent softening and less effect on strain-dependent softening. These trends are consistent with the results reported in the literature [14,15]. The effect of carbon on the kinetics of post-deformation softening differed depending upon the temperature. This effect could be divided into two regions: above and below Tnr. The kinetics of post-deformation softening were very similar for all alloys above the Tnr for both the straindependent and strain-independent regions (Fig. 4). However, the kinetics of softening were significantly slower than that of C–Mn steel of similar composition but without Nb (Fig. 3). This can be explained by the solute drag effect of Nb, which retards the mobility of grain boundaries (i.e., recrystallization) [10]. The expected time for 50% post-deformation softening at a temperature of 900 °C (below Tnr) was estimated by averaging the t50 of all alloys at a given thermomechanical condition above Tnr ( Fig. 5) and extrapolating to the lower temperature. For the low-carbon alloy
1
strain dependent softening
2.0
Strain
Figure 2. Stress–strain curves of steels at different thermomechanical conditions. A, B and C represent steel A, steel B and steel C, respectively.
Steel A Steel B Steel C [15] C-Mn Steel [Ref. ]
ε
47.4 s
=0.1
29.8 s
.2
ε =0 .3 ε =0
10
1 7.2
7.4
7.6
7.8 8.0 8.2 4 -1 10 /T (K )
8.4
8.6
Figure 5. Time to 50% softening as a function of inverse temperature for different strains using the average of t50 for all steels at deformation temperatures above Tnr at a strain rate of 1 s1. The values represent the estimated t50 at 900 °C and a strain rate of 1 s1.
(steel A), the values were close to those expected by extrapolation for strains greater than 0.7 at a strain rate of 1 s1. For steel A at low strains and steels B and C the
1062
H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062 Strain rate:1/s Strain rate: 0.1/s
C B A B A
ε*
1
7.0
7.2
7.4
7.6
7.8 8.0 8.2 -1 10000/T (K )
8.4
8.6
Figure 6. e* as a function of inverse temperature using the average of e* for all steels at deformation temperatures above Tnr for different strain rates. A, B, and C represent e* at 900 °C for steel A, steel B and steel C, respectively.
t50 results were much longer than expected. This can only be due to strain-induced precipitation as the solute drag effects are similar for all three steels. It is interesting to note that the effect in the strain-independent region is not significant at this strain rate and the t50 values are similar to those expected for all steels. There is a strong effect on e*, though, which can also be predicted in the absence of strain-induced precipitation through extrapolation of the high-temperature data (Fig. 6). Again the values for steel A are close to that expected, while those of the other steels are much higher; especially considering this is a logarithmic scale. At the lower strain rate of 0.1 s1 the results are more complex. The main difference is in the strain-independent region where there is now a marked effect of Nb. Even steel A shows a marked retardation in e* and the t50 value for the strain-independent region. It is expected from extrapolation of the high-temperature data that t50 would be in the range of 8–10 s, compared with the observed 20–30 s. Increasing the carbon to 0.1 (steel B) increased this by over an order of magnitude. A dramatic change in the post-deformation softening with carbon can be explained by the formation of Nb(C,N) at a deformation temperature lower than Tnr. Nb(C,N) precipitates have a much stronger effect on grain boundary mobility through pinning than Nb will have in solution through solute drag. Carbon significantly increases the driving force for Nb(C,N) formation at a given thermomechanical condition. Therefore, both the rate and the amount of Nb(C,N) precipitates increase with