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Influences of alloying elements on warm deformation behavior of high-Mn TRIP steel with martensitic structure
Materials and Design 89 (2016) 665–675
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Influences of alloying elements on warm deformation behavior of high-Mn TRIP steel with martensitic structure Zhikai Guo, Longfei Li ⁎ State Key Laboratory for Advanced Metals and Materials, University of Science & Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 1 August 2015 Received in revised form 15 September 2015 Accepted 3 October 2015 Available online 9 October 2015 Keywords: Warm deformation behavior High-Mn TRIP steel Deformation activation energy DRX of ferrite Alloying elements
a b s t r a c t Warm deformation behaviors of high-Mn TRIP steels with martensitic structure were investigated using a Gleeble 1500 hot simulator within the temperature range of 550–650 °C and the strain rate range of 0.001–0.1 s−1. The influences of common alloying elements, i.e. C, Mn, Si and Al, on the flow stress and the microstructural evolution during warm deformation were analyzed. The results showed that the flow stress was increased with the increase in C content, Mn content, Si content or Al content at the same condition, and the deformation activation energy Q was increased with the increase in Mn content and decreased by the increase in C content. The increase in Si content or Al content showed slight influence on the deformation activation energy Q. During warm deformation the decomposition of martensite into ferrite and carbides occurred firstly, and then dynamic recrystallization (DRX) of ferrite and the reverse transformation of austenite took place. DRX of ferrite was retarded by the increase in Mn content or Si content and accelerated by the increase in C content. Meanwhile, the reverse transformation of austenite was promoted by the increase in Mn content and retarded by the increase in C content, Si content or Al content. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction High-Mn TRIP steels, having a microstructure which consists of ultrafine-grained ferrite matrix and a large amount of retained austenite (RA) with suitable stability, demonstrate a well balance of strength and ductility [1]. For example, Furukawa et al. reported that a 0.1C–5.0Mn steel with about 30 vol.% RA, obtained by annealing at 600 °C for 3 h after hot rolling of austenite, demonstrated the tensile strength and uniform elongation up to about 1000 MPa and 24%, respectively [2]. The researches of Gibbs et al. showed that a tensile strength of 876 MPa and a total elongation of 42% can be achieved in a 0.1C–7.0Mn steel by annealing at 600 °C for 1 week after cold rolling of martensite, with the product of tensile strength and total elongation up to 37 GPa% [3]. Thus, as one potential contender of the third generation of advanced high strength steel, high-Mn TRIP steels with 5.0–8.0 wt.% Mn have received a lot of attention [1–8]. The common production of high-Mn TRIP steels requires long-time intercritical annealing in the austenite plus ferrite twophase region with the duration of about several hours to several days, directly after hot rolling of austenite or after hot rolling of austenite plus cold rolling of martensite [2–3,5–6]. Such longtime intercritical annealing is unfavorable to the mass production of high-Mn TRIP steels. In the previous work [7], a new thermomechanical process based on warm deformation of martensite ⁎ Corresponding author. E-mail address: [email protected] (L. Li).
http://dx.doi.org/10.1016/j.matdes.2015.10.010 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
and short-time intercritical annealing was developed to product high-Mn TRIP steels with less cost. The decomposition of martensite into ferrite and carbides, dynamic recrystallization (DRX) of ferrite and the reverse transformation of austenite were accelerated by warm deformation, thus leading to the shorter time for subsequent annealing in the intercritical region, with the length of about one hour. As a promising process to obtain ultrafine ferrite grains through DRX of ferrite, warm deformation behavior of steels has been well investigated in IF steels [9], low-carbon steels [10–12], medium-carbon steels [13–15] and high-carbon steels [16–17], with the initial microstructure of single-phase ferrite grains, ferrite grains plus pearlite colonies, pearlite colonies or lath martensite structure. Among them, warm deformation of lath martensite structure has a noticeable advantage to obtain ultrafine ferrite grains with the reduced critical strain, owning to the high density of dislocations, the supersaturated solute C atoms and the ultrafine laths [13,18]. Commonly, the kinetics of DRX of ferrite and the average size of DRX grains are influenced definitely by the deformation temperature and the strain rate. The Zener–Hollomon parameter Z is commonly used to summarize the effects of the deformation temperature and the strain rate on DRX, which can be expressed as Z ¼ ε • expðQ=RTÞ , where, ε• , Q, R and T refer to strain rate (s− 1), deformation activation energy (J/mol), gas constant (8.3145 J/mol/K) and absolute temperature (K), respectively [19]. It is clearly that the value of Z has a direct relationship with the deformation activation energy Q, which is well known to be influenced by the element contents of the matrix, such as C and Mn contents in austenite grains or ferrite grains [20–21]. For example, the increase in Mn content from 0.48% to
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Table 1 Chemical compositions (wt.%) and equilibrium transformation temperatures (°C) of the used steels. Steel
C, %
Mn, %
Si, %
Al, %
Fe, %
A1, °C
A3, °C
A B C D E
0.083 0.089 0.18 0.081 0.11
4.62 7.60 4.60 4.48 4.89
0.021 0.022 – 1.52 –
– – – – 1.01
Bal. Bal. Bal. Bal. Bal.
