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Microstructural evolution at the initial stages of continuous annealing of cold rolled dual-phase steel
Materials Science and Engineering A 391 (2005) 296–304
Microstructural evolution at the initial stages of continuous annealing of cold rolled dual-phase steel R.O. Rochaa , T.M.F. Meloa , E.V. Perelomab , D.B. Santosc,∗ a Usiminas SA, Ipatinga, MG, Brazil School of Physics and Materials Engineering, Monash University, Vic. 3800, Australia Department of Metallurgical and Materials Engineering, Federal University of Minas Gerais, Belo Horizonte, MG, Brazil b
c
Received 26 March 2004; received in revised form 30 August 2004; accepted 30 August 2004
Abstract The performance of cold rolled dual-phase (DP) steels depends on their microstructure, which results from the thermomechanical processing conditions, involving hot rolling, cold rolling, and continuous annealing. The knowledge on the influence of each annealing stage on the microstructure formation is essential for manufacturing high-quality DP steels. In the present work, the effects of some intercritical annealing parameters (heating rate, soaking temperature, soaking time, and quench temperature) on the microstructure and mechanical properties of a cold rolled DP steel (0.08% C–1.91% Mn) were studied. The microstructure of specimens quenched after each annealing stage, simulated on a Gleeble, was analyzed using optical, scanning, and transmission electron microscopy. The tensile properties, determined for specimens submitted to complete annealing cycles, are influenced by the volume fractions of martensite, bainite, martensite/austenite (MA) constituent, and carbides, which depend on annealing processing parameters. The results obtained showed that the yield strength (YS) increase and the ultimate tensile strength (UTS) decrease with the increasing intercritical temperature. This can be explained by the increased formation of granular bainite associated with the increased volume fraction of austenite formed at the higher temperatures. The experimental data also showed that, for the annealing cycles carried out, UTS values in excess of 600 MPa could be obtained with the steel investigated. © 2004 Elsevier B.V. All rights reserved. Keywords: Continuous annealing; Dual-phase steel; Microstructural characterization; Mechanical properties
1. Introduction Fuel economy and, thereby, weight reduction are extremely important factors for the automotive industry [1–6]. The development of lightweight vehicles with high passive safety has been accomplished through the use of highstrength steels, like multiphase steels. Among these, the dualphase (DP) steels, whose microstructure consists mainly of ferrite and martensite, are an excellent choice for applications where low yield strength, high tensile strength, continuous yielding, and good uniform elongation are required [7–15]. Depending on the volume fraction of martensite in the microstructure, several grades of DP steels with different strength levels can be produced. DP steels containing 10–20% ∗
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of martensite typically have an ultimate tensile strength of around 600 MPa, yield strength from 300 to 400 MPa, and a relatively high ductility (24–30% of total elongation) [2,8]. The continuous annealing process to produce cold rolled DP steels typically have the following stages: heating to the intercritical temperature region, soaking in order to allow the nucleation and growth of austenite, slow cooling to the quench temperature, rapid cooling to transform the austenite into martensite, overaging, and air cooling. The amount and morphology of the constituents formed depend on the such annealing parameters [16–21]. The effects of the retained austenite, ferrite, and martensite morphologies on the mechanical behavior of DP steels have been intensively investigated [12–14,18,19]. On the other hand, some authors have investigated the influence of the intercritical annealing parameters on the microstructure and mechanical properties of these steels [16,17,22]. However,
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the microstructural evolution at the initial steps of the continuous annealing process is still not completely understood. For the simulation of the annealing cycle, some researchers have used dilatometer tests, with small specimen on which only microstructural and hardness evaluations can be performed [20,22]. However, simulations in a Gleeble can use larger samples, which also allow for subsequent tensile tests [23]. The purpose of the present research was to study the microstructural evolution of a cold rolled DP (0.08% C–1.91% Mn) steel at the initial stages (heating, soaking, and slow cooling) of an industrial annealing cycle simulated using a Gleeble. Selected samples were also subjected to the complete annealing cycle for the determination of the final tensile properties.
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Table 2 Processing parameters of industrial hot and cold rolling Thickness (mm)
Reduction per pass (%)
Hot rolling 29.68 17.01 10.62 6.98 5.14 3.94 3.23
42.7 37.6 34.3 26.3 23.2 18.2
Cold rolling 3.23 2.35 1.85 1.54 1.26 1.23
27.06 21.31 16.84 17.89 2.08
2. Experimental procedure An industrial heat was melted using a LD converter and cast into slabs on a two-strand continuous caster. The chemical composition obtained is given in Table 1. The slabs (252 mm of initial thickness) were heated to 1209 ◦ C and then hot rolled on two reversing four high roughing mill (29.67 mm) and a six-strand four high finishing train to the thickness of 3.23 mm. The finishing temperature was kept above Ar3 , 870 ◦ C, and after cooling the strip was coiled at 600 ◦ C. After pickling in a hydrochloric descaling line, the hot rolled coils were cold rolled on a five-strand, three four high, and two six-high, tandem mill to the final thickness of 1.23 mm. Table 2 shows the thickness and reduction per pass applied during the hot and cold rolling. After cold rolling, the 150 mm × 45 mm × 1.2 mm samples were machined and used for the simulation of the industrial continuous annealing (Fig. 1), utilizing a Gleeble [23]. Annealing was carried out within the intercritical temperature range, which was determined theoretically from the chemical composition of the material to be Ac3 = 854 ◦ C and Ac1 = 704 ◦ C [24]. The variables studied were the temperature and time of soaking and the quench temperature. The processing schedules and parameters used are given in Fig. 2 and Table 3. After each annealing stage, the samples were water quenched for microstructural and hardness examination. Transverse and longitudinal cold rolling plane sections of samples after rolling and annealing, were prepared following
Fig. 1. Schematic diagram of the continuous annealing process.
standard metallographic procedures and examined by optical, scanning electron microscopy (SEM) and transmission electron microscopy (TEM) [20,22]. For observation of the grain structure, the polished specimens were etched with 2% Nital. To reveal the presence of carbides, a 4% Picral etchant was used, while LePera etchant was applied to highlight the martensite/austenite (MA) constituent [20]. The volume fractions of the constituents were measured by manual counting
Table 1 Chemical composition of steel (wt.%) C Mn Si P S Al N
0.08 1.91 0.04 0.018 0.006 0.035 0.005
Fig. 2. Schedule of the annealing processing simulation. H, S, and SC are mean heating, soaking, and slow cooling steps, respectively.
