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Microstructures and mechanical properties of hot-rolled Nb-microalloyed TRIP steels by different thermo-mechanical processes
Materials Science & Engineering A 605 (2014) 14–21
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructures and mechanical properties of hot-rolled Nb-microalloyed TRIP steels by different thermo-mechanical processes Qingxiao Feng a, Longfei Li a,n, Wangyue Yang b, Zuqing Sun a a b
State Key Laboratory for Advanced Metals and Materials, University of Science & Technology Beijing, Beijing 100083, China School of Materials Science and Engineering, University of Science & Technology Beijing, Beijing 100083, China
art ic l e i nf o
a b s t r a c t
Article history: Received 20 November 2013 Received in revised form 28 February 2014 Accepted 10 March 2014 Available online 19 March 2014
The microstructural evolution and mechanical properties of hot-rolled Nb-microalloyed TRIP steel with different states of austenite (recrystallized/pancaked) prior to the transformation of austenite to ferrite were investigated under the thermo-mechanical process based on controlled cooling or based on dynamic transformation of undercooled austenite (DTUA) by hot compression tests using a Gleeble-1500 hot simulator, in combination with OM, SEM, TEM, XRD and ICP–OES. The results indicated that the influence of the states of austenite prior to the transformation of austenite to ferrite on the multi-phase microstructures and the mechanical properties of the tested steel became weak in the case of the process based on DTUA, in comparison with the case of the process based on controlled cooling. The much refined and uniform microstructures with smaller sizes of ferrite and bainitic packets and the more stable retained austenite were obtained by the process based on DTUA, resulting in the markedly improved combination of strength and ductility of the tested steel. & 2014 Elsevier B.V. All rights reserved.
Keywords: Hot-rolled TRIP steel Nb-microalloyed Dynamic transformation of undercooled austenite Retained austenite Mechanical properties
1. Introduction There are two principal ways for producing TRIP steels, continuous annealing of cold-rolled material and hot strip rolling. Most previous studies have concentrated on the production of cold-rolled TRIP steels and made a lot of achievements [1–3], but the production of hot-rolled TRIP steels is relatively simple and cost-effective [4,5]. One key step for obtaining the multiphase microstructure of TRIP steels, which consists of polygonal ferrite matrix, bainite, retained austenite and some martensite, is to obtain the mixed microstructure consisting of ferrite and austenite with nearly same volume fraction of about 50%, before the isothermal bainitic treatment [6]. For conventional hot-rolled process of TRIP steels, the control of such mixed microstructure is commonly developed by adjusting cooling rate and time through multi-stage cooling after hot rolling of austenite [7,8], that is relatively difficult to be realized precisely in the industry, limiting the mass production of hot-rolled TRIP steel. Recently, a novel thermo-mechanical process to manufacture hotrolled TRIP steel was developed in our previous studies based on dynamic transformation of undercooled austenite (DTUA) [9–12]. The typical feature of this process is that the formation of the mixed microstructure of ferrite and austenite can be well controlled by the
n
Corresponding author. Tel.: þ 86 10 62334862; fax: þ 86 10 62333447. E-mail address: [email protected] (L. Li).
http://dx.doi.org/10.1016/j.msea.2014.03.051 0921-5093/& 2014 Elsevier B.V. All rights reserved.
applied strain of hot deformation of undercooled austenite. In comparison with controlling the cooling rate of undercooled austenite, the strain is relatively easy to be controlled accurately in the industry, and the multi-phase microstructure of hot-rolled TRIP steel could be obtained easily and be refined markedly by the thermo-mechanical process based on DTUA [9]. Hot-rolled TRIP steels obtained by such process demonstrate a good combination of strength and ductility, because of the refined and uniform multi-phase microstructure. For example, the tensile strength and the total elongation of C–Mn–Al–Si TRIP steel are 780 MPa and 34% respectively [10], which is nearly same as the mechanical properties of cold-rolled TRIP steel with the similar composition. With the addition of microalloying element Nb in C–Mn– Al–Si TRIP steel [12], the tensile strength can be enhanced to 845 MPa without any deterioration in the total elongation, due to the further refinement of the multi-phase microstructure. The product of tensile strength and total elongation is up to 28.7 GPa%, which is close to the requirement of the mechanical properties for the “Third-Generation” advanced high-strength steel, i.e. 30 GPa% [13]. Commonly, the main aim of the addition of Nb in steels is to adjust the state of austenite prior to the transformation of austenite to ferrite (γ-α) and thus accelerate the formation of ferrite grains, resulting in the structural refinement in the final products. For example, in the case of conventional hot-rolled 0.2C– 1.55Mn–1.55Si TRIP steel [14], it is found that the deformation in the nonrecrystallized austenite region can lead to the refinement of ferrite grains, and affect the transformation kinetics and morphology of bainite as well as the distribution of retained
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austenite in the multi-phase microstructure. As a result of that, TRIP steels with different states of austenite prior to the transformation of γ-α exhibited different mechanical properties. Moreover, it is found that in the case of conventional hot-rolled Nbmicroalloyed C–Mn–Si TRIP steel [15], the deformation in the nonrecrystallized austenite region can promote the precipitation of Nb-containing particles and result in a decrease in the volume fraction of retained austenite due to a reduction of Nb in the solid solution and, therefore, a decrease in parent austenite stability. In the previous work [9–12], the state of austenite prior to the dynamic transformation of undercooled austenite is only equiaxed recrystallized austenite grains for C–Mn–Al–Si hot-rolled TRIP steels with and without Nb. The key objective of this work, therefore, is to study the effect of the different states of austenite prior to the transformation of γ-α on the microstructural evolution and the mechanical properties of hot-rolled Nb-microalloyed TRIP steel based on DTUA processing, in comparison with that based on controlled-cooling processing.
2. Materials and experimental method The used material is a C–Mn–Al–Si–Nb steel by vacuum melting, containing 0.20C, 1.50Mn, 0.50Si, 1.04Al, 0.038Nb and the balance Fe (wt%). The casting blanks were first reheated to 1200 1C and held for 2 h, then hot-forged at 1150–900 1C followed by air-cooling to room temperature, and then machined into different specimens for hot-simulation test. The hot-simulation tests were performed using a Gleeble-1500 hot simulator, with the thermo-mechanical processes shown in Fig. 1. After austenitization at 1250 1C for 5 min, the samples were subjected to a 3-stage deformation at the austenitic recrystallization region from 1080 1C to 1000 1C; in each stage the applied strain was 30% at 2 s 1 and then held for 10 s to obtain recrystallized austenite (R) through static recrystallization. Then the recrystallized austenite was cooled to austenitic nonrecrystallization region, i.e. 900 1C, deformed to strain of 30% at 2 s 1 to obtain pancaked austenite (P). Next, the recrystallized austenite or pancaked austenite underwent rest of the thermo-mechanical process based on controlled cooling (Fig. 1a) or based on DTUA (Fig. 1b). After the isothermal bainitic treatment at 450 1C for 5 min, the samples were quenched by water. For the used steel, the A3 temperature is calculated to be 974 1C using Thermo-Calc, and the Ar3 temperature was measured to be 708 1C for recrystallized austenite and 723 1C for pancaked austenite at the cooling rate of 15 1C/s using thermal dilatometer. Cylindrical specimens of 6 mm in diameter and 15 mm in length were prepared for microstructure observation before isothermal bainitic treatment. Special compression samples were prepared for obtaining the final multi-phase microstructures as shown in Fig. 2, and the compression loading direction was indicated by P. The specimens for tensile tests were cut from the hot-deformed samples with a gage length of 4 mm 10 mm and thickness of 1.8 mm, as indicated by the dashed lines in Fig. 2. The samples were quenched by water from various stages of the hot compression schedule for the microstructure examination. The state of austenite prior to the transformation of γ-α was examined using a LEICA DC 100 optical microscope (OM) after hotetching using supersaturated picric acid with sodiumdodecylbenzene sulfonate. Microstructure examination for the following process was conducted using a Zeiss SUPRA 55 field-emission scanning electron microscope (SEM) and the microstructures were revealed by etching with 4% nital. The volume fraction of retained austenite (RA) and its lattice parameter was determined by a TTRIII multi-function X-ray diffraction (XRD) using Cu Kα radiation. The mean lattice parameter (a) determined using the (220)γ and (311) γ peaks was converted to carbon content by using the relationship
Fig. 1. Schematic diagrams of thermo-mechanical processing schedule for hot-rolled TRIP steel: (a) based on controlled cooling and (b) based on DTUA. (Tnr is the austenite no-recrystallization temperature, R is the recrystallized austenite, P is the pancaked austenite, and W.Q. is the water quenching.)