592 563 635 622 644
736 681 723 771 799
1.49% resulted in the increase in the value of Q from 260.3 kJ/mol to 450.0 kJ/mol for low-carbon steels with C content of about 0.06% during warm deformation in the ferrite phase region, retarding the occurrence of DRX of ferrite and refining the average size of DRX ferrite grains [21]. The addition of Al with the content of 0.95% resulted in the increase in the value of Q from 333.6 kJ/mol to 387.6 kJ/mol for high-carbon steels with C content of about 1.0% during warm deformation of martensite, impeding the decomposition of martensite and refining the average sizes of ferrite grains and cementite particles [16]. As the common elements in high-Mn TRIP steels, the increase in Mn content or C content can result in the improvement in the mechanical properties of high-Mn TRIP steels, due to the increase in the stability of RA [1,2–3,7]. Meanwhile, the addition of Si and Al can delay the precipitation of carbides and thus increase the stability of RA in the conventional TRIP steels [22]. Recently, ferrite–austenite duplex lightweight steel, such as 0.3C–8.5Mn–5.6Al steel, is developed to further reduce vehicle weight in order to reduce exhaust emissions and improve fuel efficiency, which shows the operation of both TRIP and TWIP mechanisms during tensile straining and thus demonstrates well balance of strength and elongation [23–26]. In this work, the influences of the common element, i.e. C, Mn, Si and Al, on the warm deformation behavior of high-Mn TRIP steel were investigated by hot uniaxial compression tests, in order to optimizing the aforementioned thermo-mechanical process based on warm deformation and short-time intercritical annealing with deep understanding of the warm deformation behavior of martensite in high-Mn TRIP steel.
2. Experimental Table 1 shows the chemical compositions (wt.%) of the used steels with the corresponding equilibrium transformation temperatures A1 and A3. The increase in C content or Mn content leads to the narrowing of the intercritical region, while the increase in Si content or Al content leads to the broadening of the intercritical region. The casting blanks were reheated to 1200 °C and held for 2 h, then hot-forged at 1150– 900 °C followed by air-cooling to room temperature (RT). Due to the high Mn contents, martensitic structure can be obtained by air-cooling for the used steels [7]. Cylindrical specimens of 6 mm in diameter and 15 mm in length were machined from the blanks. After austenization at 800 °C for 30 min and air cooling to RT, warm compression tests were performed using a Gleeble 1500 hot simulator. The temperature of the sample during warm deformation was automatically controlled by computer using the feedback signals of a pair of thermocouples, which were welded on the specimen during deformation. The specimens were reheated to 550, 600 or 650 °C at 20 °C/s, isothermally held for 1 min, and then deformed to a true strain of 0.10, 0.22 and 0.69 at 0.001, 0.01 or 0.1 s−1, finally air-cooled to RT. The microstructure observation was carried out using a ZEISS SUPRA55 field-emission scanning electronic microscope (FE-SEM) equipped with an electron backscattered diffraction (EBSD) attachment on the sections parallel to the compression direction after etching with 4 vol.% Nital. The volume fraction of ferrite, and the size of ferrite and M/A islands were measured using Image-Pro Plus image-analysis software with scanning electron micrographs. The volume fraction of RA was determined by a TTRIII X-ray diffractometer using Cu Kα radiation [27].
3. Results and discussions 3.1. Flow curves of the used steels during warm deformation Fig. 1 shows the flow curves of the used steels deformed at 600 °C at 0.01 s−1 to a strain of 0.69. The flow curves of the used steels had the similar shape, i.e. the flow stress increased rapidly to a peak and then
Fig. 1. Flow curves of the used steels during warm deformation of martensite at 600 °C at 0.01 s−1.
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Fig. 2. Peak stresses of the used steels during warm deformation of martensite under different conditions: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
decreased continuously with straining, suggesting the appearance of dynamic softening. Fig. 2 summarizes the effect of the deformation temperature or the strain rate on the peak stress of the used steels deformed at various conditions, which is clear that the peak stress was decreased
with the increase in the deformation temperature or the decrease in the strain rate for each steel. Meanwhile, owning to the solid solute strengthening effect of alloying elements [28–29], the peak stress of each of the steels B–E was higher than that of the steel A at the same condition.
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Table 2 Values of deformation activation energy Q and constants B and β for the used steels during warm deformation of martensite. Steel
Q, kJ/mol
B, s−1
β
A B C D E
336.6 419.3 311.0 331.5 332.0
1.554 × 1014 4.431 × 1018 9.147 × 1012 1.668 × 1014 1.201 × 1014
0.02163 0.02264 0.01907 0.01858 0.01942
As mentioned above, the Zener–Hollomon parameter Z is commonly used to summarize the effects of the deformation temperature (T) and the strain rate (ε• ) on hot or warm deformation behavior of metal materials. At relatively low deformation temperature, the relationship
of the Zener–Hollomon parameter Z and the peak stress σm can be expressed as [30–32]: Z ¼ B expðβσ m Þ
ð1Þ
where, B and β are constants. Using a mathematical method based on the least square method [32], with the data of σm at various deformation conditions and the two equations about the relationships of Z and T, ε•, σm, the deformation activation energy Q and the constants B and β were calculated for the used steels, as shown in Table 2. In comparison with that of Steel A, the deformation activation energy Q of Steel B was increased and that of Steel C was decreased, respectively, due to the increase in Mn content or C content. Meanwhile, the deformation activation energy Q of Steel D
Fig. 3. Relationship of the deformation activation energy Q of Steels A, C and E with the carbon content.
Fig. 4. Relationships between the peak stresses and the Zener–Hollomon parameters for the used steels during warm deformation of martensite.