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Table 3 Processing parameters of the annealing process simulation. H, S and SC mean heating, soaking, and slow cooling steps, respectively, according to Fig. 2 Code
Heating T
(◦ C)
Soaking
Slow cooling
Time (s)
Time (s)
T (◦ C)
Time (s)
H1 S1 SC1
820
155
– 56 56
– – 750
– – 23
H2 S2 SC2
780
207
– 75 75
– – 710
– – 31
H3 S3 SC3
750
310
– 113 113
– – 675
– – 46
TM
and by the IMAGE PRO-PLUS image analyzer software according to the ASTM standard E 562-02 [25]. Selected samples were analyzed using transmission electron microscopy. Thin foils of 3 mm diameter were punched out from about 0.2 mm thick slices, mechanically thinned to 0.12 mm and electro polished in a solution of methanol with 5% perchloric acid at −35 ◦ C in a Struers Tenupol double jet unit, set at 30 V. Thin foils were examined using a transmission electron microscope Philips CM 20 operated at 200 kV. Bright-field (BF), dark-field (DF), and selected-area electron diffraction (SAED) techniques were utilized. Vickers microhardness tests were conducted with a load of 2.94 N. Data from Vickers microhardness represent the mean of 20 impressions from the whole section analyzed. The tensile specimens (length: 135 mm, gauge length: 25 mm, and thickness: 1.23 mm) were machined from the Gleeble samples and three tests were carried out per condition at a crosshead speed of 5 mm/min. The results represent the mean of these three values. 3. Results and discussion The microstructure of the steel after hot rolling consists of banded polygonal ferrite and pearlite along the thickness
of the strip (Fig. 3(a)). The banding is mainly due to the high manganese content used. After cold rolling, the microstructure consists of elongated grains of ferrite and deformed colonies of pearlite (Fig. 3(b)). Some small islands of MA constituent, shown as white areas in Fig. 3(c), are also present in the microstructure. 3.1. Microstructure of annealed samples (Stages I–III) The microstructures of the samples after quenching from various stages of the processing schedule (Fig. 2 and Table 3) are given in Fig. 4. The microstructure consists of a mixture of ferrite, martensite, bainite, MA constituent, and Fe3 C carbides. Increasing the soaking temperature from 750 to 820 ◦ C resulted in lower amounts of ferrite (dark gray areas) in the microstructure (Fig. 4(a, c)). This is due to the larger amount of austenite formed at higher temperatures, which transforms into martensite, bainite, and carbides (light gray and white regions) on quenching. For the samples heated to lower soaking temperature and quenched after holding at this temperature (Fig. 4(f)) or after slow cooling (Fig. 4(i)), the amount of ferrite is higher and the amount of second constituents (martensite, MA, and carbides) is lower. The formation of austenite from ferrite is a diffusion-controlled phase transformation. Thus, the volume fraction of austenite and its coarseness increase with increasing soaking temperature and time. The microstructure is finer for samples heated to or annealed and quenched from 820 ◦ C, because the heating rate is higher and the time spent on heating and soaking is shorter (Table 3). These samples also contain higher amount of second constituents (martensite, cementite, and bainite) (Fig. 4(a, d)). The evolution of the ferrite recrystallization on heating to the soaking temperatures depends on the heating rate employed (Fig. 5), where experimental points were linked to show the tendency. Recrystallization was completed during heating to about 685, 710, and 730 ◦ C for samples treated according to routes H3, H2, and H1, respectively. This means that in all routes when the intercritical soaking stage is reached, the ferrite grains are fully recrystallized. This is in
Fig. 3. Microstructures of steel after hot rolling (a), and cold rolling (b, c). (a, b) Nital 2% etched, (c) LePera etched; small white regions are MA constituent. Ferrite—F; Pearlite—P.
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Fig. 4. SEM micrographs of samples after various processing routes designated in Table 2: (a) H1; (b) H2; (c) H3; (d) S1; (e) S2; (f) S3; (g) SC1; (h) SC2; and (i) SC3. Ferrite—F; Carbide—C; Bainite—B; Martensite–Austenite—M.
agreement with earlier reported data for different chemical compositions [16,17,19,20]. The distribution of MA constituent and carbides in the samples quenched after various stages of processing route 2 (heating to 780 ◦ C, Table 3) is shown in Fig. 6. Both MA islands, white (Fig. 6(a–c)); and carbide particles, dark (Fig. 6(d–f)) are primary located along the ferrite grain boundaries. This distribution suggests that the formation of austenite islands took place after the complete recrystalliza-
Fig. 5. Volume fraction of recrystallized ferrite as a function of continuous heating temperature for each annealing schedule.
