- Email: [email protected]
Effects of auto-tempering on microstructure and mechanical properties in hot rolled plain C-Mn dual phase steels
Materials Science & Engineering A 665 (2016) 98–107
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effects of auto-tempering on microstructure and mechanical properties in hot rolled plain C-Mn dual phase steels Cheng-ning Li a, Guo Yuan a,n, Feng-qin Ji a, Dong-sheng Ren b, Guo-dong Wang a a State Key Laboratory of Rolling and Automation, Northeastern University, P.O. Box. 105, No. 11, Lane 3, Wenhua Road, Heping District, Shenyang 110819, Liaoning, China b Compact Strip Production Plant, Baotou Iron & Steel (Group) Co., Ltd., Baotou 014010, China
art ic l e i nf o
a b s t r a c t
Article history: Received 30 November 2015 Received in revised form 4 April 2016 Accepted 12 April 2016 Available online 13 April 2016
In this paper, the behavior of auto-tempering in hot rolled plain C-Mn dual phase steels was investigated. The steels with thickness of 11 mm were fabricated on the compact strip production line and coiled at 150 °C and 260 °C, respectively. The auto-tempering in the steel with coiling temperature of 260 °C caused the softening of ferrite and martensite and the reduction in hardness difference between these two phases, ultimately led to the decrease in tensile strength and strain hardening ability and the increase in post-necking elongation. The Charpy impact test results indicated that the toughness of autotempered steel was superior to the non-tempered steel. When the tensile test or impact test was conducted on the auto-tempered steel, both ferrite grains and tempered martensite islands were readily elongated as fibrous structure. The acquired fibrous structure could inhibit the formation of microvoids and propagation of cracks in tensile test, and delayed the tensile fracture. In addition, the fibrous structure could inhibit the propagation of main crack during impact test, whereas induced the branching cracks, leading to high impact energy. & 2016 Elsevier B.V. All rights reserved.
Keywords: Hot rolled dual phase steel Auto-tempering Fibrous structure Fracture Thermo mechanical processing
1. Introduction There are two main approaches to improve the strength of steels. The first one is the addition of alloy elements, especially microalloying elements; the second one is the improvement of manufacture processes to expose the potential properties of steels. The latter is more attractive due to the advantages of productive resource usage and low-cost. Dual phase (DP) treatment is an economical method to improve the mechanical properties of steels. The mechanical properties of DP steel, such as high tensile strength, continuous yielding behavior, low yield strength and high strain hardening rate and so on, are closely associated with the transformation and morphology of martensite [1–4]. Because tempering can change the morphology of martensite and other structure characteristics derived from martensite transformation, some of the mechanical properties of DP steel are sensitive to the tempering process. The microstructure evolution and changes in strength and plasticity of heat treatment DP steel during tempering have been fully investigated through the isothermal tempering treatment [5– 9]. Chang [5] summarized that the microstructure exhibited n
Corresponding author. E-mail address: [email protected] (G. Yuan).
http://dx.doi.org/10.1016/j.msea.2016.04.038 0921-5093/& 2016 Elsevier B.V. All rights reserved.
different characteristics under the four isothermal tempering regimes: below 225 °C, 225–350 °C, 350–535 °C and above 535 °C, respectively. Kuang [6] illustrated that the microstructure evolution of DP steel during tempering included three stages: the relief of residual stress, the precipitation of carbides in both ferrite and martensite, and the dissolution of carbides in the ferrite. At the higher tempering temperature (above 300 °C), both tensile strength and yield strength decreased, which was opposite to the elongation. When the tempering temperature lowered than 200 °C, the changes in mechanical properties highly depended on the tempering temperature and time. Except for tensile properties, few studies focused on the effects of tempering on impact toughness of DP steel. It is clear that the tempering of martensite readily occurs in steels with lean chemical compositions, particularly when the martensite forms near the martensite start temperature (Ms) [10,11]. Currently, the tempering of DP steels reported in previous studies mainly stems from the reheating offline, and few papers consider the tempering behavior occurred on hot rolling line. In fact, the tempering process after coiling can occur on the hot strip rolling line in the steels with thick gauge. Because the residual heat is difficult to transfer into surroundings, the tempering occurs in the steel coils, called “auto-tempering”. Therefore, it is expected that the auto-tempering would occur in the hot rolled plain C-Mn DP steel with thick gauge when the steel are coiled at relative high temperature.
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
99
This work studies the phenomenon of auto-tempering on CSP line in a hot rolled plain C-Mn DP steel with thickness of 11 mm. The microstructure, tensile properties and impact properties were investigated in the DP steels with the coiling temperature at 150 °C and 260 °C, respectively. The fracture behaviors during tensile test and impact test were discussed in details to provide fundamental understanding of the effects of auto-tempering on strength, plasticity and toughness.
2. Experimental procedure 2.1. Thermo-mechanical controlled processing The experimental steel was a plain C-Mn steel with a simple chemical composition of 0.06 C, 0.45 Si, and 1.55 Mn (wt%), without the addition of Ni, Cr and Mo (the measured amounts of Ni, Cr and Mo were 0.009%, 0.02% and 0.005%, respectively). The experiments were carried out on the CSP line of Baotou Steel in China. The schematic diagram and the corresponding temperature regime are given in Fig. 1. The steels were casted into casting slabs with gauge of 70 mm and then held for 12 min at 1180 °C in a tunnel furnace. The slabs were continuously hot-rolled into 11 mm thick strips with the finish rolling temperature of 860 °C. After hot rolling, the strips were immediately cooled to 670 °C by laminar cooling system with the cooling rate of 30 °C/s, followed by air cooling, then fast cooled to 150 °C and 260 °C respectively through ultra fast cooling system with the cooling rate of 120 °C/s, and eventually coiled into coils. After coiling, the coils were cooled in the air. The cooling rate of the coils in the air decreased with time, presented in Fig. 1b. The air cooling process of coil from 260 °C to 150 °C took about 4 h, while that from 150 °C to room temperature took about 35 h. For convenience, the steels strips with coiling temperatures of 150 °C and 260 °C were abbreviated as CT1 and CT2, respectively. 2.2. Microstructure observation and mechanical properties test The steel coils were cooled in the air for 48 h before being prepared into the specimens for metallography, tensile test and V-notch Charpy impact. The metallography specimens were mechanically polished using standard metallographic procedure, etched by 4% nital solution and observed using Leica-DMIRM optical microscope (OM) and ZEISS ULTRA 55 field emission scanning electron microscopy (SEM), respectively. Additionally, the 3 mm discs were punched and twin-jet electro polished in a mixture of 12.5% perchloric acid and 87.5% ethanol at 30 °C with a voltage of 32 V. They were observed in a FEI Tecnai G2 F20 field-emission transmission electron microscope (TEM) operated at 200 kV. Micro-hardness measurements were carried out by FM-700 Vickers hardness testing machine with a load of 10 g. The tensile samples with gauge length of 50 mm and diameter of 8 mm and the V-notch Charpy impact specimens with the size of 10 mm 10 mm 55 mm were prepared along the rolling direction. CMT5105-SANS computerized tensile testing machine was used to detect the tensile properties with a crosshead speed of 3 mm/min at room temperature. The Charpy impact test was implemented at 20 to 100 °C using Instron Dynatup 9200 series instrumented drop weight impact tester.