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or Steel E was decreased slightly, i.e. about 1.5% or about 1.3% respectively, suggesting that the increase in Si content for Steel D or the increase in Al content for Steel E shows slight influence on the deformation activation energy Q of the used steels. Generally, the deformation activation energy Q can be regarded as an indicator of deformation difficulty degree in plasticity deformation [33], and is close to the lattice self-diffusion activation energy of pure metals [32]. For alloys, such as steel, Q is sensitive to even small alterations in the chemical composition of the material [20], owning to the solution strengthening and the solute drag effect [9,32,34]. In the present work, the deformation activation energy Q for each of the used steels, as shown in Table 2, is much larger than the self-diffusion activation energy of pure α-iron, i.e. 239 kJ/mol [9,32], which should be mainly attributed to the high Mn contents of the used steels. The increase in Mn content retarded the diffusion of Fe atoms, thus increased the deformation activation energy Q of the used steels [21]. Moreover, since the peak stress of the flow curves of the used steels during warm deformation of martensite corresponded to the decomposition of martensite into ferrite and carbides [7], the increase in C content was of benefit to the decomposition of martensite and thus led to the reduction of the deformation activation energy Q of Steel C. Although both of Si and Al have solute strengthening effect and the increase in Si content or Al content could retard the decomposition of martensite into ferrite and carbides, the increase in Si content or Al content showed slight influence on the deformation activation energy Q of the used
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steels. In special, the values of the deformation activation energy Q of Steel A, Steel C and Steel E were in a linear relationship with C contents, as shown in Fig. 3, suggesting that the slight decrease in the deformation activation energy Q of Steel E may be mainly attributed to the increase in C content. Such phenomena were quite different with that in the investigations about the influences of Si content and Al content on the deformation activation energy Q of austenite [20,33,35–36]. The work on the influence of Si content or Al content on the warm deformation behavior of martensite is insufficient, further systemic investigations is required, as the work of Medina et al. on the deformation activation energy Q of hot deformation of austenite in steels [20]. As a whole, the constitutive equations for warm deformation of the used steels could be expressed as follows: • 336600 Z ¼ ε exp ¼ 1:554 1014 expð0:02163σ m ÞðSteel AÞ 8:3145T
ð2Þ
• 419300 ¼ 4:431 1018 expð0:02264σ m Þ ðSteel BÞ Z ¼ ε exp 8:3145T
ð3Þ
• 311000 ¼ 9:147 1012 expð0:01907σ m Þ ðSteel CÞ Z ¼ ε exp 8:3145T
ð4Þ
• 331500 ¼ 1:668 1014 expð0:01858σ m ÞðSteel DÞ Z ¼ ε exp 8:3145T
ð5Þ
Fig. 5. Microstructures of steels deformed to 0.10 at 600 °C at 0.01 s−1 and air cooled: (a) steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
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Fig. 6. Microstructures of steels deformed to 0.22 at 600 °C at 0.01 s−1 and air cooled: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
•
Z ¼ ε exp
332000 8:3145T
¼ 1:201 1014 expð0:01942σ m Þ ðSteel EÞ
ð6Þ
As shown in Fig. 4, the Z–σm curves obtained from the Eqs. (2)–(6) correspond well with the experimental data of the used steel. 3.2. Microstructure evolution of the used steels during warm deformation Figs. 5–7 demonstrate the microstructures of the used steels by warm deformation of martensite at 600 °C at 0.01 s−1 to a strain of 0.10, 0.22 or 0.69 and air cooled to RT, the corresponding evolutions of the volume fractions of DRX ferrite grains, M/A islands and RA are presented in Fig. 8. Fig. 9 shows EBSD maps of Steel A deformed at 600 °C at 0.01 s−1 to 0.22 and 0.69. Low angle grain boundary (LAGB) with misorientation between 2° to 15° is presented by thin green line, and high angle grain boundary (HAGB) with misorientation more than 15° is shown by thick black line. Fig. 10 shows EBSD maps of Steels B–E deformed at 600 °C at 0.01 s−1 to 0.69, and the HAGB fraction of the microstructures of Steels A–E deformed to 0.69 is shown in Fig. 11. The microstructural evolutions of steels A–E during warm deformation of martensite were similar, i.e. the decomposition of martensite into ferrite and carbides took place firstly, DRX of ferrite and the reverse transformation of austenite occurred subsequently [7]. At the strain of 0.10 which was just beyond the peak strain of each of the used steels, the decomposition of martensite was the main process,
as shown in Fig. 5. In comparison with that of Steel A, the increase in C content (Steel C) promoted the decomposition of martensite and resulted in more carbide particles while the increase in Mn content (Steel B), Si content (Steel D) or Al content (Steel E) retarded the decomposition of martensite and resulted in less carbide particles. The effect of Mn should be attributed to the solute drag effect on the diffusion of Fe atoms and the effect of Si or Al should be due to their retarding effect on the formation of carbide particles. With the strain up to 0.22, DRX of ferrite occurred and small DRX grains surrounded by HAGB were easily formed around carbide particles [37], as shown in Figs. 6 and 9a. With further straining to 0.69, the development of DRX of ferrite resulted in the increase in the volume fractions and sizes of DRX grains, as well as the fraction of HAGB, as shown in Figs. 7, 8a and 9b. Meanwhile, the reverse transformation of austenite resulted in the formation of M/A islands in the final aircooled microstructures, as shown in Figs. 7 and 8b. In comparison with that of Steel A, DRX of ferrite was promoted in Steel C or Steel E and was retarded in Steel B or Steel D, as illustrated by Figs. 8a and 10. The promoting effect of the increase in C content should be attributed to the marked increase in the amount of carbide particles which can simulate the formation of DRX nuclei [37], and the retarding effect of the increase in Mn content or Si content should be attributed to the solid solution drag effect of Mn atoms or Si atoms. The relative higher C content of Steel E may be the main reason for the well development of DRX in comparison with that of Steel A, since the increase in Al