tion of ferrite and the partial spheroidisation of cementite. The effect of the processing route on the volume fractions of phases present in the microstructure is shown in Figs. 7 and 8. After soaking (routes S1–S3), the microstructure becomes more homogeneous, the amount of austenite formed increases, and the amount of ferrite decreases with increasing soaking temperature. This is due to the diffusion-controlled mechanism of this phase transformation, where the higher temperature is more effective than the shorter heating and soaking time. As a consequence, a higher volume fraction of martensite is formed on quenching. With the progress of the annealing route (Stages II and III, Fig. 1), the volume fraction of Fe3 C carbides decreases, as more austenite is formed (Figs. 6(d–f) and 7). After slow cooling (Stage III, Fig. 1), the microstructure of all samples consists of ferrite with some carbides, small amount of granular bainite, and martensite formed at grain boundaries of the recrystallized ferrite grains. The volume fraction of MA constituent slightly decreased in all samples as compared to the corresponding samples after Stage II (Fig. 1), and was slightly higher for the lower annealing temperatures (Fig. 8). The total amount of second constituents (martensite + bainite + carbide + MA, Fig. 7) is lower in the samples processed according to the routes SC2 and SC3, despite their
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Fig. 6. Optical micrographs of samples quenched after various stages of annealing at 780 ◦ C: (a, d) H2; (b, e) S2; (c, f) SC2. (a–c) Picral etched, (d–f) LePera etched.
Fig. 7. Effect of the annealing cycle on the amount of second phases (martensite + bainite + carbide + MA) present.
Fig. 8. Effect of the annealing cycle on the volume fraction of MA constituent.
higher amounts of MA constituent (Fig. 8). This effect can be explained mainly by the changes in the volume fractions of martensite and bainite. During slow cooling after soaking at the higher temperatures, part of the austenite is transformed back to ferrite by diffusion-controlled phase transformation, while the remaining austenite with a lower carbon content transforms to granular bainite and martensite on subsequent quenching. The formation of ferrite on cooling affects the volume fraction of the second constituent formed (Fig. 8). The longer the slow cooling (lower slow cooling temperature), the more austenite transforms to ferrite and the lower is the amount of austenite available to form martensite on subsequent quenching. This was accordingly reflected in the volume fractions of constituents present in the microstructure (Figs. 7 and 8). Although the amount of carbides decreased substantially, there was probably not enough time for their complete dissolution. TEM studies have confirmed the SEM observations. The typical TEM micrographs of samples after various stages of annealing process are shown in Fig. 9. The microstructure of samples quenched on reaching the soaking temperature consisted of small crystals of martensite, retained austenite (RA), and/or MA constituent located at ferrite grain boundaries (Fig. 9(a)). The size and volume fraction of such constituents are the lowest at this stage of annealing and increase with each further stage (Fig. 9(a, d)). The martensite has predominantly lath morphology with high dislocation density, but some crystals contain a fine twinned structure (Fig. 9(f)). Fine spherical Fe3 C precipitates were observed in the ferrite matrix (Fig. 9(a, c)), and in some retained austenite crystals (Fig. 9(b, d)).
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The Mn content in the steel was high enough to stabilize the austenite. The TEM studies confirm that the carbide spheroidisation took place already during the reheating to the soaking temperature. Coarsening of carbide particles with each subsequent stage of annealing was also visible. In the microstructure of samples quenched after slow cooling (Stage III), granular bainite was observed. Granular bainite consists of equiaxed ferrite grains having higher dislocation density than polygonal ferrite and martensite/retained austenite islands (Fig. 9(e)). Granular bainite forms during continuous cooling at lower temperatures than polygonal ferrite or at higher cooling rates [26–28].
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3.2. Microhardness of annealed specimens (Stages I–III) Microhardness (HVN) value is smaller for samples quenched after heating to the soaking temperature (routes H1–H3) as compared to the samples quenched after soaking (routes S1–S3) (Fig. 10). Samples after slow cooling (routes SC1–SC3) exhibited intermediate values of microhardness. This is directly related to the microstructural changes occurred during the process. The lowest microhardness values for samples processed according to the routes H1–H3 are due to the smaller amount of second constituent in the microstructure, as less ferrite transformed to austenite on heating to the soaking tem-
Fig. 9. Thin foils TEM micrographs and SAED inserts of samples: (a) H1; (b, c) SC1; (d) S1; (e) S2; (f) S3. Zone axes are: (a) [1 1 1]␣ //[1 1 0]␥ ; (b) ¯ c ; (c) [1 1 1]␣ //[1 0 0]c ; (d) [1 1 0]//[2¯ 1 1] ¯ c ; (f) [1 3 5]␣ . ␣ denotes bcc lattice, ␥ fcc lattice, and c Fe3 C. [1 1 4]␥ //[1 0 1]
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Fig. 10. Effect of the annealing cycle on microhardness of samples processed according to the schedules shown in Fig. 2.