Fig. 2. Schematic diagrams of samples for compressive and tensile tests (unit: mm and P is the direction of compression loading).
a¼ 3.585þ 0.033 pct C [16]. The quantitative analysis of the volume fraction and the size of ferrite, bainite, or retained austenite/martensite (RA/M) constituents were performed using UTHSCSA image-analysis software (UT Health and Science center,
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San Antonio, San Antonio, TX) with scanning electron micrographs. The volume fraction of martensite was obtained as the difference between the volume fraction of RA/M and that of RA. In order to make the distinction possible between RA and martensite in the final multi-phase microstructures by SEM observation, samples were first annealed for 2 h at 200 1C and then etched with 4% nital [17]. After such a combination of heat treatment and etching, martensite showed rough surface and RA showed smooth surface. The content of Nb dissolved in the matrix was determined using a Varian 715-ES inductively coupled plasma–optical emission spectrometer (ICP–OES) by the method according to Ref. [18]. Carbon extraction replicas for precipitation analysis were prepared and examined by JEM-2010 transmission electron microscopy (TEM) and the sizes of the particles were measured using the software “Image Tool”. The room temperature tensile tests were performed in a Reger 3010 tensile tester at 1.2 mm/min, using the samples machined from the deformed samples with the dimension shown in Fig. 2. The stress–strain curves were drawn using the data collected by the extensometer of the tensile tester. The total elongation was determined according to the gage length of the tensile specimen before and after the tensile test.
3. Results and discussion 3.1. Microstructural evolution By austenitization at 1250 1C for 5 min, the average size of austenite grains of the tested steel was about 180 μm. By the 3-stage deformation in the austenitic recrystallization region, i.e. in the range of 1080–1000 1C, the average size of austenite was refined markedly to about 18 μm, as shown in Fig. 3a. With further deformation in austenitic nonrecrystallization region, i.e. at 900 1C, pancaked austenite was obtained as shown in Fig. 3b.
3.1.1. The thermo-mechanical process based on controlled cooling It is well known that the deformation in the austenitic nonrecrystallization region can accelerate the transformation of γ-α, resulting in finer ferrite grains. For the purpose of obtaining nearly the same volume fraction of ferrite, recrystallized austenite and pancaked austenite were cooled to 620 1C and 650 1C at a slow cooling rate of 3 1C/s respectively, as shown in Fig. 1a; the corresponding microstructures are shown in Fig. 4. The volume fractions of ferrite in these microstructures are about 40% and 42%, and the average grain sizes of ferrite are about 8.8 μm and 5.9 μm for recrystallized austenite and pancaked austenite, respectively. Then, the samples were cooled to 450 1C at 25 1C/s and held for 5 min, finally it was quenched by water. Fig. 5 shows the typical multi-phase microstructures for TRIP steels obtained by the thermo-mechanical process based on controlled cooling; the corresponding parameters of different constituents in the multiphase microstructures are summarized in Table 1. The volume fractions of ferrite are about 45% and 48% for recrystallized austenite and pancaked austenite respectively, indicating that only a small part of remaining austenite transformed into ferrite during cooling to 450 1C at 25 1C/s. At the same time, the average size of ferrite grains changed slightly. In the case of recrystallized austenite (Fig. 5a), bainite packets consisted of longer bainitic ferrite laths with nearly single orientation. In contrast, the deformation in the austenite nonrecrystallization region led to the refinement of bainite packets, which consisted of shorter bainitic ferrite laths with chaotic orientation, as shown in Fig. 5b. Moreover, the volume fraction of RA in the case of pancaked austenite is relatively higher, but its carbon content is relatively lower. Besides those located between bainitic ferrite laths, there are some blocky RA/M islands located at the polygonal ferrite/ bainite interface in the microstructures shown in Fig. 5. Fig. 6 shows the morphologies of blocky RA/M islands in the multiphase microstructures, treated by tempering at 200 1C for 2 h. It is clear that most of the blocky RA/M islands in the case of
Fig. 3. Microstructures of the tested steel prior to the transformation of austenite to ferrite: (a) recrystallized austenite and (b) pancaked austenite.
Fig. 4. Microstructures of the tested steel with different states of austenite cooled to different temperatures at 3 1C/s and water quenched:(a) 620 1C for recrystallized austenite and (b) 650 1C for pancaked austenite (F is ferrite, M is martensite).
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Fig. 5. Multi-phase microstructures of the tested steel with different states of austenite based on controlled-cooling processing: (a) recrystallized austenite and (b) pancaked austenite. (F is ferrite, B is bainite, and RA/M is blocky retained austenite/martensite.)
Table 1 Microstructure parameters of different constituents in the multi-phase microstructures of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing. Austenite state
F (%)
B (%)
M (%)
RA (%)
C of RA (%)
dF (mm)
R P
45 48
40 38
8.8 6.9
6.2 7.1
1.31 1.26
8.8 5.9
Note: F is ferrite, B is bainite, M is martensite, RA is retained austenite, C of RA is the carbon content of retained austenite, and dF is the average grain size of ferrite.
Fig. 6. Morphologies of RA/M islands in the multi-phase microstructures shown in Fig. 5 tempered at 200 1C for 2 h: (a) recrystallized austenite and (b) pancaked austenite. (Retained austenite shows smooth surface and martensite shows rough surface.)
recrystallized austenite (Fig. 6a) were martensite, and quite a part of such blocky RA/M islands in the case of pancaked austenite was granular RA with a smooth surface as indicated by arrows in Fig. 6b. That should be attributed to the refinement of ferrite grains in the case of pancaked austenite [14]. Additionally, after austenitization at 1250 1C for 5 min, the amount of Nb dissolved in austenite was 0.0264%, i.e. about 70% of the total added amount of Nb was dissolved into austenite. For the recrystallized austenite shown in Fig. 3a, the amount of Nb dissolved in austenite was 0.0256%, i.e. about 97% of the total dissolved amount of Nb was still dissolved in austenite. For the pancaked austenite shown in Fig. 3b, the amount of Nb dissolved in austenite was 0.0232%, i.e. about 88% of the total dissolved amount of Nb was still dissolved in austenite. In the final multiphase microstructures of the tested steel based on controlledcooling processing, the content of Nb still in solution was 0.0156% and 0.0124% for recrystallized austenite and pancaked austenite respectively, indicating that the deformation in the austenitic nonrecrystallization region obviously promoted the precipitation of Nb. As shown in Fig. 7, there were many Nb-containing particles with the sizes of 5–35 nm in the final multi-phase microstructures.