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Fig. 7. Microstructures of steels deformed to 0.69 at 600 °C at 0.01 s−1 and air cooled: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
content may also retard the development of DRX of ferrite due to the solid solution drag effect of Al atoms. However, the effect of the increase in Al content on DRX of ferrite requires further investigation. As a whole, due to the high Mn contents of 5.0–8.0 wt.% of the used steels, the development of DRX of ferrite during warm deformation was sluggish in comparison with that in common low-carbon steels [21]. Meanwhile, the reverse transformation of austenite was promoted by the increase in Mn content (Steel B) and retarded by the increase in C content (Steel C), Si content (Steel D) or Al content (Steel E), as illustrated by Fig. 8b. As shown in Table 1, the increase in Mn content leads to the decrease in the equilibrium transformation temperature A1, resulting in the higher chemical driving force for the reverse transformation of austenite during warm deformation of Steel B at 600 °C, in comparison with that of Steel A. On the contrary, the increase in C content, Si content or Al content leads to the increase in the equilibrium transformation temperature A1, which is higher than 600 °C, thus there was no chemical driving force for the reverse transformation of austenite during warm deformation of Steel C, Steel D or Steel E at 600 °C. That is, the reverse transformation of austenite during warm deformation of Steel C, Steel D or Steel E at 600 °C should be mainly attributed to the introduction of the mechanical driving force, i.e. the deformation storage energy. In comparison with that of Steel A, Steel C, Steel D or Steel E, the higher chemical driving force and higher deformation storage energy introduced by the higher flow stress of Steel B during warm deformation promoted the reverse transformation of austenite markedly, resulting in the much higher fraction of M/A islands in the
microstructure of Steel B deformed at 600 °C at 0.01 s−1 to 0.69. In addition, the higher Mn content of Steel B increased the thermal stability of the reversed austenite and made some reversed austenite having enough thermal stability to be retained at RT, as shown in Fig. 8c. The influence of the increase in Mn content, C content, Si content or Al content on DRX of ferrite and the reverse transformation of austenite during warm deformation of the used steels at each of the used deformation conditions were similar to that at 600 °C at 0.01 s−1. Fig. 12 summarizes the influences of the deformation temperature and the strain rate on the volume fractions of M/A islands in the air-cooled microstructures of the used steels deformed to 0.69 at different conditions. With the increase in the deformation temperature, the volume fraction of M/A islands in the air-cooled microstructure of each of the used steels was increased markedly, owning to the increased chemical driving force for the reverse transformation of austenite and the increased diffusion coefficients of C atoms, Mn atoms and Fe atoms etc. As mentioned above, the introduction of the mechanical driving force, i.e. the deformation storage energy, by warm deformation is of benefit to the reverse transformation of austenite, resulting in the formation of the reversed austenite during warm deformation of the used steels at the deformation temperature below the equilibrium transformation temperature A1, such as 550 °C or 600 °C. However, with the increase in the strain rate at a certain deformation temperature, which led to the increase in the deformation storage energy, the volume fraction of M/A islands in the air-cooled microstructure of each of the used steels was decreased markedly. Such phenomenon should be attributed to the shortening of
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Fig. 8. Volume fractions of DRX grains (a), M/A islands (b) and RA (c) in the microstructures of the used steels deformed to different strains at 600 °C at 0.01 s−1 and air cooled.
the deformation time for a certain strain by increasing the strain rate, resulting in the insufficient diffusions of C atoms, Mn atoms and Fe atoms. That means, except the thermodynamic conditions, the kinetic conditions is also an important factor which influences the reverse transformation of austenite during warm deformation of the used steels.
4. Conclusions Warm deformation behaviors of high-Mn TRIP steels with martensite as the initial structure were investigated using a Gleeble 1500 hot simulator within the temperature range of 550–650 °C and the strain rate range of 0.001–0.1 s− 1, and the influences of the
Fig. 9. EBSD maps of Steel A deformed at 600 °C at 0.01 s−1 to 0.22 (a) and 0.69 (b) and air cooled (Misorientation is between 2° to 15° for thin green line and more than 15° for thick black line, similarly hereinafter.).
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Fig. 10. EBSD maps of Steels B–E deformed at 600 °C at 0.01 s−1 to 0.69 and air cooled: (a): Steel B; (b) Steel C; (c) Steel D; (d) Steel E.
increase in C content, Mn content, Si content or Al content on the flow stress and the microstructural evolution of high-Mn TRIP steel during warm deformation of martensite were analyzed. The main results can be summarized as follows: 1. In comparison with 0.08C–4.6Mn steel, the increase in Mn content, C content, Si content or Al content resulted in the increase in the flow stress of the used steels, owning to the solid solute strengthening. The increase in Mn content resulted in the increase in the
deformation activation energy Q whereas the increase in C content resulted in the decrease in the value of Q. The increase in Si content or Al content showed slight influence on the deformation activation energy Q of the used steels. 2. In comparison with 0.08C–4.6Mn steel, the increase in C content promoted the decomposition of martensite and resulted in more carbide particles while the increase in Mn content, Si content or Al content retarded the decomposition of martensite and resulted in less carbide particles. DRX of ferrite was promoted by the increase
Fig. 11. Fraction of high angle grain boundaries of Steels A–E deformed at 600 °C at 0.01 s−1 to 0.69 and air cooled.
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Fig. 12. Volume fractions of M/A islands in the air-cooled microstructures of the used steels deformed to 0.69 at different conditions: (a) 0.001 s−1; (b) 0.01 s−1; (c) 0.1 s−1.