3. Results 3.1. Microstructural characteristics at different coiling temperatures Figs. 2 and, 3 show the microstructures of the steels observed
Fig. 1. (a) Schematic diagram of CSP line and (b) temperature regimes during thermo-mechanical controlled processing and after coiling.
by OM, SEM and TEM. As shown in Figs. 2a and 3a, the matrix of the steels CT1 and CT2 consisted of 85% polygonal ferrite with the average size of 5.5 mm. In the ferrite near the martensite/ ferrite interface, numerous dislocations were observed (Figs. 2b and 3b), which were induced by martensite transformation and affected the yielding behavior of dual phase steel [8]. The martensite morphology exhibited much difference between steels CT1 and CT2. Fig. 2c and d shows that the surface of martensite islands was smooth, suggesting no significant martensite tempering occurred at coiling temperature of 150 °C. However, presented in Fig. 3c and d, fine carbide particles were clearly observed in the martensite islands and the interface of ferrite and martensite exhibited less tetragonal, which inferred the martensite tempering at coiling temperature of 260 °C. It is mentioned that there were also few non-tempered martensites with small size in steel CT2 (Fig. 3e). The possible reason is that these small-sized martensites are highly enriched with carbon, forming at relative low temperature and difficult to be tempered. It indicated that auto-tempering occurred in the plain C-Mn dual phase steel when the coiling temperature was 260 °C. In the dual phase steel, the Ms was calculated as follows [12]:
Ms = 539 − 423C − 30.4Mn − 17.7Ni − 12.1Cr − 11.0Si − 7.0Mo
(1)
The carbon content of untransformed austenite was estimated by:
wγ = (w0 − fα wα )/(1 − fα )
(2)
where wγ, w0 and wα were the carbon contents of untransformed austenite, prior austenite and ferrite, respectively, and fα is the volume fraction of ferrite. In this work, the amounts of C, Si and Mn were 0.06%, 0.45% and 1.55% respectively, and that of Ni, Cr and Mo were ignored. The Ms of untransformed austenite estimated from (Eqs. (1) and 2) was 320 °C. Therefore, the martensite transformation accompanied with auto-tempering would occur at the coiling temperature of 260 °C, which was far more close to the Ms than 150 °C. It was reported that the carbides formation during tempering was suppressed by increasing the contents of Mn and Si [13–17]. However, the effect of Mn was weak when its content was less than 2% [15]. To suppress cementite formation, the Si content
100
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 2. Microstructures of steel CT1: (a) microstructure observed by OM, (b) dislocations observed by TEM, and (c) and (d) non-tempered martensite observed by SEM and TEM respectively.
Fig. 3. Microstructures of steel CT2: (a) microstructure observed by OM, (b) dislocations observed by TEM, (c) and (d) tempered martensite observed by SEM and TEM respectively, and (e) non-tempered martensite observed by SEM.
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
should be increased to 1.5% [16,17]. Therefore, the 1.55% Mn and 0.45% Si in the experimental steel may not play effectively role in suppressing carbide formation and increase the trends of autotempering. Auto-tempering led to the softening of both ferrite and martensite. When the coiling temperature increased from 150 °C to 260 °C, the micro-hardness of ferrite decreased from 148 HV to 139 HV, and that of martensite dropped from 346 HV to 269 HV. The softening degree of martensite was much larger than that of ferrite, causing the reduction of hardness difference between these two phases. This remarkable softening of martensite after autotempering is associated with the migration of carbon atoms, precipitation of particles and recovery of dislocation structure and so on [18].
101
Table 1 Tensile properties of the steels. Steel Yield Strength (MPa)
Tensile Strength (MPa)
Uniform Elongation (%)
Total Elongation (%)
Area Reduction (%)
CT1 CT2
598 551
16.5 13.0
35.5 36.5
71 86
334 394
3.2. Mechanical properties at different coiling temperatures 3.2.1. Tensile properties In the dual phase steel with low alloy, the auto-tempering is readily occurred at the coiling temperature close to Ms, as mentioned above, which will alter the mechanical properties of the steel. As shown in Fig. 4, steel CT1 exhibited a continuous yielding that was typical for dual phase steel. However, steel CT2 had a yield plateau with a pronounced increase in yield strength. It is known that the continuous yielding of dual phase steel was mainly attributed to the mobile dislocations produced by martensite transformation at low temperature [1,7,19,20], for example in Fig. 3a. Dislocations, one of structure defects in crystalline solid solutions, are disturbed regions with low thermodynamically stability, and attract solute segregation when diffusion is enable [21–23]. Therefore, when the coiling temperature increased to 260 °C, C and N atoms was considered to diffuse to the dislocations and pinned them. Many researches had also demonstrated that the interstitial atoms of C and N could segregate and pin the dislocations when the tempering temperature increased to 250 °C, leading to the rise in yield strength and the appearance of yield platform [7–9,19,23]. In this study, the yield plateau and the significantly increase in yield strength, to some extent, revealed that interstitial atoms of C and N segregated and pinned the dislocations during auto-tempering. As listed in Table, 1, steel CT1 was characterized by high tensile strength and strong strain hardening ability, which attributed to the high hardness of quenched martensite [24–26]. For steel CT2, the tensile strength decreased distinctly due to the softening of martensite. The reduction of hardness difference between ferrite
Fig. 4. Stress-strain curves of the steels during tensile test.
Fig. 5. True stress and strain hardening rate of the steels.
and martensite caused by auto-tempering resulted in the decrease of strain hardening ability [24,25]. When the coiling temperature increased from 150 °C to 260 °C, the uniform elongation decreased from 16.5% to 13%, which demonstrated that the auto-tempering stimulated the occurrence of necking. It is known that the necking appears when the strain hardening rate (ds/dε) is equal or less than the true stress (s) [27], and the corresponding strain is the uniform strain (εu), shown in Fig. 5. Thus, the increase of strain hardening rate or the decrease of true stress can postpone the occurrence of necking. After auto-tempering, both the strain hardening rate and the true stress reduced, but the degree of the former was more prominent. Therefore, the auto-tempering accelerated the occurrence of necking. Nevertheless, the stress-strain curve of steel CT2 had a distinct necking stage (Fig. 4), and the post-necking elongation and reduction of area reached 23.5% and 86%, respectively, higher than that of steel CT1. The data indicates that steel CT2 has the high fracture resistance. 3.2.2. Charpy impact energy and fractograph of the specimens Fig. 6 indicates that the toughness of steel CT2 was higher than that of steel CT1. The impact energy (vE) of steel CT1 was sensitive to the test temperature and decreased from 185 J to 10 J with the temperature decreasing from 20 °C to 100 °C. However, the toughness of steel CT2 under the test temperature range of 20 to 60 °C was such excellent (higher than 300 J) that the specimens cannot be completely separated. When the temperature decreased from 60 °C to 100 °C, the value of vE sharply dropped to 15 J from 320 J. The impact specimen macrofractographs of steel CT1 are given in Fig. 7. At 20 °C and 40 °C, the macrofractographs of steel CT1 contained the typical fibrous, radial, second fibrous and shear lip zones, marked by 1, 2, 3 and 4 respectively in Fig. 7a. The fibrous, second fibrous and shear lip zones of the steel CT1 at 20 °C consisted of dimples (Fig. 7d–f), reflecting the ductile fracture. The dimples in fibrous and second fibrous zones were stretched due to the shear stress. In the radial zone, the existence of cleavage facets, tearing ridges and cracks indicated the brittle fracture (Fig. 7g). As Shown in Fig. 7b and c, with decrease in test
102
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 6. Charpy V-notch impact properties: (a) impact energy versus test temperature and (b) the specimens after impact test.
Fig. 7. Impact specimen fractographs of steel CT1: (a)–(c) low magnification fractographs tested at fractographs of fibrous, second fibrous, shear lip and radial zone respectively.
temperature, the area of fibrous, second fibrous and shear lip zones decreased, replaced by the radial zone. At 60 °C, the radial zone occupied almost all the fracture surface of specimens. For steel CT2, shown in Fig. 8a, d and e, the fracture surfaces of unseparated specimen tested at 60 °C consisted of dimples, and the branching cracks were observed in the fracture surface. The macrofractures of steel CT2 at 80 °C (Fig. 8b) were familiar to that of steel CT1 at 20 °C and 40 °C. However, Fig. 8f–h shows that the dimples of steel CT2 were larger than that of steel CT1, indicating the superior toughness. At 100 °C, the fracture surface of specimens were mainly comprised of radial zone, which proven the low toughness at this temperature. Assuming that the test temperature corresponding to where the area of radial zone occupies 50% of the fracture surface is the ductile-to-brittle transition temperature (DBTT), the DBTT reduced to 80 °C from 40 °C when the steel was auto-tempered. The transition of ductile-to-brittle attributed to the competition of plastic deformation and brittle fracture at the tips of cracks or defects, and the DBTT was determined by the difference between tensile stress near the crack tips and the brittle fracture stress [28– 30]. According to Yoffee diagram [28,29], when the effective yield stress, at which significant deformation occurs, exceeds the brittle fracture stress, the brittle fracture occurs. Although the yield
20 °C,
40 °C and
60 °C, respectively, and (d)–(g) high magnification
strength of steel CT2 was higher than that of steel CT1 during tensile test, the stress of steel CT2 was lower than steel CT1 when the strain exceeds 0.01, shown in Fig. 4, which resulted from the softening of ferrite and martensite. In addition, the softening of steel CT2 caused by auto-tempering led to the extensive plastic deformation before fracture. It was clear that the extensive plastic deformation at the crack tip limited the local stress and inhibited the brittle fracture [31]. Therefore, the auto-tempering played an effective role in delaying the transition of ductile-to-brittle.
