Effects of austempering temperature on bainitic microstructure and mechanical properties of a high-C high-Si steel

Author’s Accepted Manuscript Effects of austempering temperature on bainitic microstructures and mechanical properties of a high-C high-Si steel Jiali Zhao, Bo Lv, Fucheng Zhang, Zhinan Yang, Lihe Qian, Chen Chen, Xiaoyan Long www.elsevier.com/locate/msea

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S0921-5093(18)31523-5 https://doi.org/10.1016/j.msea.2018.11.004 MSA37124

To appear in: Materials Science & Engineering A Received date: 14 September 2018 Revised date: 31 October 2018 Accepted date: 1 November 2018 Cite this article as: Jiali Zhao, Bo Lv, Fucheng Zhang, Zhinan Yang, Lihe Qian, Chen Chen and Xiaoyan Long, Effects of austempering temperature on bainitic microstructures and mechanical properties of a high-C high-Si steel, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.11.004 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effects of austempering temperature on bainitic microstructures and mechanical properties of a high-C high-Si steel Jiali Zhaoa, Bo Lvb*1, Fucheng Zhanga*, Zhinan Yangc, Lihe Qiana, Chen Chena, Xiaoyan Longa a

State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, China b

c

College of Environmental and Chemical Engineering, Yanshan University, Qinhuangdao 066004, China National Engineering Research Center for Equipment and Technology of Cold Strip Rolling, Yanshan University, Qinhuangdao 066004, China

* Corresponding author. Tel.: +86 335 8055883; E-mail address: [email protected] (B. Lv).

* Corresponding author. Tel.: +86 335 8063949; E-mail address: [email protected] (F.C. Zhang).

Abstract: High tensile strength of ~1550 MPa, uniform elongation of ~30%, and impact energy of ~80 J were achieved in a high-C high-Si bainitic steel, by austempering around the nose-tip temperature of the time–temperature–transformation (TTT) diagram. This bainitic microstructure revealed multiple lengths of nanobainite plates and various sizes and shapes of retained austenite. Its bainite transformation during austempering and microstructure evolution during tensile deformation were discussed in depth. Results demonstrated that large blocks of retained austenite (with inhomogeneous carbon distribution) transformed partly and gradually to multiple martensite grains as strain increases, due to the strong effect of carbon on mechanical stability of retained austenite. Ductile long nanobainite plates (~0.13 at.% C) suppressed the propagation of microcracks, because of stress relief effect ahead of the current microcracks. Keywords: Bainitic steel; Dualphase; Multiscale; Mechanical properties 1. Introduction During bainite transformation, supersaturated bainite plates grow firstly via a displacive mechanism. Then, the excess carbon in bainite plates is partitioned into untransformed austenite or transforms into the carbides in bainite plates [1]. Carbon-free bainite, typically, is synthesized from high-Si steels through austempering at low temperature (between Ms and Bs). Owing to sufficient amount of Si retarding the precipitation of carbides into bainite plates, the excess carbon in bainite plates is only partitioned into untransformed austenite [2]. Then, carbon-rich austenite is retained at room temperature. Thus, the bainitic microstructure will consist of carbide-free bainite plates and a large number of retained austenite. Without carbide accelerating cleavage or void nucleation under

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applied stress, carbon-free bainitic steels exhibit good strength and ductility/toughness [3, 4]. Excellent combination of strength and ductility/toughness makes carbon-free nanobainitic steels widely applied in industries such as armor and assembled railway fork, etc., where the demanding service conditions require high strength and high impact resistance [5, 6]. Considerable works have addressed the optimization of the heat treatment [7–9] and the resulting microstructures and mechanical properties of these steels [10–12]. Many traditional approaches for strengthening carbon-free bainitic steels typically come at the grain refining strengthening [13–16]. For instance, by careful design in the chemical composition, a Ms temperature as low as 125 °C is obtain. This allows austempering at relatively low temperatures. Therefore, extremely fine microstructures (~30 nm thick bainite plates) and very high ultimate tensile strength (2.5 GPa) are generated [13]. However, the weak interaction of dislocations existed in fine microstructures due to the short free path of the dislocation movement causes very low work hardening rate. The low work hardening rate is negative for improving the ductility, especially the uniform elongation. Some studies show that tailoring microstructures to multiscale grain size distribution could enhance the ductility of metals without much strength loss [17, 18]. Accordingly, multiscale grain size distribution in carbon-free bainitic steels needs further study. In carbon-free bainitic steels, the metastable retained austenite can transform into martensite during the tensile test. A progressive increase in volume fraction of the hard phase (martensite) and additional plastic deformation due to transformation strains cause the extra work hardening [2]. This phenomenon is called transformation induced plasticity (TRIP) effect [19]. Various experimental investigations [20–22] show that the TRIP effect of retained austenite depends on its initial volume fraction and mechanical stability. The mechanical stability of retained austenite is affected by its grain size, morphology, carbon content, grain orientation, and neighboring phases [23–25]. The carbon content of retained austenite is one of the most important factors governing the mechanical stability of retained austenite. Thus, microstructure studies and especially quantification measuring of