carbon, resulting in a much lower static recrystallization rate. This effect is more pronounced when Nb(C,N) precipitates form during deformation, as these retard DRX and significantly change the initial state of the austenite. This can explain the significant difference in post-deformation behaviour of steels A and B at strain rate of 0.1 s1 at 900 °C. At this condition steel A showed evidence of DRX, while little or no DRX occurred in steel B (Figs. 2 and 4b). Dutta and Sellars developed a method for strain-induced precipitation as a function of composition and process conditions [16]. This model was slightly modified by one of the current authors [17] based on torsion tests of a range of steels. It is interesting to note that using these two models gives precipitate times for T = 900 °C, e = 0.8 and strain rate of 1 s1 of 50–100 s, 3–6 s and
1–3 s for carbon contents of 0.04, 0.1 and 0.16, respectively. As recrystallization takes around 10 s for 50% by extrapolation of high-temperature data, it suggests the two higher-carbon steels will precipitate before they recrystallize, whereas steel A will be recrystallized before precipitation, which supports the current results. For steel B at 0.1 s1 the precipitation times are similar to the deformation times, which explains why all results are increased. However, for steel A the precipitation times are much longer (120–250 s for e = 1) than either the deformation or the recrystallization times. Therefore, the models do not completely explain the current results. Continuous and interrupted hot torsion tests were performed to study the influence of carbon content on the hot deformation and recrystallization behaviour of austenite in Nb-steels. Carbon appeared to have no specific influence on hot working of austenite in the temperature range 900–1100 °C. However, the postdeformation softening behaviour of austenite behaved differently in respect to the carbon content in different temperature regimes. At deformation temperatures above Tnr, the post-deformation softening was slightly influenced by carbon. However, the post-deformation softening was significantly retarded by carbon at deformation temperatures lower than Tnr due to Nb(C,N) precipitates. If these precipitates formed during deformation, then the recrystallization rate was markedly reduced, even at very high strains. This research was supported by grants through the Australian Research Council (H.B.) including an ARC Federation Fellowship (P.D.H.). The technical assistance of Mr. R. Pow is gratefully acknowledged. [1] H.W. Mead, C.E. Birchenall, Trans. AIME 206 (1956) 1336. [2] P.J. Wray, Met. Trans. A 13 (1982) 125. [3] T. Sakai, M. Ohashi, Tetsu-to-Hangane 678 (1981) 134. [4] Y. Misaka, T. Yoshimoto, JSTP 8 (1967) 414. [5] S. Shida, JSTP 10 (1969) 610. [6] M.J. Stewart, in: J.B. Balance (Ed.), The Hot Deformation of Austenite, TMS-AIME, 1977, p. 47. [7] D.C. Collinson, P.D. Hodgson, C.H.J. Davies, in: Thermec’97, Wollongong, Australia, 1997, p. 483. [8] K.J. Irvine, F.B. Pickering, T. Gladman, J. Iron Steel Inst. 205 (1967) 161. [9] F. Siciliano Jr., J.J. Jonas, Metall. Mater. Trans. A 31 (2000) 511. [10] D.Q. Bai, S. Yue, W.P. Sun, J.J. Jonas, Metall. Mater. Trans. A 11 (1980) 387. [11] R.K. Gibbs, P.D. Hodgson, B.A. Parker, E. Morris, in: P.K. Liaw, J.R. Weertman, H.L. Marcus, J.S. Santner (Eds.), Fine Symposium, TMS, Warrendale, PA, 1991, p. 73. [12] F. Borrato, R. Barbosa, S. Yue, J.J. Jonas, in: Thermec’88, Tokyo, Japan, 1988, p. 383. [13] E.I. Poliak, J.J. Jonas, ISIJ Int. 43 (2003) 684. [14] P.D. Hodgson, S.H. Zahiri, J.J. Whale, ISIJ Int. 44 (2004) 1224. [15] M.R. Cartmill, M.R. Barnett, S.H. Zahiri, P.D. Hodgson, ISIJ Int. 45 (2005) 1903. [16] B. Dutta, C.M. Sellars, Mater. Sci. Technol. 3 (1987) 197. [17] P.D. Hodgson, R.K. Gibbs, ISIJ Int. 32 (1992) 1329.