3.1.2. The thermo-mechanical process based on DTUA Fig. 8 shows the microstructures of the tested steel with different states of austenite quenched at different stages in the
Fig. 7. TEM micrograph of Nb precipitates presented on carbon extraction replicas of the tested steel with pancaked austenite state based on controlled-cooling processing.
processing schedule of the thermo-mechanical process based on DTUA. After deformed to strain of 50% at 750 1C at 1 s 1, finegrained ferrite was formed with the average size of about 2.0 μm in both cases of the recrystallized austenite and the pancaked austenite. In the case of the recrystallized austenite, ferrite grains were only formed along the prior grain boundaries of austenite, as shown in Fig. 8a. As a contrast, the deformation bands interior of austenite grains could also be the nucleation site of ferrite in the case of the pancaked austenite, as shown in Fig. 8b. Thus, the volume fraction of ferrite formed by DTUA was higher in the latter
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Fig. 8. Microstructures of the tested steel with different states of austenite quenched at different stages in the schedule based on DTUA processing: (a), (c), (e) recrystallized austenite and (b), (d), (f) pancaked austenite; (a), (b) deformed of undercooled austenite to strain of 50% at 750 1C, (c), (d) then cooled to 450 1C at 25 1C/s and (e), (f) then held for 5 min at 450 1C. (F is ferrite, B is bainite, and RA/M is blocky retained austenite/martensite.)
case than in the former case, i.e. about 16% and 25% for recrystallized austenite and pancaked austenite, respectively. By hot deformation of undercooled austenite, many crystal defects were introduced into austenite due to the severe deformation, which increased the driving force and supplied extra nucleation sites for the subsequent transformation of γ-α [6]. As shown in Fig. 8c and d, the volume fractions of ferrite increased obviously during the rapid cooling from 750 1C to 450 1C at 25 1C/s. At the same time, the average size of ferrite grains had almost no change. In both cases of the recrystallized austenite and the pancaked austenite, the untransformed austenite was divided into many small regions and the difference in the volume fraction of ferrite became smaller. By delay at 450 1C for 5 min and water quenched, refined and uniform multi-phase microstructures were obtained in both cases, as shown in Fig. 8e and f, consisting of fine-grained ferrite with the average size of about 2.0 μm, small bainite packets with shorter bainitic ferrite laths, and some retained austenite and martensite. The corresponding parameters of different constituents in the multi-phase microstructures are summarized in Table 2. It is clear that the difference in the parameters of different constituents is smaller in the case of the thermo-mechanical process based on DTUA than that in the case of the thermomechanical process based on controlled cooling, that should be attributed to the conduct of hot deformation of undercooled austenite. Additionally, in comparison with that based on controlled-cooling
Table 2 Microstructure parameters of different constituents in the multi-phase microstructures of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing. Austenite state
F (%)
B (%)
M (%)
RA (%)
C of RA (%)
dF (mm)
R P
51 54
37 35
5.8 4.6
6.2 6.4
1.2 1.24
2.0 2.0
processing (Fig. 6), granular RA with smaller average size less than 1 μm in the multi-phase microstructures based on DTUA processing located not only between the bainitic ferrite laths and at the polygonal ferrite/bainite interface, but also between the polygonal ferrite grains, as indicated by arrows in Fig. 9. In the final multi-phase microstructures of the tested steel based on DTUA processing, there were also many Nb-containing particles with the morphology similar to that based on controlledcooling processing, as shown in Fig. 10. The content of Nb still in solution was 0.0174% and 0.0163% for recrystallized austenite and pancaked austenite, respectively. That is, the amount of Nb precipitated during the thermo-mechanical process based on DTUA was less than that during the thermo-mechanical process based on controlled cooling, due to the fast cooling rate, 25 1C/s vs. 3 1C/s.
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Fig. 9. Morphologies of RA/M islands in the multi-phase microstructures shown in Fig. 8f tempered at 200 1C for 2 h.
Fig. 10. TEM micrograph of Nb precipitates presented on carbon extraction replicas of the tested steel with pancaked austenite state based on DTUA processing.
3.2. Mechanical properties 3.2.1. Hot-rolled TRIP steel based on controlled-cooling processing The tensile curves of the tested steel with different states of austenite treated by the thermo-mechanical process based on controlled cooling are shown in Fig. 11a, and the corresponding mechanical properties are illustrated in Table 3. Both the curves show a continuous yielding and similar shapes. Although the average size of ferrite grains of the tested steel with the recrystallized austenite was larger than that in the case of the pancaked austenite, the yield strength (s0.2) in the case of the recrystallized austenite was slightly higher than that in the case of the pancaked austenite. That may be attributed to the morphology of bainite in the multi-phase microstructure in the case of the recrystallized austenite, and those blocky martensite islands located at the polygonal ferrite/bainite interface [14]. However, the ultimate tensile strength (sb), uniform elongation (δu) and total elongation (δtal) in the case of the recrystallized austenite were lower than that in the case of the pancaked austenite. The lower yield strength and higher ultimate tensile strength in the case of the pancaked austenite resulted in a lower yield ratio (s0.2/sb), indicating the enhancement of the work-hardening capability. The work-hardening capability of TRIP steel is usually characterized by the incremental work-hardening exponent nincr defined as nincr ¼ d(ln s)/d(ln ε), where s is the true stress and ε is the true strain [19]. The incremental work-hardening exponenttrue strain curves corresponding to the tensile curves in Fig. 11a are shown in (b). The curves in the two cases exhibited a similar tendency, i.e. nincr value increased initially at low strain and decreased with increasing strain. The peak value of nincr was higher in the case of the pancaked austenite, indicating that more RA transformed to martensite under the same condition of plastic deformation. The difference in the behavior of RA transformed to martensite between the two cases may be attributed to the variety of the
Fig. 11. (a) Tensile curves and (b) incremental strain-hardening exponent curves of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing.
Table 3 Mechanical properties of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing. Austenite state
s0.2
sb
δu
δtal
(MPa)
(MPa)
(%)
(%)
R P
529 483
798 802
13.3 14.8
28 29.2
(GPa%)
s0.2/ sb
22.3 23.4
0.66 0.60
sb δtal
Note: s0.2 is the yield strength, sb is the ultimate tensile strength, δu is the uniform elongation, and δtal is the total elongation.
location and morphology of the RA. In the case of the pancaked austenite, a large part of RA was located at the polygonal ferrite/ bainite interface. Compared with that located between the bainitic ferrite laths, such RA had lower carbon contents due to the absence of carbon enrichment during bainite transformation [20]. In consequence, it had a higher tendency to transform to martensite during plastic deformation, therefore, expressed a stronger TRIP effect at small strain. As a result, the improvement of work-hardening capability led to an increase in the strength and ductility of the tested steel in the case of the pancaked austenite.
3.2.2. Hot-rolled TRIP steel based on DTUA processing The tensile curves of the tested steel with different states of austenite treated by the thermo-mechanical process based on DTUA are shown in Fig. 12a, and the corresponding mechanical
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Fig. 13. Incremental strain-hardening exponent curves of the tested steel with pancaked austenite based on controlled-cooling processing or based on DTUA processing.
Fig. 12. (a) Tensile curves and (b) incremental strain-hardening exponent curves of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing.