4. Discussion 4.1. Effects of auto-tempering on tensile fracture behavior Fig. 9 indicates that both the two steels fractured under ductile mechanism with cup-and-cone fracture. The steels suffered an obvious plastic deformation before fracture. However, the degree of section contraction of steel CT2 was much larger than that of steel CT1, indicating that auto-tempering provided superior ductility for DP steel. The fracture surfaces included two parts: central fiber zone and surrounding shear lip, marked in Fig. 9. During the tensile test, the microvoids virtually begin to initiate before
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
103
Fig. 8. Impact specimen fractographs of steel CT2: (a)–(c) low magnification fractographs tested at 60 °C, 80 °C and 100 °C, respectively, (d) and (e) local magnification fractographs of specimen tested at 60 °C, and (f)–(i) high magnification fractographs of fibrous, second fibrous, shear lip and radial zone respectively.
Fig. 9. Fracture surface of the tensile specimens for steel CT1 (a) and CT2 (b).
104
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 10. Subsurfaces of the main fracture surface in tensile specimens for steels CT1 (a) and CT2 (b).
necking. As the necking occurs, the stress concentration significantly increases and then induces more microvoids. The coalescence of microvoids leads to the microcracks, which grow perpendicular to the applied stress and evolve into the central fiber zone. As the necking further continues, the crack propagates along the 45° maximum shear stress direction and develops into the shear lip [28,32]. The shear lip of steel CT2 exhibited ring-like with “Λ” shape, which was caused by the severe deformation before fracture. As shown in the micrographs at high magnification, the fracture surfaces of the steels were composed of dimples. Nevertheless, the dimples in the fiber zone of steel CT2 were much larger and deeper than that of steel CT1, which may be the reason the post-necking elongation of steel CT2 was much larger than that of steel CT1. Fig. 10 shows the micrographs at subsurface under the main fracture surface of tensile specimens. The ferrite grains and martensite islands were elongated in the necking zone of steel CT1, but the elongated extent of ferrite grains was much higher than that of martensite islands due to the large hardness difference between ferrite and non-tempered martensite. The deformation incompatibility between ferrite and martensite resulted in high stress concentration, especially where the shape of martensite was tetragonal and sharp. Since the deformation between ferrite and martensite was highly incompatible (Fig. 10a) and the pure martensites were tetragonal (Fig. 2c) in steel CT1, lots of microvoids initiated around the tips of the elongated martensite islands (indicated by arrows in Fig. 10a), which led to the relatively small post-necking elongation. For steel CT2, both ferrite grains and tempered martensite islands were severely elongated as fibrous structure. The larger deformation and the better deformation compatibility between ferrite and tempered martensite were due to the softening and reduction of hardness difference between these two phases.
Fig. 10b indicates that the microvoids in steel CT2 tended to initiate near the non-tempered martensite islands and inclusions because the deformations of ferrite matrix and non-tempered martensite (or inclusion) were still incompatibility. Since most of martensite islands were auto-tempered and became less tetragonal (Fig. 3c) and the ferrite and martensite compatibly deformed as fibrous structure, the stress concentration was small compared to steel CT1. Therefore, the nucleation sites of microvoids in steel CT2 (indicated by arrows in Fig. 10b) was far less than that in steel CT1. The fewer nucleation sites of microvoids could explain why the dimples in fracture surface of steel CT2 were much larger and deeper than steel CT1. Additionally, the cracks in steel CT2 were expected to propagate more difficult to across the section of tensile specimen due to the fibrous structure normal to the propagation direction of cracks. Although both of two steels fractured with ductile mechanism, steel CT2 exhibited the higher ability to prevent the crack formation and propagation. 4.2. Effects of auto-tempering on impact fracture behavior To investigate the fracture process and the fracture behavior of the steels at different temperatures, the impact load (P), impact energy and deflection (u) were recorded by instrument during impact test. Fig. 11a and b show that the P-u curves of steel CT1 at 20 °C and 40 °C had the similar shape. At the beginning, the impact load increased rapidly to the peak value Pm. In this stage, the elastic deformation and plastic deformation occurred successively, and the absorbed energy corresponding to the Pm was crack initiate energy (vEi). The value of vEi at 20 °C and 40 °C were stable at 40–45 J in steel CT1. Then, with the increase in u, the value of P slowly decreased to Py, then dropped to Pa and slowly reduced to zero (P0), successively. The stages of PmPy, PyPa and PaP0 reflected the processes of ductile propagation, brittle fracture and
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
105
Fig. 11. Typical curves of impact load and impact energy versus deflection at different temperatures (black lines-steel CT1, red lines–steel CT2, solid lines - impact load and dashed lines - impact energy): (a) 20 °C, (b) 40 °C, (c) 60 °C (d) 80 °C and (e) 100 °C. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
ductile fracture, respectively. The total absorbed energy of these three stages was crack propagation energy (vEp). When the test temperature declined from 20 °C to 40 °C in steel CT1, the value of vEp decreased from 139 J to 95 J. This suggested that the lower temperature accelerated the crack propagation and eventually deteriorates the toughness. The microvoids and cracks underneath the fracture surface of specimen at 20 °C for steel CT1 are presented in Fig. 12a and b. Microvoids were observed underneath the fibrous zone of fracture surface, implying the ductile fracture. Underneath the radial zone, the cracks exhibited transgranular with the brittle mechanism. The characteristics of P-u curves and the observation of cracks underneath the fracture surface in steel CT1 demonstrated the mixed mechanism of ductile and brittle in the range of 20 to 40 °C. In the P-u curves at 60 to 100 °C, the ductile propagation and ductile fracture stages disappeared (Fig. 11c–e), which suggests that the fracture mechanism converted from mixture of ductile
and brittle fracture to complete brittle fracture. In Fig. 11a–c, the P-u curves of steel CT2 at 20 to 60 °C exhibited a distinct ductile propagation stage. It indicated that the auto-tempering played an effective role in suppressing the propagation of cracks. With increase in u, P decreased slowly from Pm, then rapidly declined to the first plateau stage, and the second plateau stage, and finally dropped to zero when the specimen was struck away from the specimen holder. This similar plateau-step of P-u curve was also observed in the low-carbon steel with severe elongated structure, multilayer composite steel and glass-fiber/ vinyl-ester composite [29,33,34]. In these materials, the branching cracks tended to propagate along the interfaces of elongated structure or fiber planes that were normal to the propagation direction of the main crack. In this case, the propagation of the main cracks was restrained, resulting in the plateau stage of P-u curve. The impact load dropped when the new cracks initiated. In the fractographs (Fig. 8a, d and e) and micrographs underneath the
106
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 12. Cracks underneath the fracture surface of impact specimens: (a) microvoids underneath fibrous zone of steel CT1 at 20 °C, (b) transgranular cracks underneath radial zone of steel CT1 at 20 °C, (c) and (d) fiber structure and branching cracks of steel CT2 at 20 °C and 60 °C, respectively, (e) microvoids underneath fibrous zone of steel CT2 at 80 °C, and (f) transgranular cracks underneath radial zone of steel CT2 at 80 °C.
fracture surface (Fig. 12b and c) of the impact specimens, the branching cracks were observed clearly in steel CT2 at 20 °C and 60 °C. The branching cracks were attributed to the acquired fibrous structure. Presented in Fig. 12b and c, the branching cracks propagated along the interface of fibrous structure. The shape of P-u curve of steel CT2 at 80 °C (Fig. 11d) was similar to that of steel CT1 at 20 to 40 °C, showing the similar fracture mechanism. In Fig. 12e and f, similar to micrographs of steel CT1 at 20 °C (Fig. 12a and b), the microvoids and the transgranular cracks were observed in fibrous and radial zones, respectively. At 100 °C, the P-u curve of steel CT2 exhibited the brittle type. Consequently, with the decrease in test temperature from 20 °C to 100 °C, the fracture mechanism of steel CT2 transferred from ductile fracture accompanied with branching crack to mixture of ductile and brittle fracture, and finally to brittle fracture.