perature, than for samples held at the soaking temperature (S1–S3). However, the higher the soaking temperature, the higher is the volume fraction of second constituent and the harder is the material. The samples quenched immediately after soaking showed the highest hardness values, since more austenite was formed during soaking and there is no transformation back to ferrite, as in the case of samples processed according to routes SC1–SC3. However, the reason for the drop in the hardness value of sample S2 compared to S3 is unclear. To some extent it could be explained by the growth of austenite grains during soaking, which resulted in coarser martensite islands, decreasing its hardness. It is also worth to note that the samples subjected to the shorter slow cooling times (SC1 and SC2) showed higher hardness values than the SC3 samples, as there was less time for austenite to ferrite transformation, resulting in a higher volume fraction of martensite. 3.3. Microstructure and tensile properties of samples after full annealing cycle Comparing the microstructures after the full annealing cycle and after Stage III (SC2 condition), it can be seen that there is a small difference as related to ferrite grain size or distribution of martensite and MA constituents between these two samples (Figs. 4(h) and 11). However, larger amounts of granular bainite and carbides are present in the microstructure after the full annealing cycle. The formation and coarsening
of carbides take place during Stages III and V (overaging) for both martensite and ferrite. It is also possible to identify some islands of tempered martensite in the Fig. 11 (indicated by arrows). The tensile properties obtained after the complete cycle for each processing route are shown in Fig. 12. Sample 1 followed route S1; sample 2, route S2; and sample 3, route S3; then all samples were subjected to slow cooling to 710 ◦ C, followed by rapid cooling at 50 ◦ C s−1 , then overaging at 260 ◦ C for 300 s, and air cooling to room temperature. The yield strength (YS) increases as the soaking temperature increases, while the ultimate tensile strength (UTS) decreases. The total elongation does not change significantly, though it is slightly higher at the intermediate intercritical soaking temperature of 780 ◦ C. This behavior reflects the differences in the microstructure of the samples annealed at the three different temperatures. As discussed earlier, at the soaking stage a considerable amount of completely recrystallised ferrite transforms to austenite. Although some undissolved cementite was detected in the microstructure (Fig. 4a–c), the significant amount of cementite has dissolved in the austenite. During slow cooling some austenite transforms back to ferrite, but the higher the final slow cooling temperature, the higher is the final volume fraction of austenite. On accelerated cooling at 50 ◦ C s−1 , this remaining austenite will transform to granular bainite and martensite. The higher the austenite fraction, the larger is the amount of granular bainite and the lower is the amount of martensite, because of the lower carbon and manganese content of the austenite. These factors, mainly the carbon content, led to a lower Ms temperature [29]. Pichler et al. [20] also reported the formation of bainite in a dualphase steel of similar composition after soaking at 800 ◦ C for 60 s followed by cooling at 100 ◦ C s−1 . This correlates with the lowest UTS and the highest YS values obtained. It is well known that yield strength depends on lattice friction to dislocation motion, solid solution strengthening (interaction of dislocations with interstitial atoms of C and N and substitutional atoms of Mn and Si), precipitation hardening arisen from interaction of dislocations with carbides and grain boundary strengthening, which is defined by a free path for dislocation movement (ferrite grain size or martensite lath width). No significant differences in ferrite grain size and volume fraction of MA constituent were observed in the studied
Fig. 11. Microstructure of the sample annealed at 750 (a), 780 (b), and 820 ◦ C (c) for 113, 75 and 56 s, slow cooled to 710 ◦ C, rapid cooled at 50 ◦ C s−1 to 260 ◦ C, held for 300 s and then air cooled to room temperature. Tempering martensite—TM.
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samples. The main changes have occurred in volume fraction of martensite and bainite. Despite this decreasing hardness of the second phase, the effect of increasing volume fraction of hard phases prevailed and therefore with increasing the annealing temperatures a higher strength levels were attained. A decrease in the soaking temperature (and in the final slow cooling temperature) results in the formation of less austenite. Thus, less bainite is formed on quenching, because the austenite is richer in hardening elements, leading to an increase in the UTS and a decrease in the YS, which results in a lower yield ratio. The formation of carbides inside the martensite during the overaging (Stage V, Figs. 1 and 11) also contributes to a decrease in strength [20]. The ductility increases with an increase in the volume fraction of ferrite and a decrease in the amount of martensite. The modified law of mixtures [30,31] could be used to describe the flow behavior of dual-phase steel containing metastable retained austenite (usually when there is around 20% of martensite) and various amounts (up to 20%) of retained austenite in the microstructure of ferrite. Ferrite contributes to flow stress due to the: (i) normal work hardening of ferrite according to the Hollomon equation; (ii) hardening of ferrite in the areas close to hard martensite islands due to dislocations generated in ferrite due to the transformation of austenite to martensite on quenching; (iii) hardening of ferrite due to the dislocation generated due to the strain induced transformation of retained austenite to martensite during tensile testing. Austenite and martensite contribute to flow stress due to the normal work hardening. Contribution to strength due to the load transfer between the individual phases could be accounted for using an intermediate law of mixtures. In this work, the volume fractions of second constituent were around 19% as shown in Fig. 12. The volume fraction of MA increased from 2.5 to 3.5 and to 5.5% for annealing temperatures of 750, 780 and 820 ◦ C, respectively (Fig. 11a–c). A similar tendency was observed for bainite volume frac-
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tion (∼1% increase). These variations in volume fractions of phases present in the microstructure of studied samples could explain the slight differences in their mechanical properties. On the other hand, there was a refinement in ferrite grain size, which could also account for higher yield strength of samples processed at high annealing temperature. Whereas there is twice more of retained austenite (∼5%) in the sample annealed at 820 ◦ C compared to the samples annealed at two lower temperatures (∼2.5%), these amounts are still low to have any significant effect on the ductility (elongation) of samples after three conditions of heating and soaking, which show similar values for % elongation. In this case, the ductility in particular and mechanical properties in general are defined by the combination of all phases present in the microstructure of studied samples.
4. Conclusions The microstructural evolution in a DP steel subjected to various intercritical annealing routes has been investigated through simulation in a Gleeble. The results show that the microstructure after quenching from the first three stages of annealing cycle consists predominantly of ferrite with fine carbides, martensite, and small amounts of granular bainite with MA constituent. The volume fractions of phases present and the homogeneity of the microstructure depend on the processing parameters, such as heating rate, soaking temperature, and soaking time, and the duration and slow cooling final temperature process. During heating to the annealing temperature, the austenite islands form at the completely recrystallized ferrite grain boundaries and at spherical cementite particles. During soaking and slow cooling, the microstructure becomes more homogeneous while the volume fraction of ferrite and carbides decrease. The results obtained also show that the yield strength (YS) and ultimate tensile strength (UTS) change as a function of the intercritical temperature. The increase in YS and decrease in UTS with the decreasing annealing temperature can be explained by the lower content of hardening elements in the higher volume fraction of austenite, leading to the formation of increasing amounts of granular bainite. The UTS of all samples after the complete processing routes studied has significantly exceeded 600 MPa, typical for this class of steel.