Table 4 Mechanical properties of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing. Austenite state
s0.2 (MPa)
R P
590 594
sb (MPa)
835 841
δu (%) 14.6 16.1
δtal (%) 32.2 33.4
(GPa%)
s0.2/ sb
26.9 28.1
0.71 0.71
sb δtal
properties are illustrated in Table 4. Both the curves also show a continuous yielding and similar shapes. The yield strengths in the two cases were similar, but the ultimate tensile strength, the uniform elongation and the total elongation in the case of the recrystallized austenite were lower than that in the case of the pancaked austenite. It indicated that the work-hardening capability of the tested steel in the case of the recrystallized austenite was slightly lower, as shown in Fig. 12b. The volume fraction and morphology of each constituent were similar in the multi-phase microstructures in the two cases; such difference in the workhardening capability may be attributed to the difference in the detailed state of RA in the multi-phase microstructures in both cases. In comparison with the tested steels by the thermo-mechanical process based on controlled cooling, the improvement of the mechanical properties of the test steel by the thermomechanical process based on DTUA should be mainly attributed to the refinement of the multi-phase microstructures as well as
the more uniform distribution of various constituents, i.e. ferrite, bainite, RA and martensite. Moreover, the stabilities of RA in the multi-phase microstructures of the tested steels by the thermomechanical process based on controlled cooling or based on DTUA were different, as indicated by the incremental strain-hardening exponent curves of the tested steel with the pancaked austenite shown in Fig. 13. In the case based on controlled-cooling processing, the value of nincr increased rapidly at small strain and then decreased gradually, suggesting that RA had low stability and the transformation of RA to martensite occurred easily [21]. On the contrary, the value of nincr increased gradually with strain and was up to the peak at relatively large strain in the case based on DTUA processing, suggesting that there were significant quantities of stable RA in the multi-phase microstructures of the tested steels by the thermo-mechanical process based on DTUA. In both cases mentioned above, the carbon contents of RA were similar, 1.26% and 1.24%, thus their chemical stabilities should be similar. The remarkable difference in their stabilities during the plastic deformation should be attributed to their different morphologies. The average size of RA in the case based on controlled-cooling processing was about 1.5 μm, compared to that less than 1 μm in the case based on DTUA processing. Smaller RA crystals provided less potential nucleation sites for the formation of martensite [22–24]; consequently, a greater total driving force for the nucleation of martensite was required. As a whole, in comparison with that based on controlled cooling, the influence of the state of austenite prior to the transformation of γ-α on the final multi-phase microstructures became weak for the tested steel by the thermo-mechanical process based on DTUA, due to the deformation activation energy introduced by hot deformation of undercooled austenite. It can be concluded that by the thermo-mechanical process based on DTUA, the requirement on the microstructure control of austenite in high temperature region is decreased and the processing stability is increased. Moreover, the refinement of the multi-phase microstructures by the thermo-mechanical process based on DTUA resulted in the marked improvement of the mechanical properties of hot-rolled Nb-microalloyed TRIP steels.
4. Conclusions Hot rolled Nb-microalloyed TRIP steels were heat treated by the thermo-mechanical process based on controlled cooling or based on dynamic transformation of undercooled austenite (DTUA), with the different states of austenite prior to the transformation of
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austenite to ferrite (γ-α), i.e. recrystallized austenite or pancaked austenite. In the case of the process based on controlled cooling, the transformation of γ-α was faster with the state of pancaked austenite, resulting in the refined multi-phase microstructures of the tested steel, as well as the improved combination of strength and ductility. On the contrary, in the case of the process based on DTUA, the influence of the states of austenite prior to the transformation of γ-α on the multi-phase microstructures and the mechanical properties of the tested steel became weak. The much refined and uniform microstructures with smaller sizes of ferrite and bainitic packets and the more stable retained austenite were obtained by the process based on DTUA, regardless of the state of austenite prior to the transformation of γ-α, resulting in the markedly improved combination of strength and ductility of the tested steel. Acknowledgments Financial supports of The National Basic Research Program of China (2010CB630801) and the State Key Laboratory for Advanced Metals and Materials are gratefully acknowledged. References [1] A.K. Srivastava, D. Bhattacharjee, G. Jha, N. Gope, S.B. Singh, Mater. Sci. Eng. A 445–446 (2007) 549–557. [2] P.J. Jacques, E. Girault, A. Mertens, B. Verlinden, J. Van Humbeeck, F. Delannay, ISIJ Int. 41 (2001) 1068–1074. [3] D.W. Suh, S.J. Park, C.S. Oh, S.J. Kim, Scr. Mater. 57 (2007) 1097–1100. [4] H.B. Ryu, J.G. Speer, J.P. Wise, Metall. Mater. Trans. A 33A (2002) 2811–2816.
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[5] I.B. Timokhina, E.V. Pereloma, P.D. Hodgson, Mater. Sci. Technol. 17 (2001) 135–140. [6] L.F. Li, X.J. Zhang, W.Y. Yang, Z.Q. Sun, Metall. Mater. Trans. A 44A (2013) 4337–4345. [7] S. Hashimoto, S. Ikeda, K.I. Sugimoto, S. Miyake, ISIJ Int. 44 (2004) 1590–1598. [8] E.V. Pereloma, I.B. Timokhina, P.D. Hodgson, Mater. Sci. Eng. A 273–275 (1999) 448–452. [9] Y.Y. Yin, W.Y. Yang, L.F. Li, Z.Q. Sun, X.T. Wang, Acta Metall. Sin. 44 (2008) 1299–1304. [10] W.Y. Yang, L.F. Li, Y.Y. Yin, Z.Q. Sun, X.T. Wang, Mater. Sci. Forum 654–656 (2010) 250–253. [11] Y.Y. Yin, W.Y. Yang, L.F. Li, Z.Q. Sun, X.T. Wang, Acta Metall. Sin. 46 (2010) 155–160. [12] B. Fu, W.Y. Yang, M.Y. Lu, Q. Feng, L.F. Li, Z.Q. Sun, Mater. Sci. Eng. A 536 (2012) 265–268. [13] E. De Moor, P.J. Gibbs, J.G. Speer, D.K. Matlock, J.G. Schroth, Iron Steel Technol. 7 (11) (2010) 132–144. [14] I.B. Timokhina, P.D. Hodgson, E.V. Pereloma, Metall. Mater. Trans. A 34A (2003) 1599–1609. [15] A. Zarei-Hanzaki, P.D. Hodgson, S. Yue, Metall. Mater. Trans. A 28A (1997) 2405–2414. [16] E. Emadoddin, A. Akbarzadeh, G.H. Daneshi, Mater. Charact. 57 (2006) 408–413. [17] E. Girault, P.J. Jacques, Ph. Harlet, K. Mols, J. Van Humbeeck, E. Aernoudt, F. Delannay, Mater. Charact. 40 (1998) 111–118. [18] Y.B. Cao, F.R. Xiao, G.Y. Qiao, X.B. Zhang, B. Liao, Mater. Sci. Eng. A 530 (2011) 277–284. [19] P.J. Jacques, Q. Furnemont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681–3693. [20] I.B. Timokhina, P.D. Hodgson, E.V. Pereloma, Metall. Mater. Trans. A 35A (2004) 2331–2341. [21] E.M. Bellhouse, J.R. McDermid, Metall. Mater. Trans. A 41A (2010) 1460–1473. [22] M.L. Brandt, G.B. Olson, Iron Steelmak. 20 (1993) 55–60. [23] E. Jimenez-Melero, N.H. Van Dijk, L. Zhao, J. Sietsma, S.E. Offerman, J.P. Wright, S. Van der Zwaag, Scr. Mater. 56 (2007) 421–424. [24] M. Mukherjee, O.N. Mohanty, S.I. Hashimoto, T. Hojo, K.I. Sugimoto, ISIJ Int. 46 (2006) 316–324.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructures and mechanical properties of hot-rolled Nb-microalloyed TRIP steels by different thermo-mechanical processes Qingxiao Feng a, Longfei Li a,n, Wangyue Yang b, Zuqing Sun a a b
State Key Laboratory for Advanced Metals and Materials, University of Science & Technology Beijing, Beijing 100083, China School of Materials Science and Engineering, University of Science & Technology Beijing, Beijing 100083, China
art ic l e i nf o
a b s t r a c t
Article history: Received 20 November 2013 Received in revised form 28 February 2014 Accepted 10 March 2014 Available online 19 March 2014
The microstructural evolution and mechanical properties of hot-rolled Nb-microalloyed TRIP steel with different states of austenite (recrystallized/pancaked) prior to the transformation of austenite to ferrite were investigated under the thermo-mechanical process based on controlled cooling or based on dynamic transformation of undercooled austenite (DTUA) by hot compression tests using a Gleeble-1500 hot simulator, in combination with OM, SEM, TEM, XRD and ICP–OES. The results indicated that the influence of the states of austenite prior to the transformation of austenite to ferrite on the multi-phase microstructures and the mechanical properties of the tested steel became weak in the case of the process based on DTUA, in comparison with the case of the process based on controlled cooling. The much refined and uniform microstructures with smaller sizes of ferrite and bainitic packets and the more stable retained austenite were obtained by the process based on DTUA, resulting in the markedly improved combination of strength and ductility of the tested steel. & 2014 Elsevier B.V. All rights reserved.