Effect of carbon content on the recrystallization kinetics of Nb-steels H. Beladi* and P.D. Hodgson Centre for Material and Fibre Innovation, Deakin University, Geelong, Vic. 3217, Australia Received 16 January 2007; revised 10 February 2007; accepted 21 February 2007 Available online 21 March 2007
A rapid method was used to study the effect of carbon content on the kinetics of post-deformation softening, t50, in Nb-steels. The hot deformation behaviour of austenite was not affected by carbon. However, the t50 was influenced by the carbon with different effects in different temperature regimes. At deformation temperatures above the non-recrystallization temperature, Tnr, carbon produced a small change in the softening behaviour. However, the t50 was significantly retarded with increasing carbon content at deformation temperatures lower than Tnr, due to Nb(C,N) precipitates. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Recrystallization kinetics; Nb-Steels; Carbon effect
Understanding the high-temperature deformation behaviour of steel is necessary to control the process and dimensional accuracy of hot-rolled products as well as the final properties of these materials. A number of studies have examined the effect of thermomechanical conditions, such as temperature, strain and strain rate, on the hot deformation behaviour of steels. However, there has been little investigation of the effect of steel composition. Although carbon has received wider attention, there is still disagreement regarding the effect of carbon on the hot deformation of austenite and the recrystallization kinetics of steel. Early work by Mead and Birchenall [1] showed that carbon increases the self-diffusion coefficient of iron. These authors also reported that the activation energy for iron self-diffusion decreases with an increase in carbon content. It would, therefore, be expected that higher-carbon steels would have a higher recovery rate. This finding was supported by other workers, who showed a reduction in the work-hardening rate and flow stress with increasing carbon content [2,3]. In contrast, other reports have showed an increase [4,5] or little effect [6,7]. In this work, very large changes in carbon content were used and most of the focus was on the hot deformation and dynamic recrystallization behaviour. The limited work to date on post-deformation recrystalliza-
* Corresponding author. Tel.: +61 3 5227 3321; fax: +61 3 5227 1103; e-mail: [email protected]
tion [7] has shown a very strong effect at low temperatures for large changes in carbon content, but little effect at high temperatures. To date, research has mainly focused on the effect of carbon on hot deformation of plain carbon steels, and no results have been reported for microalloyed steels. In microalloyed steels, carbon interacts with microalloying elements (e.g., Nb), affecting the hot deformation behaviour. The aim of the current study was to investigate the effect of carbon on the kinetics of recrystallization for Nb-microalloyed steels. In this case, the carbon levels have been restricted to those encountered in weldable microalloyed steels. Three Nb-microalloyed steels were prepared in a vacuum furnace using electrolytically pure iron with alloying additions. The alloying additions were kept constant apart from the carbon, which varied from 0.04% to 0.16%C (Table 1). A hot torsion deformation simulator was employed to deform the samples. The specimens had a gauge length of 20 mm and a gauge diameter of 6.7 mm. The entire length of the specimen was enclosed in a quartz tube, in which a positive pressure of argon gas was maintained to prevent oxidation of the specimen during testing. The maximum temperature difference between the two ends of the gauge length was 10 °C. The reheating temperature was chosen based on the Nb(C,N) solution temperature (Ts), which was estimated by a modified Irvine et al. [8] equation to take into account the Mn and Si levels [9] (Table 2). Therefore,
1359-6462/$ - see front matter Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2007.02.029
H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062
Table 1. Chemical composition of steels (wt.%) Steel
%C
%Mn
%Si
%Nb
%N
A B C
0.042 0.105 0.16
0.83 0.86 0.86
0.28 0.29 0.29
0.031 0.037 0.034
0.0039 0.0041 0.0033
a
dc (lm)
Ar3 (°C)
Tnr (°C)
Ts (°C)
A B C
41 33 30
826 793 786
932 970 987
1051 1181 1225
Ar3
808
844
882
150
778 °C
920 958
100 1020
50
Table 2. The temperature and deformation characteristics of steels Steel
200 Tnr
Stress (MPa)
1060
0 0.0
0.5
1.0
1.5
2.0
Strain
b 180
Ar3 =826 °C