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Influences of alloying elements on warm deformation behavior of high-Mn TRIP steel with martensitic structure Zhikai Guo, Longfei Li ⁎ State Key Laboratory for Advanced Metals and Materials, University of Science & Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 1 August 2015 Received in revised form 15 September 2015 Accepted 3 October 2015 Available online 9 October 2015 Keywords: Warm deformation behavior High-Mn TRIP steel Deformation activation energy DRX of ferrite Alloying elements
a b s t r a c t Warm deformation behaviors of high-Mn TRIP steels with martensitic structure were investigated using a Gleeble 1500 hot simulator within the temperature range of 550–650 °C and the strain rate range of 0.001–0.1 s−1. The influences of common alloying elements, i.e. C, Mn, Si and Al, on the flow stress and the microstructural evolution during warm deformation were analyzed. The results showed that the flow stress was increased with the increase in C content, Mn content, Si content or Al content at the same condition, and the deformation activation energy Q was increased with the increase in Mn content and decreased by the increase in C content. The increase in Si content or Al content showed slight influence on the deformation activation energy Q. During warm deformation the decomposition of martensite into ferrite and carbides occurred firstly, and then dynamic recrystallization (DRX) of ferrite and the reverse transformation of austenite took place. DRX of ferrite was retarded by the increase in Mn content or Si content and accelerated by the increase in C content. Meanwhile, the reverse transformation of austenite was promoted by the increase in Mn content and retarded by the increase in C content, Si content or Al content. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction High-Mn TRIP steels, having a microstructure which consists of ultrafine-grained ferrite matrix and a large amount of retained austenite (RA) with suitable stability, demonstrate a well balance of strength and ductility [1]. For example, Furukawa et al. reported that a 0.1C–5.0Mn steel with about 30 vol.% RA, obtained by annealing at 600 °C for 3 h after hot rolling of austenite, demonstrated the tensile strength and uniform elongation up to about 1000 MPa and 24%, respectively [2]. The researches of Gibbs et al. showed that a tensile strength of 876 MPa and a total elongation of 42% can be achieved in a 0.1C–7.0Mn steel by annealing at 600 °C for 1 week after cold rolling of martensite, with the product of tensile strength and total elongation up to 37 GPa% [3]. Thus, as one potential contender of the third generation of advanced high strength steel, high-Mn TRIP steels with 5.0–8.0 wt.% Mn have received a lot of attention [1–8]. The common production of high-Mn TRIP steels requires long-time intercritical annealing in the austenite plus ferrite twophase region with the duration of about several hours to several days, directly after hot rolling of austenite or after hot rolling of austenite plus cold rolling of martensite [2–3,5–6]. Such longtime intercritical annealing is unfavorable to the mass production of high-Mn TRIP steels. In the previous work [7], a new thermomechanical process based on warm deformation of martensite ⁎ Corresponding author. E-mail address: [email protected] (L. Li).
http://dx.doi.org/10.1016/j.matdes.2015.10.010 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
and short-time intercritical annealing was developed to product high-Mn TRIP steels with less cost. The decomposition of martensite into ferrite and carbides, dynamic recrystallization (DRX) of ferrite and the reverse transformation of austenite were accelerated by warm deformation, thus leading to the shorter time for subsequent annealing in the intercritical region, with the length of about one hour. As a promising process to obtain ultrafine ferrite grains through DRX of ferrite, warm deformation behavior of steels has been well investigated in IF steels [9], low-carbon steels [10–12], medium-carbon steels [13–15] and high-carbon steels [16–17], with the initial microstructure of single-phase ferrite grains, ferrite grains plus pearlite colonies, pearlite colonies or lath martensite structure. Among them, warm deformation of lath martensite structure has a noticeable advantage to obtain ultrafine ferrite grains with the reduced critical strain, owning to the high density of dislocations, the supersaturated solute C atoms and the ultrafine laths [13,18]. Commonly, the kinetics of DRX of ferrite and the average size of DRX grains are influenced definitely by the deformation temperature and the strain rate. The Zener–Hollomon parameter Z is commonly used to summarize the effects of the deformation temperature and the strain rate on DRX, which can be expressed as Z ¼ ε • expðQ=RTÞ , where, ε• , Q, R and T refer to strain rate (s− 1), deformation activation energy (J/mol), gas constant (8.3145 J/mol/K) and absolute temperature (K), respectively [19]. It is clearly that the value of Z has a direct relationship with the deformation activation energy Q, which is well known to be influenced by the element contents of the matrix, such as C and Mn contents in austenite grains or ferrite grains [20–21]. For example, the increase in Mn content from 0.48% to
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Table 1 Chemical compositions (wt.%) and equilibrium transformation temperatures (°C) of the used steels. Steel
C, %
Mn, %
Si, %
Al, %
Fe, %
A1, °C
A3, °C
A B C D E
0.083 0.089 0.18 0.081 0.11
4.62 7.60 4.60 4.48 4.89
0.021 0.022 – 1.52 –
– – – – 1.01
Bal. Bal. Bal. Bal. Bal.