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Fig. 12. Yield strength, ultimate tensile strength, and total elongation as a function of annealing temperature.
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Microstructural evolution at the initial stages of continuous annealing of cold rolled dual-phase steel R.O. Rochaa , T.M.F. Meloa , E.V. Perelomab , D.B. Santosc,∗ a Usiminas SA, Ipatinga, MG, Brazil School of Physics and Materials Engineering, Monash University, Vic. 3800, Australia Department of Metallurgical and Materials Engineering, Federal University of Minas Gerais, Belo Horizonte, MG, Brazil b
c
Received 26 March 2004; received in revised form 30 August 2004; accepted 30 August 2004
Abstract The performance of cold rolled dual-phase (DP) steels depends on their microstructure, which results from the thermomechanical processing conditions, involving hot rolling, cold rolling, and continuous annealing. The knowledge on the influence of each annealing stage on the microstructure formation is essential for manufacturing high-quality DP steels. In the present work, the effects of some intercritical annealing parameters (heating rate, soaking temperature, soaking time, and quench temperature) on the microstructure and mechanical properties of a cold rolled DP steel (0.08% C–1.91% Mn) were studied. The microstructure of specimens quenched after each annealing stage, simulated on a Gleeble, was analyzed using optical, scanning, and transmission electron microscopy. The tensile properties, determined for specimens submitted to complete annealing cycles, are influenced by the volume fractions of martensite, bainite, martensite/austenite (MA) constituent, and carbides, which depend on annealing processing parameters. The results obtained showed that the yield strength (YS) increase and the ultimate tensile strength (UTS) decrease with the increasing intercritical temperature. This can be explained by the increased formation of granular bainite associated with the increased volume fraction of austenite formed at the higher temperatures. The experimental data also showed that, for the annealing cycles carried out, UTS values in excess of 600 MPa could be obtained with the steel investigated. © 2004 Elsevier B.V. All rights reserved. Keywords: Continuous annealing; Dual-phase steel; Microstructural characterization; Mechanical properties
1. Introduction Fuel economy and, thereby, weight reduction are extremely important factors for the automotive industry [1–6]. The development of lightweight vehicles with high passive safety has been accomplished through the use of highstrength steels, like multiphase steels. Among these, the dualphase (DP) steels, whose microstructure consists mainly of ferrite and martensite, are an excellent choice for applications where low yield strength, high tensile strength, continuous yielding, and good uniform elongation are required [7–15]. Depending on the volume fraction of martensite in the microstructure, several grades of DP steels with different strength levels can be produced. DP steels containing 10–20% ∗
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of martensite typically have an ultimate tensile strength of around 600 MPa, yield strength from 300 to 400 MPa, and a relatively high ductility (24–30% of total elongation) [2,8]. The continuous annealing process to produce cold rolled DP steels typically have the following stages: heating to the intercritical temperature region, soaking in order to allow the nucleation and growth of austenite, slow cooling to the quench temperature, rapid cooling to transform the austenite into martensite, overaging, and air cooling. The amount and morphology of the constituents formed depend on the such annealing parameters [16–21]. The effects of the retained austenite, ferrite, and martensite morphologies on the mechanical behavior of DP steels have been intensively investigated [12–14,18,19]. On the other hand, some authors have investigated the influence of the intercritical annealing parameters on the microstructure and mechanical properties of these steels [16,17,22]. However,
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the microstructural evolution at the initial steps of the continuous annealing process is still not completely understood. For the simulation of the annealing cycle, some researchers have used dilatometer tests, with small specimen on which only microstructural and hardness evaluations can be performed [20,22]. However, simulations in a Gleeble can use larger samples, which also allow for subsequent tensile tests [23]. The purpose of the present research was to study the microstructural evolution of a cold rolled DP (0.08% C–1.91% Mn) steel at the initial stages (heating, soaking, and slow cooling) of an industrial annealing cycle simulated using a Gleeble. Selected samples were also subjected to the complete annealing cycle for the determination of the final tensile properties.
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Table 2 Processing parameters of industrial hot and cold rolling Thickness (mm)
Reduction per pass (%)
Hot rolling 29.68 17.01 10.62 6.98 5.14 3.94 3.23
42.7 37.6 34.3 26.3 23.2 18.2
Cold rolling 3.23 2.35 1.85 1.54 1.26 1.23
27.06 21.31 16.84 17.89 2.08
2. Experimental procedure An industrial heat was melted using a LD converter and cast into slabs on a two-strand continuous caster. The chemical composition obtained is given in Table 1. The slabs (252 mm of initial thickness) were heated to 1209 ◦ C and then hot rolled on two reversing four high roughing mill (29.67 mm) and a six-strand four high finishing train to the thickness of 3.23 mm. The finishing temperature was kept above Ar3 , 870 ◦ C, and after cooling the strip was coiled at 600 ◦ C. After pickling in a hydrochloric descaling line, the hot rolled coils were cold rolled on a five-strand, three four high, and two six-high, tandem mill to the final thickness of 1.23 mm. Table 2 shows the thickness and reduction per pass applied during the hot and cold rolling. After cold rolling, the 150 mm × 45 mm × 1.2 mm samples were machined and used for the simulation of the industrial continuous annealing (Fig. 1), utilizing a Gleeble [23]. Annealing was carried out within the intercritical temperature range, which was determined theoretically from the chemical composition of the material to be Ac3 = 854 ◦ C and Ac1 = 704 ◦ C [24]. The variables studied were the temperature and time of soaking and the quench temperature. The processing schedules and parameters used are given in Fig. 2 and Table 3. After each annealing stage, the samples were water quenched for microstructural and hardness examination. Transverse and longitudinal cold rolling plane sections of samples after rolling and annealing, were prepared following
Fig. 1. Schematic diagram of the continuous annealing process.
standard metallographic procedures and examined by optical, scanning electron microscopy (SEM) and transmission electron microscopy (TEM) [20,22]. For observation of the grain structure, the polished specimens were etched with 2% Nital. To reveal the presence of carbides, a 4% Picral etchant was used, while LePera etchant was applied to highlight the martensite/austenite (MA) constituent [20]. The volume fractions of the constituents were measured by manual counting
Table 1 Chemical composition of steel (wt.%) C Mn Si P S Al N
0.08 1.91 0.04 0.018 0.006 0.035 0.005
Fig. 2. Schedule of the annealing processing simulation. H, S, and SC are mean heating, soaking, and slow cooling steps, respectively.