the specimen was reheated at 1250 °C for 5 min, which is well above Ts for all steels. The specimen was then cooled to 1100 °C and held for 80 s, followed by two pass strains of 0.4 with an interpass time of 20 s at a strain rate of 1 s1. The deformed structure was fully recrystallized within the 20 s interpass time after each pass strain. This led to a homogeneous austenite grain distribution in the roughing condition. Afterwards, three different deformation schedules were carried out to determine different thermomechanical processing characteristics. A multipass torsion test was performed during cooling (1 °C s1) from 1100 °C to determine the characteristic temperatures for all alloys, the empirical austenite to ferrite transformation (Ar3) and the non-recrystallization temperature (Tnr). This follows the method developed by Bai et al. [10]. A pass strain of 0.15 and a strain rate of 1 s1 were applied. Fourteen passes were performed and the temperature decreased from pass to pass (20 °C). The Tnr and Ar3 temperatures for all alloys were estimated as discussed later (Table 2). The second hot torsion test was performed after the roughing process to establish the continuous flow curve for full dynamic recrystallization (DRX) under different thermomechanical conditions. This was used to identify the characteristic strains at a given thermomechanical condition: the strain to the initiation of DRX (eC) and peak stress (eP). Finally, interrupted torsion tests were carried out to study the kinetics of post-deformation softening. After the roughing condition, the specimen was cooled at 1 °C s1 to a given deformation temperature and held for 80 s before a double-hit deformation was applied. Each experiment involved deforming to the required strain followed by holding for a given time and then re-deforming to a strain of 0.2. The deformation temperature was varied between 900 and 1100 °C. The specimen was isothermally deformed at strain rates of 1 1 100, at and 0.1 s1. The softening fraction, X s ¼ rr33 r r2 a given thermomechanical condition was then estimated. r1 is the yield stress of the first deformation, r2 is the yield stress of the second deformation and r3 is the stress at the end of first deformation [11]. The multipass torsion test showed that the flow stress increased from pass to pass with decreasing temperature (Fig. 1a). There was, however, little interdeformation softening at a temperature below 940 °C. The mean flow stress (MFS) as a function of inverse temperature
MFS (MPa)
160 140 120
Tnr=932 °C
100 80 60 7.5
Region I 8.0
Region III
Region II 8.5
9.0 -4
9.5
10.0
-1
1/T X 10 (K )
Figure 1. (a) Stress–strain curves corresponding to multipass torsion tests for steel A (the temperatures of the passes are shown). (b) Dependence of the MFS on the inverse absolute temperature during a multipass torsion test of steel A.
showed two slope changes, which allowed the graph to be divided into three regions (Fig. 1b). In region I, which corresponds to high-temperature deformation, full recrystallization took place and there was no passto-pass strain accumulation. The increase in MFS was solely due to the decrease in temperature. In region II, which corresponds to deformation below the Tnr, there was only partial recrystallization or no recrystallization at all. Here, the strain accumulated from pass to pass, so that the MFS increased more rapidly with decreasing temperature. The MFS dropped significantly with temperature in region III due to the austenite to ferrite transformation (Ar3). The Ar3 temperature decreased with an increase in carbon, whereas the Tnr temperature increased with carbon (Table 2) as reported elsewhere [12]. In some ways the shift in Tnr was stronger than expected as this continuous cooling test often means the first indications are due to solute drag, which would be largely determined by the Nb content. This is somewhat reinforced by Figure 1a, which shows that for the low-carbon steel there is more softening after the 920 °C pass than in the interval before. Below 900 °C there is clearly no recrystallization between passes. Figure 2 represents some flow curves at different thermomechanical conditions for all alloys. The curves exhibit work hardening to a peak stress followed by softening to a steady state, typical for the occurrence of DRX. The characteristic strains, eC and eP, were identified along the continuous stress–strain curves. eC was estimated using the approach developed by Poliak and Jonas [13]. In this method, eC is associated with an inflection point the strain-hardening rate as a function of flow stress plot.
H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062 10 900°C, 1s
A B C B
-1
Stress (MPa)
140 120 100
A 1100 °C, 1s
80 40 0 0.0
900°C, 0.1s
-1
A
60 20
-1
1100°C, 0.1s
0.5
-1
B C A C B
1.0
1.5
Time to 50% softening (s)
180 160
1061
*
ε
strain independent softening
0.1 0.1
1
Strain
Figure 3. Time to 50% softening as a function of strain at different alloys at a strain rate of 1 s1 and a deformation temperature of 1050 °C.
Time to 50% softening (s)
a
ε•=1/s
Steel A Steel B Steel C
100
1050°C
1000°C 900°C
10 1100°C
1
0.1
1 Strain
b Time to 50% softening (s)
1000
100
Steel A Steel B Steel C 1000°C
•
ε =0.1/s
900°C
900°C
1050°C 1100 °C
10
1 0.1
1
Strain
Figure 4. Time for 50% softening for different alloys under different thermomechanical conditions.
