Avoiding the surface cracking is one of the important factors in improving the slab quality. The hot ductility is a significant parameter for estimating the possibility of cracking formation. The change of the hot ductility associates closely with changes of microstructure such as the precipitation of second phases , the size of prior austenite grain  and the transformation of γ–α [3, 4]. These changes usually occur in the straightening operation around 700—1,000°C during the continuous casting process . Alloy elements added in steel have great influences on the changes of microstructure.
Many researchers have revealed the effect of boron addition on the hot ductility. Generally, it is considered that the hot ductility will be improved with the formation of boron-containing precipitates in the low-carbon microalloyed steel [6–9]. The formation of proeutectoid ferrite film at the grain boundaries is negative for the hot ductility during continuous casting. Trace boron has great influence on the transformation of γ–α. In order to assure the influence, titanium was added in the boron-containing steel . It was reported that the hot ductility was raised with the titanium addition [11, 12]. However, the studies on the influence of titanium and trace boron on the hot ductility are limited.
The objective of this work is to study the compound effect of titanium and trace boron on the hot ductility in low-carbon Nb-containing steel. In this study, the hot tensile test was used to determine the hot ductility. The scanning electron microscopy (SEM), optical microscopy (OM) and the transmission electron microscopy (TEM) were used for observing the fracture morphology and microstructure of steel samples. The effect of boron and titanium addition on the hot ductility of low-carbon Nb-containing steel was investigated.
Two low-carbon Nb-containing steel ingots with or without boron and titanium addition were prepared by laboratory vacuum induction melting furnace, then hot rolled into 20-mm-thick plates. The Nb-containing steel ingot containing 0.0015 mass% boron and 0.0125 mass% titanium was designated as steel sample A. The Nb-containing steel ingot without boron and titanium was designated as steel sample B. The equilibrium temperature of austenite/ferrite phase boundary (Ae3) for two steel samples was calculated using Thermo-Calc (TCFE 6 database) as shown in Table 1. The chemical compositions of two steel samples were given in Table 1. Cylindrical specimens of 10 mm in diameter and 120 mm in length for hot tensile test were machined from the plates with their longitudinal axes parallel to the rolling direction, as shown in Figure 1.
The hot tensile tests were conducted using Gleeble 3500 hot deformation simulator in argon gas atmosphere. The thermomechanical cycle used in this study is shown schematically in Figure 2. The specimens were heated from room temperature to 1,350°C at 10°C/s, held for 3 min to allow microstructural homogenization and the dissolution of the most precipitates. The specimens were then cooled to the test temperature range of 700–1,100°C at a cooling rate of 3°C/s, and held at these temperatures for 1 min, and then deformed to failure at a strain rate of 2.0 × 10−3/s. After failure, the samples were quenched by water spraying to remain the microstructure at each testing temperature.
The reduction of area (RA) was determined to estimate the hot ductility of steel samples. The metallographs and fractographs of both steel samples were examined by OM and SEM. The morphology of precipitates in the fractured zone, mounted on carbon replicas film, was determined by TEM equipped with energy-dispersive X-ray spectroscope (EDS).
Figure 3 shows the hot ductility curves for both steel samples. Both steel samples have similar changes in the range of testing temperature. Mintz  found that the possibility of crack improves distinctly when the values of RA were less than 40%. If the RA was greater than 60%, the cracks possibility will be reduced [14, 15]. Therefore, in this study, the region where the RA greater than 60% is defined as region I, the value less than 40% is region III and the range between the two regions is region II.
Both steel samples have higher RA values, which are greater than 60% from 1,100°C to 950°C (region I). It indicated that both B–Ti-containing steel and B–Ti-free steel have higher hot ductility in region I. In the range of 950–900°C (region II), the RA values drop sharply. In the range of 900–700°C (region III), the RA values of two steel samples are less than 40%, except the value of steel sample A fractured at 900°C. In region II, it can be seen that the decrease in RA value of steel sample B is more significant than that of steel sample A. In region III, the steel sample A has higher hot ductility than B from 900°C to 750°C.
Figure 4 shows the longitudinal cross-sectional microstructures of two steel samples at testing temperature of 900°C, 850°C and 800°C.
At 900°C, the microstructure of both steel samples is martensite as shown in Figure 4(a) and (b). At 850°C, the film-like proeutectoid ferrite formed along the austenite grain boundary in steel sample B is shown in Figure 4(d), whereas the transformation has not occurred in steel sample A. At 800°C, proeutectoid ferrite along grain boundary also formed in the steel sample A and the intragranular ferrite were found as shown in Figure 4(e). In steel sample B, the ferrite film grew further as shown in Figure 4(f). The temperatures of Ae3 calculated by Thermol-Calc are 826°C for steel sample A and 825°C for steel sample B. However, it can be obtained from Figure 4 that the actual temperature of γ/α transformation of B–Ti-free steel is higher than 850°C, and the actual temperature for B–Ti-containing steel is between 850°C and 800°C.