5. Conclusions
1. The microstructure consisted of polygonal ferrite and non-
tempered martensite islands in the hot rolled plain C-Mn DP steel coiled at 150 °C. When the coiling temperature was 260 °C, the interface of ferrite and martensite became less tetragonal, and the fine carbide particles dispersed in the martensite islands were also observed. Both of them reflected the occurrence of auto-tempering, leading to the softening of ferrite and martensite and the reduction of hardness difference between these two phases. 2. The auto-tempering resulted in the appearance of yield platform with the increase in yield strength, while both the strain hardening rate and the tensile strength decreased. In particular, the auto-tempered DP steel exhibited a distinct necking stage, implying the superior ductility. During tensile test, the ferrite grains and tempered martensite islands were severely elongated as fibrous structure, which effectively postponed the tensile fracture. 3. Compared with the non-tempered steel, the auto-tempered DP steel had higher impact energy and lower DBTT. When the test temperature decreased from 20 to 60 °C, the fracture mechanism of non-tempering DP steel converted from mixture of ductile and brittle fracture to complete brittle fracture. At the test temperature above 60 °C, the impact specimens of auto-
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
tempered steel fractured by ductile mechanism accompanied with branching cracks. And the fracture mechanism transferred from mixture of ductile and brittle fracture to brittle fracture with the test temperature decreasing from 80 °C to 100 °C.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51234002) and Fundamental Research Funds for the Central Universities (N130407001).
References [1] M. Sarwar, R. Priestner, Influence of ferrite-martensite microstructural morphology on tensile properties of dual-phase steel, J. Mater. Sci. 31 (1996) 2091–2095. [2] A. Bag, K.K. Ray, E.S. Dwarakadasa, E.S. Dwarakadasa, Influence of martensite content and morphology on tensile and impact properties of high-martensite dual-phase steels, Metall. Mater. Trans. A 30 (1999) 1193–1202. [3] S. Oliver, T.B. Jones, G. Fourlaris, Dual phase versus TRIP strip steels: microstructural changes as a consequence of quasi-static and dynamic tensile testing, Mater. Charact. 58 (2007) 390–400. [4] M. Mazinani, W.J. Poole, Effect of martensite plasticity on the deformation behavior of a low-carbon dual-phase steel, Mater. Sci. Eng. A 38 (2007) 328–339. [5] P.H. Chang, Temper-aging of continuously annealed low carbon dual phase steel, Metall. Trans. A 15 (1984) 73–86. [6] C.F. Kuang, J. Li, S.G. Zhang, J. Wang, H.F. Liu, A.A. Volinsky, Effects of quenching and tempering on the microstructure and bake hardening behavior of ferrite and dual phase steels, Mater. Sci. Eng. A 613 (2014) 178–183. [7] S. Gunduz, Effect of chemical composition, martensite volume fraction and tempering on tensile behaviour of dual phase steels, Mater. Lett. 63 (2009) 2381–2383. [8] A.A. Sayed, S. Kheirandish, Affect of the tempering temperature on the microstructure and mechanical properties of dual phase steels, Mater. Sci. Eng. A 532 (2012) 21–25. [9] A. Kampa, S. Celottob, D.N. Hanlonb, Effects of tempering on the mechanical properties of high strength dual-phase steels, Mater. Sci. Eng. A 538 (2012) 35–41. [10] V.H.B. Hernandez, S.S. Nayak, Y. Zhou, Tempering of martensite in dual-phase steels and its effects on softening behavior, Metall. Mater. Trans. A 42 (2011) 3115–3129. [11] H. Matsuda, R. Mizuno, Y. Funakawa, K. Seto, S. Matsuoka, Y. Tanaka, Effects of auto-tempering behaviour of martensite on mechanical properties of ultra high strength steel sheets, J. Alloy. Compd. 557 (2013) 661–667. [12] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [13] S. Yan, X.H. Liu, W.J. Liu, H.F. Lan, H.Y. Wu, Comparison on mechanical properties and microstructure of a C-Mn-Si steel treated by quenching and partitioning (Q&P) and quenching and tempering (Q&T) processes, Mater. Sci. Eng. A 620 (2015) 58–66.
107
[14] J. Speer, D.K. Matlock, B.C. De Cooman, J.G. Schroth, Carbon partitioning into austenite after martensite transformation, Acta Mater. 51 (2003) 2611–2622. [15] J. Pacyna, A. Jedrzejewska-Strach, M. Strach, The effect of manganese and silicon on the kinetics of phase transformations during tempering - Continuous Heating Transformation (CHT) curves, J. Mater. Process. Technol. 64 (1997) 311–318. [16] D. De Knijf, E.P. Da Silva, Fojer, C. Fojer, R. Petrov, Study of heat treatment parameters and kinetics of quenching and partitioning cycles, Mater. Sci. Technol. 31 (2015) 817–828. [17] D. Hauserova, M. Duchek, J. Dlouhy, Z. Novy, Properties of advanced experimental CMnSiMo steel achieved by QP process, Procedia Eng. 10 (2011) 2961–2966. [18] C.E.I.C. Ohlund, E. Schlangen, S.E. Offerman, The kinetics of softening and microstructure evolution of martensite in Fe-C-Mn steel during tempering at 300 °C, Mater. Sci. Eng. A 560 (2013) 351–357. [19] S. Gündüz, Static strain ageing behaviour of dual phase steels, Mater. Sci. Eng. A 486 (2008) 63–71. [20] J. Kadkhodapour, S. Schmauder, D. Raabe, S. Ziaei-Rad, U. Weber, M. Calcagnotto, Experimental and numerical study on geometrically necessary dislocations and non-homogeneous mechanical properties of the ferrite phase in dual phase steels, Acta Mater. 59 (2011) 4387–4394. [21] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. I. Theoretical background, Acta Mater. 55 (2007) 5129–5138. [22] M. Kuzmina, M. Herbig, D. Ponge, S. Sandlobes, D. Raabe, Linear complexions: confined chemical and structural states at dislocations, Science 349 (2015) 1080–1083. [23] R.G. Davies, Early stages of yielding and strain aging of a vanadium-containing dual-phase steel, Metall. Trans. A 10 (1979) 1549–1554. [24] M.F. Ashby, Work hardening of dispersion-hardened crystals, Philos. Mag. 14 (1966) 1157–1178. [25] M.F. Ashby, The deformation of plastically non-homogeneous materials, Philos. Mag. 21 (1970) 399–424. [26] S. Ankem, H. Margolin, C.A. Greene, B.W. Neuberger, P.G. Oberson, Mechanical properties of alloys consisting of two ductile phases, Prog. Mater. Sci. 51 (2006) 632–709. [27] M. Saeglitz, G. Krauss, Deformation, fracture, and mechanical properties of low-temperature-tempered martensite in SAE 43xx steels, Metall. Mater. Trans. A 28 (1997) 377–387. [28] T. Inoue, F.X. Yin, Y. Kimura, K. Tsuzaki, S. Ochiai, Delamination effect on impact properties of ultrafine-grained low-carbon steel processed by warm caliber rolling, Metall. Mater. Trans. A 41 (2010) 341–355. [29] J.W. Morris, Stronger, tougher steels, Science 320 (2008) 1022–1023. [30] D.L. Bourell, Cleavage delamination in impact tested warm-rolled steel, Metall. Trans. A 14 (1983) 2487–2496. [31] J.W. Morris, Steels: For Low Temperature Applications, Pergamon Press,, Oxford, 1993. [32] Y.B. Xu, X.D. Tan, X.L. Yang, Z.P. Hu, D. Wu, G.D. Wang, Microstructure evolution and mechanical properties of a hot-rolled directly quenched and partitioned steel containing proeutectoid ferrite, Mater. Sci. Eng. A 607 (2014) 460–475. [33] J.F. Kalthoff, Characterization of the dynamic failure behaviour of a glass-fiber/ vinyl-ester at different temperatures by means of instrumented Charpy impact testing, Compos. Part B 35 (2004) 657–663. [34] M. Pozuelo, F. Carreno, O.A. Ruano, Delamination effect on the impact toughness of an ultrahigh carbon–mild steel laminate composite, Metall. Mater. Trans. A 41 (2010) 341–355.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effects of auto-tempering on microstructure and mechanical properties in hot rolled plain C-Mn dual phase steels Cheng-ning Li a, Guo Yuan a,n, Feng-qin Ji a, Dong-sheng Ren b, Guo-dong Wang a a State Key Laboratory of Rolling and Automation, Northeastern University, P.O. Box. 105, No. 11, Lane 3, Wenhua Road, Heping District, Shenyang 110819, Liaoning, China b Compact Strip Production Plant, Baotou Iron & Steel (Group) Co., Ltd., Baotou 014010, China
art ic l e i nf o
a b s t r a c t
Article history: Received 30 November 2015 Received in revised form 4 April 2016 Accepted 12 April 2016 Available online 13 April 2016
In this paper, the behavior of auto-tempering in hot rolled plain C-Mn dual phase steels was investigated. The steels with thickness of 11 mm were fabricated on the compact strip production line and coiled at 150 °C and 260 °C, respectively. The auto-tempering in the steel with coiling temperature of 260 °C caused the softening of ferrite and martensite and the reduction in hardness difference between these two phases, ultimately led to the decrease in tensile strength and strain hardening ability and the increase in post-necking elongation. The Charpy impact test results indicated that the toughness of autotempered steel was superior to the non-tempered steel. When the tensile test or impact test was conducted on the auto-tempered steel, both ferrite grains and tempered martensite islands were readily elongated as fibrous structure. The acquired fibrous structure could inhibit the formation of microvoids and propagation of cracks in tensile test, and delayed the tensile fracture. In addition, the fibrous structure could inhibit the propagation of main crack during impact test, whereas induced the branching cracks, leading to high impact energy. & 2016 Elsevier B.V. All rights reserved.