592 563 635 622 644
736 681 723 771 799
1.49% resulted in the increase in the value of Q from 260.3 kJ/mol to 450.0 kJ/mol for low-carbon steels with C content of about 0.06% during warm deformation in the ferrite phase region, retarding the occurrence of DRX of ferrite and refining the average size of DRX ferrite grains [21]. The addition of Al with the content of 0.95% resulted in the increase in the value of Q from 333.6 kJ/mol to 387.6 kJ/mol for high-carbon steels with C content of about 1.0% during warm deformation of martensite, impeding the decomposition of martensite and refining the average sizes of ferrite grains and cementite particles [16]. As the common elements in high-Mn TRIP steels, the increase in Mn content or C content can result in the improvement in the mechanical properties of high-Mn TRIP steels, due to the increase in the stability of RA [1,2–3,7]. Meanwhile, the addition of Si and Al can delay the precipitation of carbides and thus increase the stability of RA in the conventional TRIP steels [22]. Recently, ferrite–austenite duplex lightweight steel, such as 0.3C–8.5Mn–5.6Al steel, is developed to further reduce vehicle weight in order to reduce exhaust emissions and improve fuel efficiency, which shows the operation of both TRIP and TWIP mechanisms during tensile straining and thus demonstrates well balance of strength and elongation [23–26]. In this work, the influences of the common element, i.e. C, Mn, Si and Al, on the warm deformation behavior of high-Mn TRIP steel were investigated by hot uniaxial compression tests, in order to optimizing the aforementioned thermo-mechanical process based on warm deformation and short-time intercritical annealing with deep understanding of the warm deformation behavior of martensite in high-Mn TRIP steel.
2. Experimental Table 1 shows the chemical compositions (wt.%) of the used steels with the corresponding equilibrium transformation temperatures A1 and A3. The increase in C content or Mn content leads to the narrowing of the intercritical region, while the increase in Si content or Al content leads to the broadening of the intercritical region. The casting blanks were reheated to 1200 °C and held for 2 h, then hot-forged at 1150– 900 °C followed by air-cooling to room temperature (RT). Due to the high Mn contents, martensitic structure can be obtained by air-cooling for the used steels [7]. Cylindrical specimens of 6 mm in diameter and 15 mm in length were machined from the blanks. After austenization at 800 °C for 30 min and air cooling to RT, warm compression tests were performed using a Gleeble 1500 hot simulator. The temperature of the sample during warm deformation was automatically controlled by computer using the feedback signals of a pair of thermocouples, which were welded on the specimen during deformation. The specimens were reheated to 550, 600 or 650 °C at 20 °C/s, isothermally held for 1 min, and then deformed to a true strain of 0.10, 0.22 and 0.69 at 0.001, 0.01 or 0.1 s−1, finally air-cooled to RT. The microstructure observation was carried out using a ZEISS SUPRA55 field-emission scanning electronic microscope (FE-SEM) equipped with an electron backscattered diffraction (EBSD) attachment on the sections parallel to the compression direction after etching with 4 vol.% Nital. The volume fraction of ferrite, and the size of ferrite and M/A islands were measured using Image-Pro Plus image-analysis software with scanning electron micrographs. The volume fraction of RA was determined by a TTRIII X-ray diffractometer using Cu Kα radiation [27].
3. Results and discussions 3.1. Flow curves of the used steels during warm deformation Fig. 1 shows the flow curves of the used steels deformed at 600 °C at 0.01 s−1 to a strain of 0.69. The flow curves of the used steels had the similar shape, i.e. the flow stress increased rapidly to a peak and then
Fig. 1. Flow curves of the used steels during warm deformation of martensite at 600 °C at 0.01 s−1.
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Fig. 2. Peak stresses of the used steels during warm deformation of martensite under different conditions: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
decreased continuously with straining, suggesting the appearance of dynamic softening. Fig. 2 summarizes the effect of the deformation temperature or the strain rate on the peak stress of the used steels deformed at various conditions, which is clear that the peak stress was decreased
with the increase in the deformation temperature or the decrease in the strain rate for each steel. Meanwhile, owning to the solid solute strengthening effect of alloying elements [28–29], the peak stress of each of the steels B–E was higher than that of the steel A at the same condition.
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Table 2 Values of deformation activation energy Q and constants B and β for the used steels during warm deformation of martensite. Steel
Q, kJ/mol
B, s−1
β
A B C D E
336.6 419.3 311.0 331.5 332.0
1.554 × 1014 4.431 × 1018 9.147 × 1012 1.668 × 1014 1.201 × 1014
0.02163 0.02264 0.01907 0.01858 0.01942
As mentioned above, the Zener–Hollomon parameter Z is commonly used to summarize the effects of the deformation temperature (T) and the strain rate (ε• ) on hot or warm deformation behavior of metal materials. At relatively low deformation temperature, the relationship
of the Zener–Hollomon parameter Z and the peak stress σm can be expressed as [30–32]: Z ¼ B expðβσ m Þ
ð1Þ
where, B and β are constants. Using a mathematical method based on the least square method [32], with the data of σm at various deformation conditions and the two equations about the relationships of Z and T, ε•, σm, the deformation activation energy Q and the constants B and β were calculated for the used steels, as shown in Table 2. In comparison with that of Steel A, the deformation activation energy Q of Steel B was increased and that of Steel C was decreased, respectively, due to the increase in Mn content or C content. Meanwhile, the deformation activation energy Q of Steel D
Fig. 3. Relationship of the deformation activation energy Q of Steels A, C and E with the carbon content.
Fig. 4. Relationships between the peak stresses and the Zener–Hollomon parameters for the used steels during warm deformation of martensite.