Keywords: Hot-rolled TRIP steel Nb-microalloyed Dynamic transformation of undercooled austenite Retained austenite Mechanical properties
1. Introduction There are two principal ways for producing TRIP steels, continuous annealing of cold-rolled material and hot strip rolling. Most previous studies have concentrated on the production of cold-rolled TRIP steels and made a lot of achievements [1–3], but the production of hot-rolled TRIP steels is relatively simple and cost-effective [4,5]. One key step for obtaining the multiphase microstructure of TRIP steels, which consists of polygonal ferrite matrix, bainite, retained austenite and some martensite, is to obtain the mixed microstructure consisting of ferrite and austenite with nearly same volume fraction of about 50%, before the isothermal bainitic treatment [6]. For conventional hot-rolled process of TRIP steels, the control of such mixed microstructure is commonly developed by adjusting cooling rate and time through multi-stage cooling after hot rolling of austenite [7,8], that is relatively difficult to be realized precisely in the industry, limiting the mass production of hot-rolled TRIP steel. Recently, a novel thermo-mechanical process to manufacture hotrolled TRIP steel was developed in our previous studies based on dynamic transformation of undercooled austenite (DTUA) [9–12]. The typical feature of this process is that the formation of the mixed microstructure of ferrite and austenite can be well controlled by the
n
Corresponding author. Tel.: þ 86 10 62334862; fax: þ 86 10 62333447. E-mail address: [email protected] (L. Li).
http://dx.doi.org/10.1016/j.msea.2014.03.051 0921-5093/& 2014 Elsevier B.V. All rights reserved.
applied strain of hot deformation of undercooled austenite. In comparison with controlling the cooling rate of undercooled austenite, the strain is relatively easy to be controlled accurately in the industry, and the multi-phase microstructure of hot-rolled TRIP steel could be obtained easily and be refined markedly by the thermo-mechanical process based on DTUA [9]. Hot-rolled TRIP steels obtained by such process demonstrate a good combination of strength and ductility, because of the refined and uniform multi-phase microstructure. For example, the tensile strength and the total elongation of C–Mn–Al–Si TRIP steel are 780 MPa and 34% respectively [10], which is nearly same as the mechanical properties of cold-rolled TRIP steel with the similar composition. With the addition of microalloying element Nb in C–Mn– Al–Si TRIP steel [12], the tensile strength can be enhanced to 845 MPa without any deterioration in the total elongation, due to the further refinement of the multi-phase microstructure. The product of tensile strength and total elongation is up to 28.7 GPa%, which is close to the requirement of the mechanical properties for the “Third-Generation” advanced high-strength steel, i.e. 30 GPa% [13]. Commonly, the main aim of the addition of Nb in steels is to adjust the state of austenite prior to the transformation of austenite to ferrite (γ-α) and thus accelerate the formation of ferrite grains, resulting in the structural refinement in the final products. For example, in the case of conventional hot-rolled 0.2C– 1.55Mn–1.55Si TRIP steel [14], it is found that the deformation in the nonrecrystallized austenite region can lead to the refinement of ferrite grains, and affect the transformation kinetics and morphology of bainite as well as the distribution of retained
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austenite in the multi-phase microstructure. As a result of that, TRIP steels with different states of austenite prior to the transformation of γ-α exhibited different mechanical properties. Moreover, it is found that in the case of conventional hot-rolled Nbmicroalloyed C–Mn–Si TRIP steel [15], the deformation in the nonrecrystallized austenite region can promote the precipitation of Nb-containing particles and result in a decrease in the volume fraction of retained austenite due to a reduction of Nb in the solid solution and, therefore, a decrease in parent austenite stability. In the previous work [9–12], the state of austenite prior to the dynamic transformation of undercooled austenite is only equiaxed recrystallized austenite grains for C–Mn–Al–Si hot-rolled TRIP steels with and without Nb. The key objective of this work, therefore, is to study the effect of the different states of austenite prior to the transformation of γ-α on the microstructural evolution and the mechanical properties of hot-rolled Nb-microalloyed TRIP steel based on DTUA processing, in comparison with that based on controlled-cooling processing.
2. Materials and experimental method The used material is a C–Mn–Al–Si–Nb steel by vacuum melting, containing 0.20C, 1.50Mn, 0.50Si, 1.04Al, 0.038Nb and the balance Fe (wt%). The casting blanks were first reheated to 1200 1C and held for 2 h, then hot-forged at 1150–900 1C followed by air-cooling to room temperature, and then machined into different specimens for hot-simulation test. The hot-simulation tests were performed using a Gleeble-1500 hot simulator, with the thermo-mechanical processes shown in Fig. 1. After austenitization at 1250 1C for 5 min, the samples were subjected to a 3-stage deformation at the austenitic recrystallization region from 1080 1C to 1000 1C; in each stage the applied strain was 30% at 2 s 1 and then held for 10 s to obtain recrystallized austenite (R) through static recrystallization. Then the recrystallized austenite was cooled to austenitic nonrecrystallization region, i.e. 900 1C, deformed to strain of 30% at 2 s 1 to obtain pancaked austenite (P). Next, the recrystallized austenite or pancaked austenite underwent rest of the thermo-mechanical process based on controlled cooling (Fig. 1a) or based on DTUA (Fig. 1b). After the isothermal bainitic treatment at 450 1C for 5 min, the samples were quenched by water. For the used steel, the A3 temperature is calculated to be 974 1C using Thermo-Calc, and the Ar3 temperature was measured to be 708 1C for recrystallized austenite and 723 1C for pancaked austenite at the cooling rate of 15 1C/s using thermal dilatometer. Cylindrical specimens of 6 mm in diameter and 15 mm in length were prepared for microstructure observation before isothermal bainitic treatment. Special compression samples were prepared for obtaining the final multi-phase microstructures as shown in Fig. 2, and the compression loading direction was indicated by P. The specimens for tensile tests were cut from the hot-deformed samples with a gage length of 4 mm 10 mm and thickness of 1.8 mm, as indicated by the dashed lines in Fig. 2. The samples were quenched by water from various stages of the hot compression schedule for the microstructure examination. The state of austenite prior to the transformation of γ-α was examined using a LEICA DC 100 optical microscope (OM) after hotetching using supersaturated picric acid with sodiumdodecylbenzene sulfonate. Microstructure examination for the following process was conducted using a Zeiss SUPRA 55 field-emission scanning electron microscope (SEM) and the microstructures were revealed by etching with 4% nital. The volume fraction of retained austenite (RA) and its lattice parameter was determined by a TTRIII multi-function X-ray diffraction (XRD) using Cu Kα radiation. The mean lattice parameter (a) determined using the (220)γ and (311) γ peaks was converted to carbon content by using the relationship
Fig. 1. Schematic diagrams of thermo-mechanical processing schedule for hot-rolled TRIP steel: (a) based on controlled cooling and (b) based on DTUA. (Tnr is the austenite no-recrystallization temperature, R is the recrystallized austenite, P is the pancaked austenite, and W.Q. is the water quenching.)