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Table 3 Processing parameters of the annealing process simulation. H, S and SC mean heating, soaking, and slow cooling steps, respectively, according to Fig. 2 Code
Heating T
(◦ C)
Soaking
Slow cooling
Time (s)
Time (s)
T (◦ C)
Time (s)
H1 S1 SC1
820
155
– 56 56
– – 750
– – 23
H2 S2 SC2
780
207
– 75 75
– – 710
– – 31
H3 S3 SC3
750
310
– 113 113
– – 675
– – 46
TM
and by the IMAGE PRO-PLUS image analyzer software according to the ASTM standard E 562-02 [25]. Selected samples were analyzed using transmission electron microscopy. Thin foils of 3 mm diameter were punched out from about 0.2 mm thick slices, mechanically thinned to 0.12 mm and electro polished in a solution of methanol with 5% perchloric acid at −35 ◦ C in a Struers Tenupol double jet unit, set at 30 V. Thin foils were examined using a transmission electron microscope Philips CM 20 operated at 200 kV. Bright-field (BF), dark-field (DF), and selected-area electron diffraction (SAED) techniques were utilized. Vickers microhardness tests were conducted with a load of 2.94 N. Data from Vickers microhardness represent the mean of 20 impressions from the whole section analyzed. The tensile specimens (length: 135 mm, gauge length: 25 mm, and thickness: 1.23 mm) were machined from the Gleeble samples and three tests were carried out per condition at a crosshead speed of 5 mm/min. The results represent the mean of these three values. 3. Results and discussion The microstructure of the steel after hot rolling consists of banded polygonal ferrite and pearlite along the thickness
of the strip (Fig. 3(a)). The banding is mainly due to the high manganese content used. After cold rolling, the microstructure consists of elongated grains of ferrite and deformed colonies of pearlite (Fig. 3(b)). Some small islands of MA constituent, shown as white areas in Fig. 3(c), are also present in the microstructure. 3.1. Microstructure of annealed samples (Stages I–III) The microstructures of the samples after quenching from various stages of the processing schedule (Fig. 2 and Table 3) are given in Fig. 4. The microstructure consists of a mixture of ferrite, martensite, bainite, MA constituent, and Fe3 C carbides. Increasing the soaking temperature from 750 to 820 ◦ C resulted in lower amounts of ferrite (dark gray areas) in the microstructure (Fig. 4(a, c)). This is due to the larger amount of austenite formed at higher temperatures, which transforms into martensite, bainite, and carbides (light gray and white regions) on quenching. For the samples heated to lower soaking temperature and quenched after holding at this temperature (Fig. 4(f)) or after slow cooling (Fig. 4(i)), the amount of ferrite is higher and the amount of second constituents (martensite, MA, and carbides) is lower. The formation of austenite from ferrite is a diffusion-controlled phase transformation. Thus, the volume fraction of austenite and its coarseness increase with increasing soaking temperature and time. The microstructure is finer for samples heated to or annealed and quenched from 820 ◦ C, because the heating rate is higher and the time spent on heating and soaking is shorter (Table 3). These samples also contain higher amount of second constituents (martensite, cementite, and bainite) (Fig. 4(a, d)). The evolution of the ferrite recrystallization on heating to the soaking temperatures depends on the heating rate employed (Fig. 5), where experimental points were linked to show the tendency. Recrystallization was completed during heating to about 685, 710, and 730 ◦ C for samples treated according to routes H3, H2, and H1, respectively. This means that in all routes when the intercritical soaking stage is reached, the ferrite grains are fully recrystallized. This is in
Fig. 3. Microstructures of steel after hot rolling (a), and cold rolling (b, c). (a, b) Nital 2% etched, (c) LePera etched; small white regions are MA constituent. Ferrite—F; Pearlite—P.
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Fig. 4. SEM micrographs of samples after various processing routes designated in Table 2: (a) H1; (b) H2; (c) H3; (d) S1; (e) S2; (f) S3; (g) SC1; (h) SC2; and (i) SC3. Ferrite—F; Carbide—C; Bainite—B; Martensite–Austenite—M.
agreement with earlier reported data for different chemical compositions [16,17,19,20]. The distribution of MA constituent and carbides in the samples quenched after various stages of processing route 2 (heating to 780 ◦ C, Table 3) is shown in Fig. 6. Both MA islands, white (Fig. 6(a–c)); and carbide particles, dark (Fig. 6(d–f)) are primary located along the ferrite grain boundaries. This distribution suggests that the formation of austenite islands took place after the complete recrystalliza-
Fig. 5. Volume fraction of recrystallized ferrite as a function of continuous heating temperature for each annealing schedule.