Influence of dynamic recrystallization on the hot ductility
Two steel samples have excellent hot ductility in region I (1,100–950°C) due to the occurrence of dynamic recrystallization. From Figure 5, it is shown that the stress is raised and the elongation is reduced with the decrease of temperature from 1,100°C to 700°C. It is caused by the precipitation of second phases which can pin the grain boundaries and play a role of precipitation strengthening. From 1,100°C to 1,000°C, the dynamic recrystallization occurs distinctly. Many studies indicate that dynamic recrystallization is beneficial for increasing the hot ductility [16, 17].
Figure 6 shows the SEM images of fracture surface of both steel samples fractured at 1,100°C. There are many deep dimples on the fracture surface of both steel samples tested at 1,100°C. Considerable plastic deformation occurs on the zone around the dimples. It indicates that the cavities inside the grain came into being necking after nucleation and growth in the process of deformation. The formation of necking also promotes the growth of cavities until they linked up to be microcracks. It can be concluded that the fracture type is intragranular ductile fracture for both steel samples at 1,100°C.
In region I (950–1,100°C), the temperature is high enough for the occurrence of dynamic recrystallization and provides the sufficient driving force for grain boundaries. Grain boundaries moved away from microcracks. Therefore, cracks were wrapped inside the grain. It requires great stress for intragranular cracks getting through grain to link together. So, the propagation of microcracks was suppressed. For the occurrence of dynamic recrystallization, the fracture type of both steel samples is intragranular fracture.
Influence of precipitates on the hot ductility
From 950°C to 900°C, the hot ductility of both steel samples drops suddenly. It is caused by the precipitates of pinned grain boundaries. The grain boundaries pinned by fine precipitates migrate hardly, only to move in grain sliding. If the microcracks exist on the grain boundaries, the cracks will propagate along grain boundaries and lead to intergranular fracture finally. The grain boundaries that are not pinned have greater migration ability. Accordingly, the possibility of intergranular fracture decreases. The formation sequences of precipitates in both steel samples were calculated using Thermol-Calc (TCFE6 database), and the results are shown in Figure 7.
In Figure 7(a), it is shown that the precipitation starting temperature of TiN is 1,450°C. It indicates that TiN could not dissolve completely at 1,350°C. Figure 8 shows the morphology of TiN in steel sample A which quenched after solution treatment at 1,350°C for 5 min. TiN with this size can play a role in refining the austenite grain. The refinement of grain is favorable for increasing the hot ductility [18, 19]. But the effect mainly occurred under the condition of austenite grain size less than 50 μm . From Figure 4(e) and (f), although the austenite grain size of steel sample A is finer than steel sample B, the grain size of both steel samples is greater than 400 μm. Grain refinement will not affect the hot ductility.
Figure 7(a) shows the sequence of precipitates in steel sample A. The Ti/N ratio of steel sample A is less than 3.42 which is the ideal Ti/N ratio for fixing all N in steel. Thus, there are residual N reacting with other alloy elements such as B, Nb and Al. The amount of BN is small due to Ti addition. The formation of NbN and AlN is rare for most of N reacted by Ti and B. Therefore, the Nb-containing precipitates are NbC. TiN precipitates dispersively and can play a role of nucleation site for BN and NbC. The effects make the BN and NbC also precipitating dispersively . It is useful for decreasing the precipitation amount at grain boundary and reducing the effect of pinning grain boundary. The reason for the hot ductility drop of B–Ti-containing steel is caused by the precipitation of TiN and NbC which can pin the grain boundary. Figure 9(a) shows the precipitates of TiN and NbC in steel sample A at testing temperature of 950°C.
From Figure 7(b), AlN precipitated later than Nb(C, N). The precipitation temperature is about 1,050°C which is in the region of austenite single phase. Although the thermal condition meets the precipitation of AlN, the nucleation is difficult because of the slower diffusion speed of Al in austenite [20, 21]. Therefore, fine Nb(C, N)-pinned grain boundary is the reason for the hot ductility drop of B–Ti-free steel. The Nb(C, N) precipitates in steel sample B are shown in Figure 9(b).
The hot ductility drop of B–Ti-containing steel is less than B–Ti-free steel from 950°C to 900°C. The reason is Ti addition. TiN precipitates dispersively and takes an effect as nucleation for BN and NbC. Comparing with steel sample B, the amount of precipitates on the grain boundaries in steel sample A is small, the effect of pinning grain boundary is weaker. It is helpful for increasing the migration ability of grain boundary. Therefore, the decrease in RA of steel sample A is more significant than that of steel sample B in region II.
Influence of γ/α transformation on the hot ductility
The transformation of γ/α is the key influence on the hot ductility of two steel samples. Film-like proeutectoid ferrite is formed along the austenite grain boundary. The ferrite is more deformable than austenite. Hence, the stress concentrates on the ferrite in the early of transformation of γ/α. The strength of the ferrite is approximately a quarter of austenite. If the stress exceeds the critical strength of ferrite, the cavities or microcracks will form at grain boundary and will lead to intergranular fracture finally. When the fine precipitates pin on the intergranular ferrite, the intergranular fracture will be aggravated.