Fig. 2. Schematic diagrams of samples for compressive and tensile tests (unit: mm and P is the direction of compression loading).
a¼ 3.585þ 0.033 pct C [16]. The quantitative analysis of the volume fraction and the size of ferrite, bainite, or retained austenite/martensite (RA/M) constituents were performed using UTHSCSA image-analysis software (UT Health and Science center,
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San Antonio, San Antonio, TX) with scanning electron micrographs. The volume fraction of martensite was obtained as the difference between the volume fraction of RA/M and that of RA. In order to make the distinction possible between RA and martensite in the final multi-phase microstructures by SEM observation, samples were first annealed for 2 h at 200 1C and then etched with 4% nital [17]. After such a combination of heat treatment and etching, martensite showed rough surface and RA showed smooth surface. The content of Nb dissolved in the matrix was determined using a Varian 715-ES inductively coupled plasma–optical emission spectrometer (ICP–OES) by the method according to Ref. [18]. Carbon extraction replicas for precipitation analysis were prepared and examined by JEM-2010 transmission electron microscopy (TEM) and the sizes of the particles were measured using the software “Image Tool”. The room temperature tensile tests were performed in a Reger 3010 tensile tester at 1.2 mm/min, using the samples machined from the deformed samples with the dimension shown in Fig. 2. The stress–strain curves were drawn using the data collected by the extensometer of the tensile tester. The total elongation was determined according to the gage length of the tensile specimen before and after the tensile test.
3. Results and discussion 3.1. Microstructural evolution By austenitization at 1250 1C for 5 min, the average size of austenite grains of the tested steel was about 180 μm. By the 3-stage deformation in the austenitic recrystallization region, i.e. in the range of 1080–1000 1C, the average size of austenite was refined markedly to about 18 μm, as shown in Fig. 3a. With further deformation in austenitic nonrecrystallization region, i.e. at 900 1C, pancaked austenite was obtained as shown in Fig. 3b.
3.1.1. The thermo-mechanical process based on controlled cooling It is well known that the deformation in the austenitic nonrecrystallization region can accelerate the transformation of γ-α, resulting in finer ferrite grains. For the purpose of obtaining nearly the same volume fraction of ferrite, recrystallized austenite and pancaked austenite were cooled to 620 1C and 650 1C at a slow cooling rate of 3 1C/s respectively, as shown in Fig. 1a; the corresponding microstructures are shown in Fig. 4. The volume fractions of ferrite in these microstructures are about 40% and 42%, and the average grain sizes of ferrite are about 8.8 μm and 5.9 μm for recrystallized austenite and pancaked austenite, respectively. Then, the samples were cooled to 450 1C at 25 1C/s and held for 5 min, finally it was quenched by water. Fig. 5 shows the typical multi-phase microstructures for TRIP steels obtained by the thermo-mechanical process based on controlled cooling; the corresponding parameters of different constituents in the multiphase microstructures are summarized in Table 1. The volume fractions of ferrite are about 45% and 48% for recrystallized austenite and pancaked austenite respectively, indicating that only a small part of remaining austenite transformed into ferrite during cooling to 450 1C at 25 1C/s. At the same time, the average size of ferrite grains changed slightly. In the case of recrystallized austenite (Fig. 5a), bainite packets consisted of longer bainitic ferrite laths with nearly single orientation. In contrast, the deformation in the austenite nonrecrystallization region led to the refinement of bainite packets, which consisted of shorter bainitic ferrite laths with chaotic orientation, as shown in Fig. 5b. Moreover, the volume fraction of RA in the case of pancaked austenite is relatively higher, but its carbon content is relatively lower. Besides those located between bainitic ferrite laths, there are some blocky RA/M islands located at the polygonal ferrite/ bainite interface in the microstructures shown in Fig. 5. Fig. 6 shows the morphologies of blocky RA/M islands in the multiphase microstructures, treated by tempering at 200 1C for 2 h. It is clear that most of the blocky RA/M islands in the case of
Fig. 3. Microstructures of the tested steel prior to the transformation of austenite to ferrite: (a) recrystallized austenite and (b) pancaked austenite.
Fig. 4. Microstructures of the tested steel with different states of austenite cooled to different temperatures at 3 1C/s and water quenched:(a) 620 1C for recrystallized austenite and (b) 650 1C for pancaked austenite (F is ferrite, M is martensite).
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Fig. 5. Multi-phase microstructures of the tested steel with different states of austenite based on controlled-cooling processing: (a) recrystallized austenite and (b) pancaked austenite. (F is ferrite, B is bainite, and RA/M is blocky retained austenite/martensite.)
Table 1 Microstructure parameters of different constituents in the multi-phase microstructures of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing. Austenite state
F (%)
B (%)
M (%)
RA (%)
C of RA (%)
dF (mm)
R P
45 48
40 38
8.8 6.9
6.2 7.1
1.31 1.26
8.8 5.9
Note: F is ferrite, B is bainite, M is martensite, RA is retained austenite, C of RA is the carbon content of retained austenite, and dF is the average grain size of ferrite.
Fig. 6. Morphologies of RA/M islands in the multi-phase microstructures shown in Fig. 5 tempered at 200 1C for 2 h: (a) recrystallized austenite and (b) pancaked austenite. (Retained austenite shows smooth surface and martensite shows rough surface.)
recrystallized austenite (Fig. 6a) were martensite, and quite a part of such blocky RA/M islands in the case of pancaked austenite was granular RA with a smooth surface as indicated by arrows in Fig. 6b. That should be attributed to the refinement of ferrite grains in the case of pancaked austenite [14]. Additionally, after austenitization at 1250 1C for 5 min, the amount of Nb dissolved in austenite was 0.0264%, i.e. about 70% of the total added amount of Nb was dissolved into austenite. For the recrystallized austenite shown in Fig. 3a, the amount of Nb dissolved in austenite was 0.0256%, i.e. about 97% of the total dissolved amount of Nb was still dissolved in austenite. For the pancaked austenite shown in Fig. 3b, the amount of Nb dissolved in austenite was 0.0232%, i.e. about 88% of the total dissolved amount of Nb was still dissolved in austenite. In the final multiphase microstructures of the tested steel based on controlledcooling processing, the content of Nb still in solution was 0.0156% and 0.0124% for recrystallized austenite and pancaked austenite respectively, indicating that the deformation in the austenitic nonrecrystallization region obviously promoted the precipitation of Nb. As shown in Fig. 7, there were many Nb-containing particles with the sizes of 5–35 nm in the final multi-phase microstructures.