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or Steel E was decreased slightly, i.e. about 1.5% or about 1.3% respectively, suggesting that the increase in Si content for Steel D or the increase in Al content for Steel E shows slight influence on the deformation activation energy Q of the used steels. Generally, the deformation activation energy Q can be regarded as an indicator of deformation difficulty degree in plasticity deformation [33], and is close to the lattice self-diffusion activation energy of pure metals [32]. For alloys, such as steel, Q is sensitive to even small alterations in the chemical composition of the material [20], owning to the solution strengthening and the solute drag effect [9,32,34]. In the present work, the deformation activation energy Q for each of the used steels, as shown in Table 2, is much larger than the self-diffusion activation energy of pure α-iron, i.e. 239 kJ/mol [9,32], which should be mainly attributed to the high Mn contents of the used steels. The increase in Mn content retarded the diffusion of Fe atoms, thus increased the deformation activation energy Q of the used steels [21]. Moreover, since the peak stress of the flow curves of the used steels during warm deformation of martensite corresponded to the decomposition of martensite into ferrite and carbides [7], the increase in C content was of benefit to the decomposition of martensite and thus led to the reduction of the deformation activation energy Q of Steel C. Although both of Si and Al have solute strengthening effect and the increase in Si content or Al content could retard the decomposition of martensite into ferrite and carbides, the increase in Si content or Al content showed slight influence on the deformation activation energy Q of the used
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steels. In special, the values of the deformation activation energy Q of Steel A, Steel C and Steel E were in a linear relationship with C contents, as shown in Fig. 3, suggesting that the slight decrease in the deformation activation energy Q of Steel E may be mainly attributed to the increase in C content. Such phenomena were quite different with that in the investigations about the influences of Si content and Al content on the deformation activation energy Q of austenite [20,33,35–36]. The work on the influence of Si content or Al content on the warm deformation behavior of martensite is insufficient, further systemic investigations is required, as the work of Medina et al. on the deformation activation energy Q of hot deformation of austenite in steels [20]. As a whole, the constitutive equations for warm deformation of the used steels could be expressed as follows: • 336600 Z ¼ ε exp ¼ 1:554 1014 expð0:02163σ m ÞðSteel AÞ 8:3145T
ð2Þ
• 419300 ¼ 4:431 1018 expð0:02264σ m Þ ðSteel BÞ Z ¼ ε exp 8:3145T
ð3Þ
• 311000 ¼ 9:147 1012 expð0:01907σ m Þ ðSteel CÞ Z ¼ ε exp 8:3145T
ð4Þ
• 331500 ¼ 1:668 1014 expð0:01858σ m ÞðSteel DÞ Z ¼ ε exp 8:3145T
ð5Þ
Fig. 5. Microstructures of steels deformed to 0.10 at 600 °C at 0.01 s−1 and air cooled: (a) steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
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Fig. 6. Microstructures of steels deformed to 0.22 at 600 °C at 0.01 s−1 and air cooled: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
•
Z ¼ ε exp
332000 8:3145T
¼ 1:201 1014 expð0:01942σ m Þ ðSteel EÞ
ð6Þ
As shown in Fig. 4, the Z–σm curves obtained from the Eqs. (2)–(6) correspond well with the experimental data of the used steel. 3.2. Microstructure evolution of the used steels during warm deformation Figs. 5–7 demonstrate the microstructures of the used steels by warm deformation of martensite at 600 °C at 0.01 s−1 to a strain of 0.10, 0.22 or 0.69 and air cooled to RT, the corresponding evolutions of the volume fractions of DRX ferrite grains, M/A islands and RA are presented in Fig. 8. Fig. 9 shows EBSD maps of Steel A deformed at 600 °C at 0.01 s−1 to 0.22 and 0.69. Low angle grain boundary (LAGB) with misorientation between 2° to 15° is presented by thin green line, and high angle grain boundary (HAGB) with misorientation more than 15° is shown by thick black line. Fig. 10 shows EBSD maps of Steels B–E deformed at 600 °C at 0.01 s−1 to 0.69, and the HAGB fraction of the microstructures of Steels A–E deformed to 0.69 is shown in Fig. 11. The microstructural evolutions of steels A–E during warm deformation of martensite were similar, i.e. the decomposition of martensite into ferrite and carbides took place firstly, DRX of ferrite and the reverse transformation of austenite occurred subsequently [7]. At the strain of 0.10 which was just beyond the peak strain of each of the used steels, the decomposition of martensite was the main process,
as shown in Fig. 5. In comparison with that of Steel A, the increase in C content (Steel C) promoted the decomposition of martensite and resulted in more carbide particles while the increase in Mn content (Steel B), Si content (Steel D) or Al content (Steel E) retarded the decomposition of martensite and resulted in less carbide particles. The effect of Mn should be attributed to the solute drag effect on the diffusion of Fe atoms and the effect of Si or Al should be due to their retarding effect on the formation of carbide particles. With the strain up to 0.22, DRX of ferrite occurred and small DRX grains surrounded by HAGB were easily formed around carbide particles [37], as shown in Figs. 6 and 9a. With further straining to 0.69, the development of DRX of ferrite resulted in the increase in the volume fractions and sizes of DRX grains, as well as the fraction of HAGB, as shown in Figs. 7, 8a and 9b. Meanwhile, the reverse transformation of austenite resulted in the formation of M/A islands in the final aircooled microstructures, as shown in Figs. 7 and 8b. In comparison with that of Steel A, DRX of ferrite was promoted in Steel C or Steel E and was retarded in Steel B or Steel D, as illustrated by Figs. 8a and 10. The promoting effect of the increase in C content should be attributed to the marked increase in the amount of carbide particles which can simulate the formation of DRX nuclei [37], and the retarding effect of the increase in Mn content or Si content should be attributed to the solid solution drag effect of Mn atoms or Si atoms. The relative higher C content of Steel E may be the main reason for the well development of DRX in comparison with that of Steel A, since the increase in Al
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Fig. 7. Microstructures of steels deformed to 0.69 at 600 °C at 0.01 s−1 and air cooled: (a) Steel A; (b) Steel B; (c) Steel C; (d) Steel D; (e) Steel E.