Figure 10(a) and (b) shows the fracture surface of both steel samples at 800°C. At this temperature, γ/α transformation has occurred in both steel samples. The fracture surfaces of both steel samples manifest the feature of intergranular fracture. But there are some differences between two steel samples. It can be seen that the grain surfaces of the steel sample A undergone greater deformation before break. The grain surfaces have shallow dimples on them, as shown in Figure 10(a). Nevertheless, the grain surfaces of the steel sample B are smooth and the cracks are found between grains, as shown in Figure 10(b). The results show that the rupture type of the B–Ti-containing steel is intergranular fracture with the characters of shallow dimple fracture and the type of the B–Ti-free steel is intergranular brittle fracture at 800°C.
In steel sample A, the solute boron segregated on the austenite grain boundary. The ferrite was inhibited greatly for the occupation of ferrite nucleation sites and energy decrease of grain boundary. At 850°C, the intergranular ferrite was formed in steel sample B as shown in Figure 4(d), whereas the transformation of γ/α has not occurred as shown in Figure 4(c). In region III, the lowest hot ductility of B–Ti-free steel appears earlier than B–Ti-containing steel. Boron addition makes the hot ductility trough to move to the low temperature.
At 800°C, the intergranular ferrite in steel sample B further grew as shown in Figure 4(f). It is found that the ferrite formed along austenite in steel sample A. As mentioned in “Results” section, a certain content of intragranular ferrite was found in steel sample A as shown in Figure 4(e). The formation of intragranular ferrite was associated with TiN. TiN distributed dispersively and take an effect of nucleation site for intragranular ferrite. The result is consistent with the finding reported by H. H. Jin . The amount of intragranular ferrite increases with the decrease of intergranular ferrite. The stress concentration on the intergranular ferrite reduced. The austenite grain deformed uniformly. It is favorable for increasing the hot ductility. Due to the greater amount of AlN formed in steel sample B shown in Figure 7(b), the AlN precipitated on the intergranular ferrite film aggravate the tendency of intergranular fracture. Therefore, the fracture type of steel sample B is intergranular brittle fracture as shown in Figure 10(b).
The high temperature and deformation lead to the occurrence of dynamic recrystallization and the improvement of migration ability of austenite grain boundary. With the improvement in the hot ductility, the intergranular cracks were inhibited in low-carbon Nb-containing steel.
The addition of boron in low-carbon microalloyed steel inhibits the transformation of γ/α, which makes the hot ductility trough to move to the low temperature.
The addition of Ti in low-carbon microalloyed steel decreases the amount of AlN, NbN and BN.
TiN can take an effect of nucleation site for fine precipitates which can pin the grain boundary. Therefore, the hot ductility of low-carbon Nb-containing steel improves caused by the amount of the precipitates decrease at grain boundary due to the dispersive distribution of TiN. TiN can also take an effect of nucleation site for intragranular ferrite. The austenite grains deform uniformly because of the decrease of intergranular ferrite. Both effects caused by Ti addition are beneficial for improvement of hot ductility.
Mintz B, Arrowsmith JM. Hot ductility behavior of C-Mn-Nb-Al steels and its relationship to crack propagation during the straightening of continuously cast strand. Met Technol 1979;6:24–32. CrossrefGoogle Scholar
Mintz B, Yue S, Jonas JJ. Hot ductility of steels and its relationship to the problem of transverse cracking during continuous casting. Int Mater Rev 1991;36:187–220. CrossrefWeb of ScienceGoogle Scholar
Grabke HJ. Impurities in engineering materials. New York, NY: Marcel Dekker Inc, 1999. Google Scholar
Yamamoto K, Susuki HG, Oono Y, et al. Formation mechanism and prevention method of facial cracks of continuously cast steel slabs containing boron. Trans Iron Steel Inst Jpn 1987;73:115–122. Google Scholar
Cai KK, Dang ZJ, Zhang Y, et al. High temperature mechanical properties in continuous casting. J Univ Sci Technol Beijing 1993;15:24–6. Google Scholar
Suzuki HG, Nishimura S, Imamura J, et al. Hot ductility in steels in the temperature range between 900 and 600 deg C: related to the transverse facial cracks in continuously cast slabs. Trans Iron Steel Inst Jpn 1981;67:1180–9. Google Scholar
Mejia I, Altamirano G, Bedolla-Jacuinde A, et al. Effect of boron on the hot ductility behavior of a low carbon advanced ultra-high strength steel (A-UHSS). Metall Mater Trans A 2013;44:5165–76. CrossrefWeb of ScienceGoogle Scholar
Gladman T, Mcivor ID, Pickering FB. Effect of carbide and nitride particles on the recrystallization of ferrite. J Iron Steel Inst 1971;209:380–90. Google Scholar
Gladman T. Proceedings of the Second International Conference on HSLA Steels, Processing, Properties and Applications, October 28–November 2, 1990, Beijing, China, 1992:3–14. Google Scholar
About the article
Published Online: 2015-01-14
Published in Print: 2015-12-01