3.1.2. The thermo-mechanical process based on DTUA Fig. 8 shows the microstructures of the tested steel with different states of austenite quenched at different stages in the
Fig. 7. TEM micrograph of Nb precipitates presented on carbon extraction replicas of the tested steel with pancaked austenite state based on controlled-cooling processing.
processing schedule of the thermo-mechanical process based on DTUA. After deformed to strain of 50% at 750 1C at 1 s 1, finegrained ferrite was formed with the average size of about 2.0 μm in both cases of the recrystallized austenite and the pancaked austenite. In the case of the recrystallized austenite, ferrite grains were only formed along the prior grain boundaries of austenite, as shown in Fig. 8a. As a contrast, the deformation bands interior of austenite grains could also be the nucleation site of ferrite in the case of the pancaked austenite, as shown in Fig. 8b. Thus, the volume fraction of ferrite formed by DTUA was higher in the latter
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Fig. 8. Microstructures of the tested steel with different states of austenite quenched at different stages in the schedule based on DTUA processing: (a), (c), (e) recrystallized austenite and (b), (d), (f) pancaked austenite; (a), (b) deformed of undercooled austenite to strain of 50% at 750 1C, (c), (d) then cooled to 450 1C at 25 1C/s and (e), (f) then held for 5 min at 450 1C. (F is ferrite, B is bainite, and RA/M is blocky retained austenite/martensite.)
case than in the former case, i.e. about 16% and 25% for recrystallized austenite and pancaked austenite, respectively. By hot deformation of undercooled austenite, many crystal defects were introduced into austenite due to the severe deformation, which increased the driving force and supplied extra nucleation sites for the subsequent transformation of γ-α [6]. As shown in Fig. 8c and d, the volume fractions of ferrite increased obviously during the rapid cooling from 750 1C to 450 1C at 25 1C/s. At the same time, the average size of ferrite grains had almost no change. In both cases of the recrystallized austenite and the pancaked austenite, the untransformed austenite was divided into many small regions and the difference in the volume fraction of ferrite became smaller. By delay at 450 1C for 5 min and water quenched, refined and uniform multi-phase microstructures were obtained in both cases, as shown in Fig. 8e and f, consisting of fine-grained ferrite with the average size of about 2.0 μm, small bainite packets with shorter bainitic ferrite laths, and some retained austenite and martensite. The corresponding parameters of different constituents in the multi-phase microstructures are summarized in Table 2. It is clear that the difference in the parameters of different constituents is smaller in the case of the thermo-mechanical process based on DTUA than that in the case of the thermomechanical process based on controlled cooling, that should be attributed to the conduct of hot deformation of undercooled austenite. Additionally, in comparison with that based on controlled-cooling
Table 2 Microstructure parameters of different constituents in the multi-phase microstructures of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing. Austenite state
F (%)
B (%)
M (%)
RA (%)
C of RA (%)
dF (mm)
R P
51 54
37 35
5.8 4.6
6.2 6.4
1.2 1.24
2.0 2.0
processing (Fig. 6), granular RA with smaller average size less than 1 μm in the multi-phase microstructures based on DTUA processing located not only between the bainitic ferrite laths and at the polygonal ferrite/bainite interface, but also between the polygonal ferrite grains, as indicated by arrows in Fig. 9. In the final multi-phase microstructures of the tested steel based on DTUA processing, there were also many Nb-containing particles with the morphology similar to that based on controlledcooling processing, as shown in Fig. 10. The content of Nb still in solution was 0.0174% and 0.0163% for recrystallized austenite and pancaked austenite, respectively. That is, the amount of Nb precipitated during the thermo-mechanical process based on DTUA was less than that during the thermo-mechanical process based on controlled cooling, due to the fast cooling rate, 25 1C/s vs. 3 1C/s.
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Fig. 9. Morphologies of RA/M islands in the multi-phase microstructures shown in Fig. 8f tempered at 200 1C for 2 h.
Fig. 10. TEM micrograph of Nb precipitates presented on carbon extraction replicas of the tested steel with pancaked austenite state based on DTUA processing.
3.2. Mechanical properties 3.2.1. Hot-rolled TRIP steel based on controlled-cooling processing The tensile curves of the tested steel with different states of austenite treated by the thermo-mechanical process based on controlled cooling are shown in Fig. 11a, and the corresponding mechanical properties are illustrated in Table 3. Both the curves show a continuous yielding and similar shapes. Although the average size of ferrite grains of the tested steel with the recrystallized austenite was larger than that in the case of the pancaked austenite, the yield strength (s0.2) in the case of the recrystallized austenite was slightly higher than that in the case of the pancaked austenite. That may be attributed to the morphology of bainite in the multi-phase microstructure in the case of the recrystallized austenite, and those blocky martensite islands located at the polygonal ferrite/bainite interface [14]. However, the ultimate tensile strength (sb), uniform elongation (δu) and total elongation (δtal) in the case of the recrystallized austenite were lower than that in the case of the pancaked austenite. The lower yield strength and higher ultimate tensile strength in the case of the pancaked austenite resulted in a lower yield ratio (s0.2/sb), indicating the enhancement of the work-hardening capability. The work-hardening capability of TRIP steel is usually characterized by the incremental work-hardening exponent nincr defined as nincr ¼ d(ln s)/d(ln ε), where s is the true stress and ε is the true strain [19]. The incremental work-hardening exponenttrue strain curves corresponding to the tensile curves in Fig. 11a are shown in (b). The curves in the two cases exhibited a similar tendency, i.e. nincr value increased initially at low strain and decreased with increasing strain. The peak value of nincr was higher in the case of the pancaked austenite, indicating that more RA transformed to martensite under the same condition of plastic deformation. The difference in the behavior of RA transformed to martensite between the two cases may be attributed to the variety of the
Fig. 11. (a) Tensile curves and (b) incremental strain-hardening exponent curves of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing.
Table 3 Mechanical properties of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on controlled-cooling processing. Austenite state
s0.2
sb
δu
δtal
(MPa)
(MPa)
(%)
(%)
R P
529 483
798 802
13.3 14.8
28 29.2
(GPa%)
s0.2/ sb
22.3 23.4
0.66 0.60
sb δtal
Note: s0.2 is the yield strength, sb is the ultimate tensile strength, δu is the uniform elongation, and δtal is the total elongation.
location and morphology of the RA. In the case of the pancaked austenite, a large part of RA was located at the polygonal ferrite/ bainite interface. Compared with that located between the bainitic ferrite laths, such RA had lower carbon contents due to the absence of carbon enrichment during bainite transformation [20]. In consequence, it had a higher tendency to transform to martensite during plastic deformation, therefore, expressed a stronger TRIP effect at small strain. As a result, the improvement of work-hardening capability led to an increase in the strength and ductility of the tested steel in the case of the pancaked austenite.
3.2.2. Hot-rolled TRIP steel based on DTUA processing The tensile curves of the tested steel with different states of austenite treated by the thermo-mechanical process based on DTUA are shown in Fig. 12a, and the corresponding mechanical
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Fig. 13. Incremental strain-hardening exponent curves of the tested steel with pancaked austenite based on controlled-cooling processing or based on DTUA processing.
Fig. 12. (a) Tensile curves and (b) incremental strain-hardening exponent curves of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing.