tion of ferrite and the partial spheroidisation of cementite. The effect of the processing route on the volume fractions of phases present in the microstructure is shown in Figs. 7 and 8. After soaking (routes S1–S3), the microstructure becomes more homogeneous, the amount of austenite formed increases, and the amount of ferrite decreases with increasing soaking temperature. This is due to the diffusion-controlled mechanism of this phase transformation, where the higher temperature is more effective than the shorter heating and soaking time. As a consequence, a higher volume fraction of martensite is formed on quenching. With the progress of the annealing route (Stages II and III, Fig. 1), the volume fraction of Fe3 C carbides decreases, as more austenite is formed (Figs. 6(d–f) and 7). After slow cooling (Stage III, Fig. 1), the microstructure of all samples consists of ferrite with some carbides, small amount of granular bainite, and martensite formed at grain boundaries of the recrystallized ferrite grains. The volume fraction of MA constituent slightly decreased in all samples as compared to the corresponding samples after Stage II (Fig. 1), and was slightly higher for the lower annealing temperatures (Fig. 8). The total amount of second constituents (martensite + bainite + carbide + MA, Fig. 7) is lower in the samples processed according to the routes SC2 and SC3, despite their
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Fig. 6. Optical micrographs of samples quenched after various stages of annealing at 780 ◦ C: (a, d) H2; (b, e) S2; (c, f) SC2. (a–c) Picral etched, (d–f) LePera etched.
Fig. 7. Effect of the annealing cycle on the amount of second phases (martensite + bainite + carbide + MA) present.
Fig. 8. Effect of the annealing cycle on the volume fraction of MA constituent.
higher amounts of MA constituent (Fig. 8). This effect can be explained mainly by the changes in the volume fractions of martensite and bainite. During slow cooling after soaking at the higher temperatures, part of the austenite is transformed back to ferrite by diffusion-controlled phase transformation, while the remaining austenite with a lower carbon content transforms to granular bainite and martensite on subsequent quenching. The formation of ferrite on cooling affects the volume fraction of the second constituent formed (Fig. 8). The longer the slow cooling (lower slow cooling temperature), the more austenite transforms to ferrite and the lower is the amount of austenite available to form martensite on subsequent quenching. This was accordingly reflected in the volume fractions of constituents present in the microstructure (Figs. 7 and 8). Although the amount of carbides decreased substantially, there was probably not enough time for their complete dissolution. TEM studies have confirmed the SEM observations. The typical TEM micrographs of samples after various stages of annealing process are shown in Fig. 9. The microstructure of samples quenched on reaching the soaking temperature consisted of small crystals of martensite, retained austenite (RA), and/or MA constituent located at ferrite grain boundaries (Fig. 9(a)). The size and volume fraction of such constituents are the lowest at this stage of annealing and increase with each further stage (Fig. 9(a, d)). The martensite has predominantly lath morphology with high dislocation density, but some crystals contain a fine twinned structure (Fig. 9(f)). Fine spherical Fe3 C precipitates were observed in the ferrite matrix (Fig. 9(a, c)), and in some retained austenite crystals (Fig. 9(b, d)).
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The Mn content in the steel was high enough to stabilize the austenite. The TEM studies confirm that the carbide spheroidisation took place already during the reheating to the soaking temperature. Coarsening of carbide particles with each subsequent stage of annealing was also visible. In the microstructure of samples quenched after slow cooling (Stage III), granular bainite was observed. Granular bainite consists of equiaxed ferrite grains having higher dislocation density than polygonal ferrite and martensite/retained austenite islands (Fig. 9(e)). Granular bainite forms during continuous cooling at lower temperatures than polygonal ferrite or at higher cooling rates [26–28].
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3.2. Microhardness of annealed specimens (Stages I–III) Microhardness (HVN) value is smaller for samples quenched after heating to the soaking temperature (routes H1–H3) as compared to the samples quenched after soaking (routes S1–S3) (Fig. 10). Samples after slow cooling (routes SC1–SC3) exhibited intermediate values of microhardness. This is directly related to the microstructural changes occurred during the process. The lowest microhardness values for samples processed according to the routes H1–H3 are due to the smaller amount of second constituent in the microstructure, as less ferrite transformed to austenite on heating to the soaking tem-
Fig. 9. Thin foils TEM micrographs and SAED inserts of samples: (a) H1; (b, c) SC1; (d) S1; (e) S2; (f) S3. Zone axes are: (a) [1 1 1]␣ //[1 1 0]␥ ; (b) ¯ c ; (c) [1 1 1]␣ //[1 0 0]c ; (d) [1 1 0]//[2¯ 1 1] ¯ c ; (f) [1 3 5]␣ . ␣ denotes bcc lattice, ␥ fcc lattice, and c Fe3 C. [1 1 4]␥ //[1 0 1]
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Fig. 10. Effect of the annealing cycle on microhardness of samples processed according to the schedules shown in Fig. 2.