91.4 s
100 Time to 50% Softening (s)
The peak stresses and strains increased with a decrease in temperature and increasing strain rate for all alloys. Carbon appears to not have any specific effect on the stress–strain behaviour (Fig. 2). This is not consistent with the early reports that showed a reduction in work-hardening rate and flow stress with carbon [1–3]. However, Collinson et al. [7] also reported very little effect of carbon on the stress–strain behaviour in plain carbon steels for Zener–Hollomon (Z) conditions similar to the current study. They have shown that the carbon effect only appears in steels with >0.4 carbon and under high Z conditions. As noted below, one interesting change in behaviour was for the 0.1C steel at 0.1 s1 and 900 °C. Here there was no clear peak in the stress–strain curve. A rapid method of determining post-deformation softening kinetics through the study of the time to 50% softening (t50) was used for the current analysis (the details of which are described elsewhere [14]). The graph obviously consisted of two regions: strain-dependent softening and strain-independent softening that changed at the transition strain (e*). t50 depends strongly on thermomechanical parameters as well as steel composition. The softening kinetics were accelerated by increasing temperature in all alloys (Fig. 4), although the strain-independent softening is significantly less affected by temperature than straindependent softening (Fig. 4). However, the strain rate showed a strong influence over strain-independent softening and less effect on strain-dependent softening. These trends are consistent with the results reported in the literature [14,15]. The effect of carbon on the kinetics of post-deformation softening differed depending upon the temperature. This effect could be divided into two regions: above and below Tnr. The kinetics of post-deformation softening were very similar for all alloys above the Tnr for both the straindependent and strain-independent regions (Fig. 4). However, the kinetics of softening were significantly slower than that of C–Mn steel of similar composition but without Nb (Fig. 3). This can be explained by the solute drag effect of Nb, which retards the mobility of grain boundaries (i.e., recrystallization) [10]. The expected time for 50% post-deformation softening at a temperature of 900 °C (below Tnr) was estimated by averaging the t50 of all alloys at a given thermomechanical condition above Tnr ( Fig. 5) and extrapolating to the lower temperature. For the low-carbon alloy
1
strain dependent softening
2.0
Strain
Figure 2. Stress–strain curves of steels at different thermomechanical conditions. A, B and C represent steel A, steel B and steel C, respectively.
Steel A Steel B Steel C [15] C-Mn Steel [Ref. ]
ε
47.4 s
=0.1
29.8 s
.2
ε =0 .3 ε =0
10
1 7.2
7.4
7.6
7.8 8.0 8.2 4 -1 10 /T (K )
8.4
8.6
Figure 5. Time to 50% softening as a function of inverse temperature for different strains using the average of t50 for all steels at deformation temperatures above Tnr at a strain rate of 1 s1. The values represent the estimated t50 at 900 °C and a strain rate of 1 s1.
(steel A), the values were close to those expected by extrapolation for strains greater than 0.7 at a strain rate of 1 s1. For steel A at low strains and steels B and C the
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H. Beladi, P. D. Hodgson / Scripta Materialia 56 (2007) 1059–1062 Strain rate:1/s Strain rate: 0.1/s
C B A B A
ε*
1
7.0
7.2
7.4
7.6
7.8 8.0 8.2 -1 10000/T (K )
8.4
8.6
Figure 6. e* as a function of inverse temperature using the average of e* for all steels at deformation temperatures above Tnr for different strain rates. A, B, and C represent e* at 900 °C for steel A, steel B and steel C, respectively.