Keywords: Hot rolled dual phase steel Auto-tempering Fibrous structure Fracture Thermo mechanical processing
1. Introduction There are two main approaches to improve the strength of steels. The first one is the addition of alloy elements, especially microalloying elements; the second one is the improvement of manufacture processes to expose the potential properties of steels. The latter is more attractive due to the advantages of productive resource usage and low-cost. Dual phase (DP) treatment is an economical method to improve the mechanical properties of steels. The mechanical properties of DP steel, such as high tensile strength, continuous yielding behavior, low yield strength and high strain hardening rate and so on, are closely associated with the transformation and morphology of martensite [1–4]. Because tempering can change the morphology of martensite and other structure characteristics derived from martensite transformation, some of the mechanical properties of DP steel are sensitive to the tempering process. The microstructure evolution and changes in strength and plasticity of heat treatment DP steel during tempering have been fully investigated through the isothermal tempering treatment [5– 9]. Chang [5] summarized that the microstructure exhibited n
Corresponding author. E-mail address: [email protected] (G. Yuan).
http://dx.doi.org/10.1016/j.msea.2016.04.038 0921-5093/& 2016 Elsevier B.V. All rights reserved.
different characteristics under the four isothermal tempering regimes: below 225 °C, 225–350 °C, 350–535 °C and above 535 °C, respectively. Kuang [6] illustrated that the microstructure evolution of DP steel during tempering included three stages: the relief of residual stress, the precipitation of carbides in both ferrite and martensite, and the dissolution of carbides in the ferrite. At the higher tempering temperature (above 300 °C), both tensile strength and yield strength decreased, which was opposite to the elongation. When the tempering temperature lowered than 200 °C, the changes in mechanical properties highly depended on the tempering temperature and time. Except for tensile properties, few studies focused on the effects of tempering on impact toughness of DP steel. It is clear that the tempering of martensite readily occurs in steels with lean chemical compositions, particularly when the martensite forms near the martensite start temperature (Ms) [10,11]. Currently, the tempering of DP steels reported in previous studies mainly stems from the reheating offline, and few papers consider the tempering behavior occurred on hot rolling line. In fact, the tempering process after coiling can occur on the hot strip rolling line in the steels with thick gauge. Because the residual heat is difficult to transfer into surroundings, the tempering occurs in the steel coils, called “auto-tempering”. Therefore, it is expected that the auto-tempering would occur in the hot rolled plain C-Mn DP steel with thick gauge when the steel are coiled at relative high temperature.
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
99
This work studies the phenomenon of auto-tempering on CSP line in a hot rolled plain C-Mn DP steel with thickness of 11 mm. The microstructure, tensile properties and impact properties were investigated in the DP steels with the coiling temperature at 150 °C and 260 °C, respectively. The fracture behaviors during tensile test and impact test were discussed in details to provide fundamental understanding of the effects of auto-tempering on strength, plasticity and toughness.
2. Experimental procedure 2.1. Thermo-mechanical controlled processing The experimental steel was a plain C-Mn steel with a simple chemical composition of 0.06 C, 0.45 Si, and 1.55 Mn (wt%), without the addition of Ni, Cr and Mo (the measured amounts of Ni, Cr and Mo were 0.009%, 0.02% and 0.005%, respectively). The experiments were carried out on the CSP line of Baotou Steel in China. The schematic diagram and the corresponding temperature regime are given in Fig. 1. The steels were casted into casting slabs with gauge of 70 mm and then held for 12 min at 1180 °C in a tunnel furnace. The slabs were continuously hot-rolled into 11 mm thick strips with the finish rolling temperature of 860 °C. After hot rolling, the strips were immediately cooled to 670 °C by laminar cooling system with the cooling rate of 30 °C/s, followed by air cooling, then fast cooled to 150 °C and 260 °C respectively through ultra fast cooling system with the cooling rate of 120 °C/s, and eventually coiled into coils. After coiling, the coils were cooled in the air. The cooling rate of the coils in the air decreased with time, presented in Fig. 1b. The air cooling process of coil from 260 °C to 150 °C took about 4 h, while that from 150 °C to room temperature took about 35 h. For convenience, the steels strips with coiling temperatures of 150 °C and 260 °C were abbreviated as CT1 and CT2, respectively. 2.2. Microstructure observation and mechanical properties test The steel coils were cooled in the air for 48 h before being prepared into the specimens for metallography, tensile test and V-notch Charpy impact. The metallography specimens were mechanically polished using standard metallographic procedure, etched by 4% nital solution and observed using Leica-DMIRM optical microscope (OM) and ZEISS ULTRA 55 field emission scanning electron microscopy (SEM), respectively. Additionally, the 3 mm discs were punched and twin-jet electro polished in a mixture of 12.5% perchloric acid and 87.5% ethanol at 30 °C with a voltage of 32 V. They were observed in a FEI Tecnai G2 F20 field-emission transmission electron microscope (TEM) operated at 200 kV. Micro-hardness measurements were carried out by FM-700 Vickers hardness testing machine with a load of 10 g. The tensile samples with gauge length of 50 mm and diameter of 8 mm and the V-notch Charpy impact specimens with the size of 10 mm 10 mm 55 mm were prepared along the rolling direction. CMT5105-SANS computerized tensile testing machine was used to detect the tensile properties with a crosshead speed of 3 mm/min at room temperature. The Charpy impact test was implemented at 20 to 100 °C using Instron Dynatup 9200 series instrumented drop weight impact tester.