perature, than for samples held at the soaking temperature (S1–S3). However, the higher the soaking temperature, the higher is the volume fraction of second constituent and the harder is the material. The samples quenched immediately after soaking showed the highest hardness values, since more austenite was formed during soaking and there is no transformation back to ferrite, as in the case of samples processed according to routes SC1–SC3. However, the reason for the drop in the hardness value of sample S2 compared to S3 is unclear. To some extent it could be explained by the growth of austenite grains during soaking, which resulted in coarser martensite islands, decreasing its hardness. It is also worth to note that the samples subjected to the shorter slow cooling times (SC1 and SC2) showed higher hardness values than the SC3 samples, as there was less time for austenite to ferrite transformation, resulting in a higher volume fraction of martensite. 3.3. Microstructure and tensile properties of samples after full annealing cycle Comparing the microstructures after the full annealing cycle and after Stage III (SC2 condition), it can be seen that there is a small difference as related to ferrite grain size or distribution of martensite and MA constituents between these two samples (Figs. 4(h) and 11). However, larger amounts of granular bainite and carbides are present in the microstructure after the full annealing cycle. The formation and coarsening
of carbides take place during Stages III and V (overaging) for both martensite and ferrite. It is also possible to identify some islands of tempered martensite in the Fig. 11 (indicated by arrows). The tensile properties obtained after the complete cycle for each processing route are shown in Fig. 12. Sample 1 followed route S1; sample 2, route S2; and sample 3, route S3; then all samples were subjected to slow cooling to 710 ◦ C, followed by rapid cooling at 50 ◦ C s−1 , then overaging at 260 ◦ C for 300 s, and air cooling to room temperature. The yield strength (YS) increases as the soaking temperature increases, while the ultimate tensile strength (UTS) decreases. The total elongation does not change significantly, though it is slightly higher at the intermediate intercritical soaking temperature of 780 ◦ C. This behavior reflects the differences in the microstructure of the samples annealed at the three different temperatures. As discussed earlier, at the soaking stage a considerable amount of completely recrystallised ferrite transforms to austenite. Although some undissolved cementite was detected in the microstructure (Fig. 4a–c), the significant amount of cementite has dissolved in the austenite. During slow cooling some austenite transforms back to ferrite, but the higher the final slow cooling temperature, the higher is the final volume fraction of austenite. On accelerated cooling at 50 ◦ C s−1 , this remaining austenite will transform to granular bainite and martensite. The higher the austenite fraction, the larger is the amount of granular bainite and the lower is the amount of martensite, because of the lower carbon and manganese content of the austenite. These factors, mainly the carbon content, led to a lower Ms temperature [29]. Pichler et al. [20] also reported the formation of bainite in a dualphase steel of similar composition after soaking at 800 ◦ C for 60 s followed by cooling at 100 ◦ C s−1 . This correlates with the lowest UTS and the highest YS values obtained. It is well known that yield strength depends on lattice friction to dislocation motion, solid solution strengthening (interaction of dislocations with interstitial atoms of C and N and substitutional atoms of Mn and Si), precipitation hardening arisen from interaction of dislocations with carbides and grain boundary strengthening, which is defined by a free path for dislocation movement (ferrite grain size or martensite lath width). No significant differences in ferrite grain size and volume fraction of MA constituent were observed in the studied
Fig. 11. Microstructure of the sample annealed at 750 (a), 780 (b), and 820 ◦ C (c) for 113, 75 and 56 s, slow cooled to 710 ◦ C, rapid cooled at 50 ◦ C s−1 to 260 ◦ C, held for 300 s and then air cooled to room temperature. Tempering martensite—TM.
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samples. The main changes have occurred in volume fraction of martensite and bainite. Despite this decreasing hardness of the second phase, the effect of increasing volume fraction of hard phases prevailed and therefore with increasing the annealing temperatures a higher strength levels were attained. A decrease in the soaking temperature (and in the final slow cooling temperature) results in the formation of less austenite. Thus, less bainite is formed on quenching, because the austenite is richer in hardening elements, leading to an increase in the UTS and a decrease in the YS, which results in a lower yield ratio. The formation of carbides inside the martensite during the overaging (Stage V, Figs. 1 and 11) also contributes to a decrease in strength [20]. The ductility increases with an increase in the volume fraction of ferrite and a decrease in the amount of martensite. The modified law of mixtures [30,31] could be used to describe the flow behavior of dual-phase steel containing metastable retained austenite (usually when there is around 20% of martensite) and various amounts (up to 20%) of retained austenite in the microstructure of ferrite. Ferrite contributes to flow stress due to the: (i) normal work hardening of ferrite according to the Hollomon equation; (ii) hardening of ferrite in the areas close to hard martensite islands due to dislocations generated in ferrite due to the transformation of austenite to martensite on quenching; (iii) hardening of ferrite due to the dislocation generated due to the strain induced transformation of retained austenite to martensite during tensile testing. Austenite and martensite contribute to flow stress due to the normal work hardening. Contribution to strength due to the load transfer between the individual phases could be accounted for using an intermediate law of mixtures. In this work, the volume fractions of second constituent were around 19% as shown in Fig. 12. The volume fraction of MA increased from 2.5 to 3.5 and to 5.5% for annealing temperatures of 750, 780 and 820 ◦ C, respectively (Fig. 11a–c). A similar tendency was observed for bainite volume frac-
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tion (∼1% increase). These variations in volume fractions of phases present in the microstructure of studied samples could explain the slight differences in their mechanical properties. On the other hand, there was a refinement in ferrite grain size, which could also account for higher yield strength of samples processed at high annealing temperature. Whereas there is twice more of retained austenite (∼5%) in the sample annealed at 820 ◦ C compared to the samples annealed at two lower temperatures (∼2.5%), these amounts are still low to have any significant effect on the ductility (elongation) of samples after three conditions of heating and soaking, which show similar values for % elongation. In this case, the ductility in particular and mechanical properties in general are defined by the combination of all phases present in the microstructure of studied samples.
4. Conclusions The microstructural evolution in a DP steel subjected to various intercritical annealing routes has been investigated through simulation in a Gleeble. The results show that the microstructure after quenching from the first three stages of annealing cycle consists predominantly of ferrite with fine carbides, martensite, and small amounts of granular bainite with MA constituent. The volume fractions of phases present and the homogeneity of the microstructure depend on the processing parameters, such as heating rate, soaking temperature, and soaking time, and the duration and slow cooling final temperature process. During heating to the annealing temperature, the austenite islands form at the completely recrystallized ferrite grain boundaries and at spherical cementite particles. During soaking and slow cooling, the microstructure becomes more homogeneous while the volume fraction of ferrite and carbides decrease. The results obtained also show that the yield strength (YS) and ultimate tensile strength (UTS) change as a function of the intercritical temperature. The increase in YS and decrease in UTS with the decreasing annealing temperature can be explained by the lower content of hardening elements in the higher volume fraction of austenite, leading to the formation of increasing amounts of granular bainite. The UTS of all samples after the complete processing routes studied has significantly exceeded 600 MPa, typical for this class of steel.
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Fig. 12. Yield strength, ultimate tensile strength, and total elongation as a function of annealing temperature.
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