t50 results were much longer than expected. This can only be due to strain-induced precipitation as the solute drag effects are similar for all three steels. It is interesting to note that the effect in the strain-independent region is not significant at this strain rate and the t50 values are similar to those expected for all steels. There is a strong effect on e*, though, which can also be predicted in the absence of strain-induced precipitation through extrapolation of the high-temperature data (Fig. 6). Again the values for steel A are close to that expected, while those of the other steels are much higher; especially considering this is a logarithmic scale. At the lower strain rate of 0.1 s1 the results are more complex. The main difference is in the strain-independent region where there is now a marked effect of Nb. Even steel A shows a marked retardation in e* and the t50 value for the strain-independent region. It is expected from extrapolation of the high-temperature data that t50 would be in the range of 8–10 s, compared with the observed 20–30 s. Increasing the carbon to 0.1 (steel B) increased this by over an order of magnitude. A dramatic change in the post-deformation softening with carbon can be explained by the formation of Nb(C,N) at a deformation temperature lower than Tnr. Nb(C,N) precipitates have a much stronger effect on grain boundary mobility through pinning than Nb will have in solution through solute drag. Carbon significantly increases the driving force for Nb(C,N) formation at a given thermomechanical condition. Therefore, both the rate and the amount of Nb(C,N) precipitates increase with carbon, resulting in a much lower static recrystallization rate. This effect is more pronounced when Nb(C,N) precipitates form during deformation, as these retard DRX and significantly change the initial state of the austenite. This can explain the significant difference in post-deformation behaviour of steels A and B at strain rate of 0.1 s1 at 900 °C. At this condition steel A showed evidence of DRX, while little or no DRX occurred in steel B (Figs. 2 and 4b). Dutta and Sellars developed a method for strain-induced precipitation as a function of composition and process conditions [16]. This model was slightly modified by one of the current authors [17] based on torsion tests of a range of steels. It is interesting to note that using these two models gives precipitate times for T = 900 °C, e = 0.8 and strain rate of 1 s1 of 50–100 s, 3–6 s and
1–3 s for carbon contents of 0.04, 0.1 and 0.16, respectively. As recrystallization takes around 10 s for 50% by extrapolation of high-temperature data, it suggests the two higher-carbon steels will precipitate before they recrystallize, whereas steel A will be recrystallized before precipitation, which supports the current results. For steel B at 0.1 s1 the precipitation times are similar to the deformation times, which explains why all results are increased. However, for steel A the precipitation times are much longer (120–250 s for e = 1) than either the deformation or the recrystallization times. Therefore, the models do not completely explain the current results. Continuous and interrupted hot torsion tests were performed to study the influence of carbon content on the hot deformation and recrystallization behaviour of austenite in Nb-steels. Carbon appeared to have no specific influence on hot working of austenite in the temperature range 900–1100 °C. However, the postdeformation softening behaviour of austenite behaved differently in respect to the carbon content in different temperature regimes. At deformation temperatures above Tnr, the post-deformation softening was slightly influenced by carbon. However, the post-deformation softening was significantly retarded by carbon at deformation temperatures lower than Tnr due to Nb(C,N) precipitates. If these precipitates formed during deformation, then the recrystallization rate was markedly reduced, even at very high strains. This research was supported by grants through the Australian Research Council (H.B.) including an ARC Federation Fellowship (P.D.H.). The technical assistance of Mr. R. Pow is gratefully acknowledged. [1] H.W. Mead, C.E. Birchenall, Trans. AIME 206 (1956) 1336. [2] P.J. Wray, Met. Trans. A 13 (1982) 125. [3] T. Sakai, M. Ohashi, Tetsu-to-Hangane 678 (1981) 134. [4] Y. Misaka, T. Yoshimoto, JSTP 8 (1967) 414. [5] S. Shida, JSTP 10 (1969) 610. [6] M.J. Stewart, in: J.B. Balance (Ed.), The Hot Deformation of Austenite, TMS-AIME, 1977, p. 47. [7] D.C. Collinson, P.D. Hodgson, C.H.J. Davies, in: Thermec’97, Wollongong, Australia, 1997, p. 483. [8] K.J. Irvine, F.B. Pickering, T. Gladman, J. Iron Steel Inst. 205 (1967) 161. [9] F. Siciliano Jr., J.J. Jonas, Metall. Mater. Trans. A 31 (2000) 511. [10] D.Q. Bai, S. Yue, W.P. Sun, J.J. Jonas, Metall. Mater. Trans. A 11 (1980) 387. [11] R.K. Gibbs, P.D. Hodgson, B.A. Parker, E. Morris, in: P.K. Liaw, J.R. Weertman, H.L. Marcus, J.S. Santner (Eds.), Fine Symposium, TMS, Warrendale, PA, 1991, p. 73. [12] F. Borrato, R. Barbosa, S. Yue, J.J. Jonas, in: Thermec’88, Tokyo, Japan, 1988, p. 383. [13] E.I. Poliak, J.J. Jonas, ISIJ Int. 43 (2003) 684. [14] P.D. Hodgson, S.H. Zahiri, J.J. Whale, ISIJ Int. 44 (2004) 1224. [15] M.R. Cartmill, M.R. Barnett, S.H. Zahiri, P.D. Hodgson, ISIJ Int. 45 (2005) 1903. [16] B. Dutta, C.M. Sellars, Mater. Sci. Technol. 3 (1987) 197. [17] P.D. Hodgson, R.K. Gibbs, ISIJ Int. 32 (1992) 1329.