3. Results 3.1. Microstructural characteristics at different coiling temperatures Figs. 2 and, 3 show the microstructures of the steels observed
Fig. 1. (a) Schematic diagram of CSP line and (b) temperature regimes during thermo-mechanical controlled processing and after coiling.
by OM, SEM and TEM. As shown in Figs. 2a and 3a, the matrix of the steels CT1 and CT2 consisted of 85% polygonal ferrite with the average size of 5.5 mm. In the ferrite near the martensite/ ferrite interface, numerous dislocations were observed (Figs. 2b and 3b), which were induced by martensite transformation and affected the yielding behavior of dual phase steel [8]. The martensite morphology exhibited much difference between steels CT1 and CT2. Fig. 2c and d shows that the surface of martensite islands was smooth, suggesting no significant martensite tempering occurred at coiling temperature of 150 °C. However, presented in Fig. 3c and d, fine carbide particles were clearly observed in the martensite islands and the interface of ferrite and martensite exhibited less tetragonal, which inferred the martensite tempering at coiling temperature of 260 °C. It is mentioned that there were also few non-tempered martensites with small size in steel CT2 (Fig. 3e). The possible reason is that these small-sized martensites are highly enriched with carbon, forming at relative low temperature and difficult to be tempered. It indicated that auto-tempering occurred in the plain C-Mn dual phase steel when the coiling temperature was 260 °C. In the dual phase steel, the Ms was calculated as follows [12]:
Ms = 539 − 423C − 30.4Mn − 17.7Ni − 12.1Cr − 11.0Si − 7.0Mo
(1)
The carbon content of untransformed austenite was estimated by:
wγ = (w0 − fα wα )/(1 − fα )
(2)
where wγ, w0 and wα were the carbon contents of untransformed austenite, prior austenite and ferrite, respectively, and fα is the volume fraction of ferrite. In this work, the amounts of C, Si and Mn were 0.06%, 0.45% and 1.55% respectively, and that of Ni, Cr and Mo were ignored. The Ms of untransformed austenite estimated from (Eqs. (1) and 2) was 320 °C. Therefore, the martensite transformation accompanied with auto-tempering would occur at the coiling temperature of 260 °C, which was far more close to the Ms than 150 °C. It was reported that the carbides formation during tempering was suppressed by increasing the contents of Mn and Si [13–17]. However, the effect of Mn was weak when its content was less than 2% [15]. To suppress cementite formation, the Si content
100
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 2. Microstructures of steel CT1: (a) microstructure observed by OM, (b) dislocations observed by TEM, and (c) and (d) non-tempered martensite observed by SEM and TEM respectively.
Fig. 3. Microstructures of steel CT2: (a) microstructure observed by OM, (b) dislocations observed by TEM, (c) and (d) tempered martensite observed by SEM and TEM respectively, and (e) non-tempered martensite observed by SEM.
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
should be increased to 1.5% [16,17]. Therefore, the 1.55% Mn and 0.45% Si in the experimental steel may not play effectively role in suppressing carbide formation and increase the trends of autotempering. Auto-tempering led to the softening of both ferrite and martensite. When the coiling temperature increased from 150 °C to 260 °C, the micro-hardness of ferrite decreased from 148 HV to 139 HV, and that of martensite dropped from 346 HV to 269 HV. The softening degree of martensite was much larger than that of ferrite, causing the reduction of hardness difference between these two phases. This remarkable softening of martensite after autotempering is associated with the migration of carbon atoms, precipitation of particles and recovery of dislocation structure and so on [18].
101
Table 1 Tensile properties of the steels. Steel Yield Strength (MPa)
Tensile Strength (MPa)
Uniform Elongation (%)
Total Elongation (%)
Area Reduction (%)
CT1 CT2
598 551
16.5 13.0
35.5 36.5
71 86
334 394
3.2. Mechanical properties at different coiling temperatures 3.2.1. Tensile properties In the dual phase steel with low alloy, the auto-tempering is readily occurred at the coiling temperature close to Ms, as mentioned above, which will alter the mechanical properties of the steel. As shown in Fig. 4, steel CT1 exhibited a continuous yielding that was typical for dual phase steel. However, steel CT2 had a yield plateau with a pronounced increase in yield strength. It is known that the continuous yielding of dual phase steel was mainly attributed to the mobile dislocations produced by martensite transformation at low temperature [1,7,19,20], for example in Fig. 3a. Dislocations, one of structure defects in crystalline solid solutions, are disturbed regions with low thermodynamically stability, and attract solute segregation when diffusion is enable [21–23]. Therefore, when the coiling temperature increased to 260 °C, C and N atoms was considered to diffuse to the dislocations and pinned them. Many researches had also demonstrated that the interstitial atoms of C and N could segregate and pin the dislocations when the tempering temperature increased to 250 °C, leading to the rise in yield strength and the appearance of yield platform [7–9,19,23]. In this study, the yield plateau and the significantly increase in yield strength, to some extent, revealed that interstitial atoms of C and N segregated and pinned the dislocations during auto-tempering. As listed in Table, 1, steel CT1 was characterized by high tensile strength and strong strain hardening ability, which attributed to the high hardness of quenched martensite [24–26]. For steel CT2, the tensile strength decreased distinctly due to the softening of martensite. The reduction of hardness difference between ferrite
Fig. 4. Stress-strain curves of the steels during tensile test.
Fig. 5. True stress and strain hardening rate of the steels.
and martensite caused by auto-tempering resulted in the decrease of strain hardening ability [24,25]. When the coiling temperature increased from 150 °C to 260 °C, the uniform elongation decreased from 16.5% to 13%, which demonstrated that the auto-tempering stimulated the occurrence of necking. It is known that the necking appears when the strain hardening rate (ds/dε) is equal or less than the true stress (s) [27], and the corresponding strain is the uniform strain (εu), shown in Fig. 5. Thus, the increase of strain hardening rate or the decrease of true stress can postpone the occurrence of necking. After auto-tempering, both the strain hardening rate and the true stress reduced, but the degree of the former was more prominent. Therefore, the auto-tempering accelerated the occurrence of necking. Nevertheless, the stress-strain curve of steel CT2 had a distinct necking stage (Fig. 4), and the post-necking elongation and reduction of area reached 23.5% and 86%, respectively, higher than that of steel CT1. The data indicates that steel CT2 has the high fracture resistance. 3.2.2. Charpy impact energy and fractograph of the specimens Fig. 6 indicates that the toughness of steel CT2 was higher than that of steel CT1. The impact energy (vE) of steel CT1 was sensitive to the test temperature and decreased from 185 J to 10 J with the temperature decreasing from 20 °C to 100 °C. However, the toughness of steel CT2 under the test temperature range of 20 to 60 °C was such excellent (higher than 300 J) that the specimens cannot be completely separated. When the temperature decreased from 60 °C to 100 °C, the value of vE sharply dropped to 15 J from 320 J. The impact specimen macrofractographs of steel CT1 are given in Fig. 7. At 20 °C and 40 °C, the macrofractographs of steel CT1 contained the typical fibrous, radial, second fibrous and shear lip zones, marked by 1, 2, 3 and 4 respectively in Fig. 7a. The fibrous, second fibrous and shear lip zones of the steel CT1 at 20 °C consisted of dimples (Fig. 7d–f), reflecting the ductile fracture. The dimples in fibrous and second fibrous zones were stretched due to the shear stress. In the radial zone, the existence of cleavage facets, tearing ridges and cracks indicated the brittle fracture (Fig. 7g). As Shown in Fig. 7b and c, with decrease in test
102
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 6. Charpy V-notch impact properties: (a) impact energy versus test temperature and (b) the specimens after impact test.
Fig. 7. Impact specimen fractographs of steel CT1: (a)–(c) low magnification fractographs tested at fractographs of fibrous, second fibrous, shear lip and radial zone respectively.
temperature, the area of fibrous, second fibrous and shear lip zones decreased, replaced by the radial zone. At 60 °C, the radial zone occupied almost all the fracture surface of specimens. For steel CT2, shown in Fig. 8a, d and e, the fracture surfaces of unseparated specimen tested at 60 °C consisted of dimples, and the branching cracks were observed in the fracture surface. The macrofractures of steel CT2 at 80 °C (Fig. 8b) were familiar to that of steel CT1 at 20 °C and 40 °C. However, Fig. 8f–h shows that the dimples of steel CT2 were larger than that of steel CT1, indicating the superior toughness. At 100 °C, the fracture surface of specimens were mainly comprised of radial zone, which proven the low toughness at this temperature. Assuming that the test temperature corresponding to where the area of radial zone occupies 50% of the fracture surface is the ductile-to-brittle transition temperature (DBTT), the DBTT reduced to 80 °C from 40 °C when the steel was auto-tempered. The transition of ductile-to-brittle attributed to the competition of plastic deformation and brittle fracture at the tips of cracks or defects, and the DBTT was determined by the difference between tensile stress near the crack tips and the brittle fracture stress [28– 30]. According to Yoffee diagram [28,29], when the effective yield stress, at which significant deformation occurs, exceeds the brittle fracture stress, the brittle fracture occurs. Although the yield
20 °C,
40 °C and
60 °C, respectively, and (d)–(g) high magnification
strength of steel CT2 was higher than that of steel CT1 during tensile test, the stress of steel CT2 was lower than steel CT1 when the strain exceeds 0.01, shown in Fig. 4, which resulted from the softening of ferrite and martensite. In addition, the softening of steel CT2 caused by auto-tempering led to the extensive plastic deformation before fracture. It was clear that the extensive plastic deformation at the crack tip limited the local stress and inhibited the brittle fracture [31]. Therefore, the auto-tempering played an effective role in delaying the transition of ductile-to-brittle.