Table 4 Mechanical properties of the tested steel with recrystallized austenite (R) or pancaked austenite (P) based on DTUA processing. Austenite state
s0.2 (MPa)
R P
590 594
sb (MPa)
835 841
δu (%) 14.6 16.1
δtal (%) 32.2 33.4
(GPa%)
s0.2/ sb
26.9 28.1
0.71 0.71
sb δtal
properties are illustrated in Table 4. Both the curves also show a continuous yielding and similar shapes. The yield strengths in the two cases were similar, but the ultimate tensile strength, the uniform elongation and the total elongation in the case of the recrystallized austenite were lower than that in the case of the pancaked austenite. It indicated that the work-hardening capability of the tested steel in the case of the recrystallized austenite was slightly lower, as shown in Fig. 12b. The volume fraction and morphology of each constituent were similar in the multi-phase microstructures in the two cases; such difference in the workhardening capability may be attributed to the difference in the detailed state of RA in the multi-phase microstructures in both cases. In comparison with the tested steels by the thermo-mechanical process based on controlled cooling, the improvement of the mechanical properties of the test steel by the thermomechanical process based on DTUA should be mainly attributed to the refinement of the multi-phase microstructures as well as
the more uniform distribution of various constituents, i.e. ferrite, bainite, RA and martensite. Moreover, the stabilities of RA in the multi-phase microstructures of the tested steels by the thermomechanical process based on controlled cooling or based on DTUA were different, as indicated by the incremental strain-hardening exponent curves of the tested steel with the pancaked austenite shown in Fig. 13. In the case based on controlled-cooling processing, the value of nincr increased rapidly at small strain and then decreased gradually, suggesting that RA had low stability and the transformation of RA to martensite occurred easily [21]. On the contrary, the value of nincr increased gradually with strain and was up to the peak at relatively large strain in the case based on DTUA processing, suggesting that there were significant quantities of stable RA in the multi-phase microstructures of the tested steels by the thermo-mechanical process based on DTUA. In both cases mentioned above, the carbon contents of RA were similar, 1.26% and 1.24%, thus their chemical stabilities should be similar. The remarkable difference in their stabilities during the plastic deformation should be attributed to their different morphologies. The average size of RA in the case based on controlled-cooling processing was about 1.5 μm, compared to that less than 1 μm in the case based on DTUA processing. Smaller RA crystals provided less potential nucleation sites for the formation of martensite [22–24]; consequently, a greater total driving force for the nucleation of martensite was required. As a whole, in comparison with that based on controlled cooling, the influence of the state of austenite prior to the transformation of γ-α on the final multi-phase microstructures became weak for the tested steel by the thermo-mechanical process based on DTUA, due to the deformation activation energy introduced by hot deformation of undercooled austenite. It can be concluded that by the thermo-mechanical process based on DTUA, the requirement on the microstructure control of austenite in high temperature region is decreased and the processing stability is increased. Moreover, the refinement of the multi-phase microstructures by the thermo-mechanical process based on DTUA resulted in the marked improvement of the mechanical properties of hot-rolled Nb-microalloyed TRIP steels.
4. Conclusions Hot rolled Nb-microalloyed TRIP steels were heat treated by the thermo-mechanical process based on controlled cooling or based on dynamic transformation of undercooled austenite (DTUA), with the different states of austenite prior to the transformation of
Q. Feng et al. / Materials Science & Engineering A 605 (2014) 14–21
austenite to ferrite (γ-α), i.e. recrystallized austenite or pancaked austenite. In the case of the process based on controlled cooling, the transformation of γ-α was faster with the state of pancaked austenite, resulting in the refined multi-phase microstructures of the tested steel, as well as the improved combination of strength and ductility. On the contrary, in the case of the process based on DTUA, the influence of the states of austenite prior to the transformation of γ-α on the multi-phase microstructures and the mechanical properties of the tested steel became weak. The much refined and uniform microstructures with smaller sizes of ferrite and bainitic packets and the more stable retained austenite were obtained by the process based on DTUA, regardless of the state of austenite prior to the transformation of γ-α, resulting in the markedly improved combination of strength and ductility of the tested steel. Acknowledgments Financial supports of The National Basic Research Program of China (2010CB630801) and the State Key Laboratory for Advanced Metals and Materials are gratefully acknowledged. References [1] A.K. Srivastava, D. Bhattacharjee, G. Jha, N. Gope, S.B. Singh, Mater. Sci. Eng. A 445–446 (2007) 549–557. [2] P.J. Jacques, E. Girault, A. Mertens, B. Verlinden, J. Van Humbeeck, F. Delannay, ISIJ Int. 41 (2001) 1068–1074. [3] D.W. Suh, S.J. Park, C.S. Oh, S.J. Kim, Scr. Mater. 57 (2007) 1097–1100. [4] H.B. Ryu, J.G. Speer, J.P. Wise, Metall. Mater. Trans. A 33A (2002) 2811–2816.
21
[5] I.B. Timokhina, E.V. Pereloma, P.D. Hodgson, Mater. Sci. Technol. 17 (2001) 135–140. [6] L.F. Li, X.J. Zhang, W.Y. Yang, Z.Q. Sun, Metall. Mater. Trans. A 44A (2013) 4337–4345. [7] S. Hashimoto, S. Ikeda, K.I. Sugimoto, S. Miyake, ISIJ Int. 44 (2004) 1590–1598. [8] E.V. Pereloma, I.B. Timokhina, P.D. Hodgson, Mater. Sci. Eng. A 273–275 (1999) 448–452. [9] Y.Y. Yin, W.Y. Yang, L.F. Li, Z.Q. Sun, X.T. Wang, Acta Metall. Sin. 44 (2008) 1299–1304. [10] W.Y. Yang, L.F. Li, Y.Y. Yin, Z.Q. Sun, X.T. Wang, Mater. Sci. Forum 654–656 (2010) 250–253. [11] Y.Y. Yin, W.Y. Yang, L.F. Li, Z.Q. Sun, X.T. Wang, Acta Metall. Sin. 46 (2010) 155–160. [12] B. Fu, W.Y. Yang, M.Y. Lu, Q. Feng, L.F. Li, Z.Q. Sun, Mater. Sci. Eng. A 536 (2012) 265–268. [13] E. De Moor, P.J. Gibbs, J.G. Speer, D.K. Matlock, J.G. Schroth, Iron Steel Technol. 7 (11) (2010) 132–144. [14] I.B. Timokhina, P.D. Hodgson, E.V. Pereloma, Metall. Mater. Trans. A 34A (2003) 1599–1609. [15] A. Zarei-Hanzaki, P.D. Hodgson, S. Yue, Metall. Mater. Trans. A 28A (1997) 2405–2414. [16] E. Emadoddin, A. Akbarzadeh, G.H. Daneshi, Mater. Charact. 57 (2006) 408–413. [17] E. Girault, P.J. Jacques, Ph. Harlet, K. Mols, J. Van Humbeeck, E. Aernoudt, F. Delannay, Mater. Charact. 40 (1998) 111–118. [18] Y.B. Cao, F.R. Xiao, G.Y. Qiao, X.B. Zhang, B. Liao, Mater. Sci. Eng. A 530 (2011) 277–284. [19] P.J. Jacques, Q. Furnemont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681–3693. [20] I.B. Timokhina, P.D. Hodgson, E.V. Pereloma, Metall. Mater. Trans. A 35A (2004) 2331–2341. [21] E.M. Bellhouse, J.R. McDermid, Metall. Mater. Trans. A 41A (2010) 1460–1473. [22] M.L. Brandt, G.B. Olson, Iron Steelmak. 20 (1993) 55–60. [23] E. Jimenez-Melero, N.H. Van Dijk, L. Zhao, J. Sietsma, S.E. Offerman, J.P. Wright, S. Van der Zwaag, Scr. Mater. 56 (2007) 421–424. [24] M. Mukherjee, O.N. Mohanty, S.I. Hashimoto, T. Hojo, K.I. Sugimoto, ISIJ Int. 46 (2006) 316–324.