content may also retard the development of DRX of ferrite due to the solid solution drag effect of Al atoms. However, the effect of the increase in Al content on DRX of ferrite requires further investigation. As a whole, due to the high Mn contents of 5.0–8.0 wt.% of the used steels, the development of DRX of ferrite during warm deformation was sluggish in comparison with that in common low-carbon steels [21]. Meanwhile, the reverse transformation of austenite was promoted by the increase in Mn content (Steel B) and retarded by the increase in C content (Steel C), Si content (Steel D) or Al content (Steel E), as illustrated by Fig. 8b. As shown in Table 1, the increase in Mn content leads to the decrease in the equilibrium transformation temperature A1, resulting in the higher chemical driving force for the reverse transformation of austenite during warm deformation of Steel B at 600 °C, in comparison with that of Steel A. On the contrary, the increase in C content, Si content or Al content leads to the increase in the equilibrium transformation temperature A1, which is higher than 600 °C, thus there was no chemical driving force for the reverse transformation of austenite during warm deformation of Steel C, Steel D or Steel E at 600 °C. That is, the reverse transformation of austenite during warm deformation of Steel C, Steel D or Steel E at 600 °C should be mainly attributed to the introduction of the mechanical driving force, i.e. the deformation storage energy. In comparison with that of Steel A, Steel C, Steel D or Steel E, the higher chemical driving force and higher deformation storage energy introduced by the higher flow stress of Steel B during warm deformation promoted the reverse transformation of austenite markedly, resulting in the much higher fraction of M/A islands in the
microstructure of Steel B deformed at 600 °C at 0.01 s−1 to 0.69. In addition, the higher Mn content of Steel B increased the thermal stability of the reversed austenite and made some reversed austenite having enough thermal stability to be retained at RT, as shown in Fig. 8c. The influence of the increase in Mn content, C content, Si content or Al content on DRX of ferrite and the reverse transformation of austenite during warm deformation of the used steels at each of the used deformation conditions were similar to that at 600 °C at 0.01 s−1. Fig. 12 summarizes the influences of the deformation temperature and the strain rate on the volume fractions of M/A islands in the air-cooled microstructures of the used steels deformed to 0.69 at different conditions. With the increase in the deformation temperature, the volume fraction of M/A islands in the air-cooled microstructure of each of the used steels was increased markedly, owning to the increased chemical driving force for the reverse transformation of austenite and the increased diffusion coefficients of C atoms, Mn atoms and Fe atoms etc. As mentioned above, the introduction of the mechanical driving force, i.e. the deformation storage energy, by warm deformation is of benefit to the reverse transformation of austenite, resulting in the formation of the reversed austenite during warm deformation of the used steels at the deformation temperature below the equilibrium transformation temperature A1, such as 550 °C or 600 °C. However, with the increase in the strain rate at a certain deformation temperature, which led to the increase in the deformation storage energy, the volume fraction of M/A islands in the air-cooled microstructure of each of the used steels was decreased markedly. Such phenomenon should be attributed to the shortening of
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Fig. 8. Volume fractions of DRX grains (a), M/A islands (b) and RA (c) in the microstructures of the used steels deformed to different strains at 600 °C at 0.01 s−1 and air cooled.
the deformation time for a certain strain by increasing the strain rate, resulting in the insufficient diffusions of C atoms, Mn atoms and Fe atoms. That means, except the thermodynamic conditions, the kinetic conditions is also an important factor which influences the reverse transformation of austenite during warm deformation of the used steels.
4. Conclusions Warm deformation behaviors of high-Mn TRIP steels with martensite as the initial structure were investigated using a Gleeble 1500 hot simulator within the temperature range of 550–650 °C and the strain rate range of 0.001–0.1 s− 1, and the influences of the
Fig. 9. EBSD maps of Steel A deformed at 600 °C at 0.01 s−1 to 0.22 (a) and 0.69 (b) and air cooled (Misorientation is between 2° to 15° for thin green line and more than 15° for thick black line, similarly hereinafter.).
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Fig. 10. EBSD maps of Steels B–E deformed at 600 °C at 0.01 s−1 to 0.69 and air cooled: (a): Steel B; (b) Steel C; (c) Steel D; (d) Steel E.
increase in C content, Mn content, Si content or Al content on the flow stress and the microstructural evolution of high-Mn TRIP steel during warm deformation of martensite were analyzed. The main results can be summarized as follows: 1. In comparison with 0.08C–4.6Mn steel, the increase in Mn content, C content, Si content or Al content resulted in the increase in the flow stress of the used steels, owning to the solid solute strengthening. The increase in Mn content resulted in the increase in the
deformation activation energy Q whereas the increase in C content resulted in the decrease in the value of Q. The increase in Si content or Al content showed slight influence on the deformation activation energy Q of the used steels. 2. In comparison with 0.08C–4.6Mn steel, the increase in C content promoted the decomposition of martensite and resulted in more carbide particles while the increase in Mn content, Si content or Al content retarded the decomposition of martensite and resulted in less carbide particles. DRX of ferrite was promoted by the increase
Fig. 11. Fraction of high angle grain boundaries of Steels A–E deformed at 600 °C at 0.01 s−1 to 0.69 and air cooled.
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Fig. 12. Volume fractions of M/A islands in the air-cooled microstructures of the used steels deformed to 0.69 at different conditions: (a) 0.001 s−1; (b) 0.01 s−1; (c) 0.1 s−1.
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