4. Discussion 4.1. Effects of auto-tempering on tensile fracture behavior Fig. 9 indicates that both the two steels fractured under ductile mechanism with cup-and-cone fracture. The steels suffered an obvious plastic deformation before fracture. However, the degree of section contraction of steel CT2 was much larger than that of steel CT1, indicating that auto-tempering provided superior ductility for DP steel. The fracture surfaces included two parts: central fiber zone and surrounding shear lip, marked in Fig. 9. During the tensile test, the microvoids virtually begin to initiate before
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
103
Fig. 8. Impact specimen fractographs of steel CT2: (a)–(c) low magnification fractographs tested at 60 °C, 80 °C and 100 °C, respectively, (d) and (e) local magnification fractographs of specimen tested at 60 °C, and (f)–(i) high magnification fractographs of fibrous, second fibrous, shear lip and radial zone respectively.
Fig. 9. Fracture surface of the tensile specimens for steel CT1 (a) and CT2 (b).
104
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 10. Subsurfaces of the main fracture surface in tensile specimens for steels CT1 (a) and CT2 (b).
necking. As the necking occurs, the stress concentration significantly increases and then induces more microvoids. The coalescence of microvoids leads to the microcracks, which grow perpendicular to the applied stress and evolve into the central fiber zone. As the necking further continues, the crack propagates along the 45° maximum shear stress direction and develops into the shear lip [28,32]. The shear lip of steel CT2 exhibited ring-like with “Λ” shape, which was caused by the severe deformation before fracture. As shown in the micrographs at high magnification, the fracture surfaces of the steels were composed of dimples. Nevertheless, the dimples in the fiber zone of steel CT2 were much larger and deeper than that of steel CT1, which may be the reason the post-necking elongation of steel CT2 was much larger than that of steel CT1. Fig. 10 shows the micrographs at subsurface under the main fracture surface of tensile specimens. The ferrite grains and martensite islands were elongated in the necking zone of steel CT1, but the elongated extent of ferrite grains was much higher than that of martensite islands due to the large hardness difference between ferrite and non-tempered martensite. The deformation incompatibility between ferrite and martensite resulted in high stress concentration, especially where the shape of martensite was tetragonal and sharp. Since the deformation between ferrite and martensite was highly incompatible (Fig. 10a) and the pure martensites were tetragonal (Fig. 2c) in steel CT1, lots of microvoids initiated around the tips of the elongated martensite islands (indicated by arrows in Fig. 10a), which led to the relatively small post-necking elongation. For steel CT2, both ferrite grains and tempered martensite islands were severely elongated as fibrous structure. The larger deformation and the better deformation compatibility between ferrite and tempered martensite were due to the softening and reduction of hardness difference between these two phases.
Fig. 10b indicates that the microvoids in steel CT2 tended to initiate near the non-tempered martensite islands and inclusions because the deformations of ferrite matrix and non-tempered martensite (or inclusion) were still incompatibility. Since most of martensite islands were auto-tempered and became less tetragonal (Fig. 3c) and the ferrite and martensite compatibly deformed as fibrous structure, the stress concentration was small compared to steel CT1. Therefore, the nucleation sites of microvoids in steel CT2 (indicated by arrows in Fig. 10b) was far less than that in steel CT1. The fewer nucleation sites of microvoids could explain why the dimples in fracture surface of steel CT2 were much larger and deeper than steel CT1. Additionally, the cracks in steel CT2 were expected to propagate more difficult to across the section of tensile specimen due to the fibrous structure normal to the propagation direction of cracks. Although both of two steels fractured with ductile mechanism, steel CT2 exhibited the higher ability to prevent the crack formation and propagation. 4.2. Effects of auto-tempering on impact fracture behavior To investigate the fracture process and the fracture behavior of the steels at different temperatures, the impact load (P), impact energy and deflection (u) were recorded by instrument during impact test. Fig. 11a and b show that the P-u curves of steel CT1 at 20 °C and 40 °C had the similar shape. At the beginning, the impact load increased rapidly to the peak value Pm. In this stage, the elastic deformation and plastic deformation occurred successively, and the absorbed energy corresponding to the Pm was crack initiate energy (vEi). The value of vEi at 20 °C and 40 °C were stable at 40–45 J in steel CT1. Then, with the increase in u, the value of P slowly decreased to Py, then dropped to Pa and slowly reduced to zero (P0), successively. The stages of PmPy, PyPa and PaP0 reflected the processes of ductile propagation, brittle fracture and
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
105
Fig. 11. Typical curves of impact load and impact energy versus deflection at different temperatures (black lines-steel CT1, red lines–steel CT2, solid lines - impact load and dashed lines - impact energy): (a) 20 °C, (b) 40 °C, (c) 60 °C (d) 80 °C and (e) 100 °C. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
ductile fracture, respectively. The total absorbed energy of these three stages was crack propagation energy (vEp). When the test temperature declined from 20 °C to 40 °C in steel CT1, the value of vEp decreased from 139 J to 95 J. This suggested that the lower temperature accelerated the crack propagation and eventually deteriorates the toughness. The microvoids and cracks underneath the fracture surface of specimen at 20 °C for steel CT1 are presented in Fig. 12a and b. Microvoids were observed underneath the fibrous zone of fracture surface, implying the ductile fracture. Underneath the radial zone, the cracks exhibited transgranular with the brittle mechanism. The characteristics of P-u curves and the observation of cracks underneath the fracture surface in steel CT1 demonstrated the mixed mechanism of ductile and brittle in the range of 20 to 40 °C. In the P-u curves at 60 to 100 °C, the ductile propagation and ductile fracture stages disappeared (Fig. 11c–e), which suggests that the fracture mechanism converted from mixture of ductile
and brittle fracture to complete brittle fracture. In Fig. 11a–c, the P-u curves of steel CT2 at 20 to 60 °C exhibited a distinct ductile propagation stage. It indicated that the auto-tempering played an effective role in suppressing the propagation of cracks. With increase in u, P decreased slowly from Pm, then rapidly declined to the first plateau stage, and the second plateau stage, and finally dropped to zero when the specimen was struck away from the specimen holder. This similar plateau-step of P-u curve was also observed in the low-carbon steel with severe elongated structure, multilayer composite steel and glass-fiber/ vinyl-ester composite [29,33,34]. In these materials, the branching cracks tended to propagate along the interfaces of elongated structure or fiber planes that were normal to the propagation direction of the main crack. In this case, the propagation of the main cracks was restrained, resulting in the plateau stage of P-u curve. The impact load dropped when the new cracks initiated. In the fractographs (Fig. 8a, d and e) and micrographs underneath the
106
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
Fig. 12. Cracks underneath the fracture surface of impact specimens: (a) microvoids underneath fibrous zone of steel CT1 at 20 °C, (b) transgranular cracks underneath radial zone of steel CT1 at 20 °C, (c) and (d) fiber structure and branching cracks of steel CT2 at 20 °C and 60 °C, respectively, (e) microvoids underneath fibrous zone of steel CT2 at 80 °C, and (f) transgranular cracks underneath radial zone of steel CT2 at 80 °C.
fracture surface (Fig. 12b and c) of the impact specimens, the branching cracks were observed clearly in steel CT2 at 20 °C and 60 °C. The branching cracks were attributed to the acquired fibrous structure. Presented in Fig. 12b and c, the branching cracks propagated along the interface of fibrous structure. The shape of P-u curve of steel CT2 at 80 °C (Fig. 11d) was similar to that of steel CT1 at 20 to 40 °C, showing the similar fracture mechanism. In Fig. 12e and f, similar to micrographs of steel CT1 at 20 °C (Fig. 12a and b), the microvoids and the transgranular cracks were observed in fibrous and radial zones, respectively. At 100 °C, the P-u curve of steel CT2 exhibited the brittle type. Consequently, with the decrease in test temperature from 20 °C to 100 °C, the fracture mechanism of steel CT2 transferred from ductile fracture accompanied with branching crack to mixture of ductile and brittle fracture, and finally to brittle fracture.
5. Conclusions
1. The microstructure consisted of polygonal ferrite and non-
tempered martensite islands in the hot rolled plain C-Mn DP steel coiled at 150 °C. When the coiling temperature was 260 °C, the interface of ferrite and martensite became less tetragonal, and the fine carbide particles dispersed in the martensite islands were also observed. Both of them reflected the occurrence of auto-tempering, leading to the softening of ferrite and martensite and the reduction of hardness difference between these two phases. 2. The auto-tempering resulted in the appearance of yield platform with the increase in yield strength, while both the strain hardening rate and the tensile strength decreased. In particular, the auto-tempered DP steel exhibited a distinct necking stage, implying the superior ductility. During tensile test, the ferrite grains and tempered martensite islands were severely elongated as fibrous structure, which effectively postponed the tensile fracture. 3. Compared with the non-tempered steel, the auto-tempered DP steel had higher impact energy and lower DBTT. When the test temperature decreased from 20 to 60 °C, the fracture mechanism of non-tempering DP steel converted from mixture of ductile and brittle fracture to complete brittle fracture. At the test temperature above 60 °C, the impact specimens of auto-
C.-n. Li et al. / Materials Science & Engineering A 665 (2016) 98–107
tempered steel fractured by ductile mechanism accompanied with branching cracks. And the fracture mechanism transferred from mixture of ductile and brittle fracture to brittle fracture with the test temperature decreasing from 80 °C to 100 °C.
Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51234002) and Fundamental Research Funds for the Central Universities (N130407001).
References [1] M. Sarwar, R. Priestner, Influence of ferrite-martensite microstructural morphology on tensile properties of dual-phase steel, J. Mater. Sci. 31 (1996) 2091–2095. [2] A. Bag, K.K. Ray, E.S. Dwarakadasa, E.S. Dwarakadasa, Influence of martensite content and morphology on tensile and impact properties of high-martensite dual-phase steels, Metall. Mater. Trans. A 30 (1999) 1193–1202. [3] S. Oliver, T.B. Jones, G. Fourlaris, Dual phase versus TRIP strip steels: microstructural changes as a consequence of quasi-static and dynamic tensile testing, Mater. Charact. 58 (2007) 390–400. [4] M. Mazinani, W.J. Poole, Effect of martensite plasticity on the deformation behavior of a low-carbon dual-phase steel, Mater. Sci. Eng. A 38 (2007) 328–339. [5] P.H. Chang, Temper-aging of continuously annealed low carbon dual phase steel, Metall. Trans. A 15 (1984) 73–86. [6] C.F. Kuang, J. Li, S.G. Zhang, J. Wang, H.F. Liu, A.A. Volinsky, Effects of quenching and tempering on the microstructure and bake hardening behavior of ferrite and dual phase steels, Mater. Sci. Eng. A 613 (2014) 178–183. [7] S. Gunduz, Effect of chemical composition, martensite volume fraction and tempering on tensile behaviour of dual phase steels, Mater. Lett. 63 (2009) 2381–2383. [8] A.A. Sayed, S. Kheirandish, Affect of the tempering temperature on the microstructure and mechanical properties of dual phase steels, Mater. Sci. Eng. A 532 (2012) 21–25. [9] A. Kampa, S. Celottob, D.N. Hanlonb, Effects of tempering on the mechanical properties of high strength dual-phase steels, Mater. Sci. Eng. A 538 (2012) 35–41. [10] V.H.B. Hernandez, S.S. Nayak, Y. Zhou, Tempering of martensite in dual-phase steels and its effects on softening behavior, Metall. Mater. Trans. A 42 (2011) 3115–3129. [11] H. Matsuda, R. Mizuno, Y. Funakawa, K. Seto, S. Matsuoka, Y. Tanaka, Effects of auto-tempering behaviour of martensite on mechanical properties of ultra high strength steel sheets, J. Alloy. Compd. 557 (2013) 661–667. [12] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [13] S. Yan, X.H. Liu, W.J. Liu, H.F. Lan, H.Y. Wu, Comparison on mechanical properties and microstructure of a C-Mn-Si steel treated by quenching and partitioning (Q&P) and quenching and tempering (Q&T) processes, Mater. Sci. Eng. A 620 (2015) 58–66.
107
[14] J. Speer, D.K. Matlock, B.C. De Cooman, J.G. Schroth, Carbon partitioning into austenite after martensite transformation, Acta Mater. 51 (2003) 2611–2622. [15] J. Pacyna, A. Jedrzejewska-Strach, M. Strach, The effect of manganese and silicon on the kinetics of phase transformations during tempering - Continuous Heating Transformation (CHT) curves, J. Mater. Process. Technol. 64 (1997) 311–318. [16] D. De Knijf, E.P. Da Silva, Fojer, C. Fojer, R. Petrov, Study of heat treatment parameters and kinetics of quenching and partitioning cycles, Mater. Sci. Technol. 31 (2015) 817–828. [17] D. Hauserova, M. Duchek, J. Dlouhy, Z. Novy, Properties of advanced experimental CMnSiMo steel achieved by QP process, Procedia Eng. 10 (2011) 2961–2966. [18] C.E.I.C. Ohlund, E. Schlangen, S.E. Offerman, The kinetics of softening and microstructure evolution of martensite in Fe-C-Mn steel during tempering at 300 °C, Mater. Sci. Eng. A 560 (2013) 351–357. [19] S. Gündüz, Static strain ageing behaviour of dual phase steels, Mater. Sci. Eng. A 486 (2008) 63–71. [20] J. Kadkhodapour, S. Schmauder, D. Raabe, S. Ziaei-Rad, U. Weber, M. Calcagnotto, Experimental and numerical study on geometrically necessary dislocations and non-homogeneous mechanical properties of the ferrite phase in dual phase steels, Acta Mater. 59 (2011) 4387–4394. [21] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. I. Theoretical background, Acta Mater. 55 (2007) 5129–5138. [22] M. Kuzmina, M. Herbig, D. Ponge, S. Sandlobes, D. Raabe, Linear complexions: confined chemical and structural states at dislocations, Science 349 (2015) 1080–1083. [23] R.G. Davies, Early stages of yielding and strain aging of a vanadium-containing dual-phase steel, Metall. Trans. A 10 (1979) 1549–1554. [24] M.F. Ashby, Work hardening of dispersion-hardened crystals, Philos. Mag. 14 (1966) 1157–1178. [25] M.F. Ashby, The deformation of plastically non-homogeneous materials, Philos. Mag. 21 (1970) 399–424. [26] S. Ankem, H. Margolin, C.A. Greene, B.W. Neuberger, P.G. Oberson, Mechanical properties of alloys consisting of two ductile phases, Prog. Mater. Sci. 51 (2006) 632–709. [27] M. Saeglitz, G. Krauss, Deformation, fracture, and mechanical properties of low-temperature-tempered martensite in SAE 43xx steels, Metall. Mater. Trans. A 28 (1997) 377–387. [28] T. Inoue, F.X. Yin, Y. Kimura, K. Tsuzaki, S. Ochiai, Delamination effect on impact properties of ultrafine-grained low-carbon steel processed by warm caliber rolling, Metall. Mater. Trans. A 41 (2010) 341–355. [29] J.W. Morris, Stronger, tougher steels, Science 320 (2008) 1022–1023. [30] D.L. Bourell, Cleavage delamination in impact tested warm-rolled steel, Metall. Trans. A 14 (1983) 2487–2496. [31] J.W. Morris, Steels: For Low Temperature Applications, Pergamon Press,, Oxford, 1993. [32] Y.B. Xu, X.D. Tan, X.L. Yang, Z.P. Hu, D. Wu, G.D. Wang, Microstructure evolution and mechanical properties of a hot-rolled directly quenched and partitioned steel containing proeutectoid ferrite, Mater. Sci. Eng. A 607 (2014) 460–475. [33] J.F. Kalthoff, Characterization of the dynamic failure behaviour of a glass-fiber/ vinyl-ester at different temperatures by means of instrumented Charpy impact testing, Compos. Part B 35 (2004) 657–663. [34] M. Pozuelo, F. Carreno, O.A. Ruano, Delamination effect on the impact toughness of an ultrahigh carbon–mild steel laminate composite, Metall. Mater. Trans. A 41 (2010) 341–355.