local

carbon

content

of

retained

austenite

are

of

fundamental

interest

for

microstructure/properties control. Different techniques are available for measuring carbon content of retained austenite. X-ray Diffraction (XRD) is traditionally used for an overall and average value. For local and discrete measurements, however, atom probe tomography (APT) can be very powerful since it provides very high spatial resolution allowing for measurements in one austenitic grain. In this paper, a dualphase and multiscale carbide-free nanobainite microstructure with high tensile strength of ~1550 MPa, uniform elongation of ~30%, and impact energy of ~80 J was proposed. The bainite transformation during austempering was studied by means of optical 2

microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and APT techniques. Moreover, the martensite transformation from residual austenite and fracture mechanism during tensile deformation were analyzed in detail. 2. Experimental The chemical composition of the steel used in this investigation was as follows: Fe-0.70 C-2.59 Si-0.63 Mn-0.59 Cr wt.% (Fe-3.09 C-4.90 Si-0.61 Mn-0.60 Cr at.%). A 40 kg cast of tested steel was produced using a vacuum induction furnace and forged into a billet ingot with sectional dimensions of 40 mm × 85 mm. The austenized samples with an average austenite grain size of ~34 µm were quenched into salt bath and held isothermally at 260 °C, 300 °C, 330 °C, 350 °C, 370 °C, and 400 °C for 2h, followed by air cooling to room temperature. Fig. 1 shown the time– temperature–transformation (TTT) diagram of the designed steel in the bainite transformation temperature range. Under the combined effects of the transformation driving force and the diffusion rate of carbon atoms, the designed steel has the shortest incubation period and transformation time at 350 °C (i.e. the nose-tip temperature of the TTT diagram).

Fig. 1. TTT diagram of the designed steel.

Tensile tests were carried out on a MTS810 servohydraulic tensile testing machine at a strain rate of 0.002 s-1. Standard tensile samples with a gage diameter of 5 mm and a gage length of 25 mm were used. To establish the microstructure–property relationship, interrupted tensile tests at custom-made deformation stage were performed. Impact tests were performed using standard-sized Charpy-U samples on a 300 J Charpy testing machine. Tensile and impact tests were conducted by using two and three samples, respectively, to check the repeatability. Afterward, the mean values of the properties were calculated. Microstructures were characterized using OM (Axiovert 200MAT), SEM (Hitachi SU-5000, 15 kV), and TEM (JOEL-2010, 200 kV). The volume fraction and carbon content of retained austenite were determined by XRD (Cu-Kα radiation, scan rate: 2 °/min, scan step size: 0.02 °) [26]. 3

The quantification measuring of local carbon content of retained austenite and adjacent bainite plates was carried out by APT (Cameca Instruments LEAP 4000X HR). The LEAP was operated in voltage-pulse mode with a specimen temperature of 50 K, a pulse repetition rate of 200 kHz, and a pulse fraction of 0.2. The APT samples were prepared through a site-specific FIB lift-out technique. The Image Visualization and Analysis Software version 3.6 was used for 3-D reconstruction and composition analyses. 3. Results Fig. 2 shown the OM, SEM, and TEM images of the designed steel austempered at 260 °C, 350 °C, and 400 °C. No carbide precipitations were observed in the TEM images. This is because the extremely high content of silicon, a non-carbide forming element, retards the carbide formation [27]. Under increasing austempering temperature, the length of bainite sheaves increased. Long bainite sheaves traversed the prior austenite grains after austempering at 350 °C and above. When the sample was austempered at 350 °C, the prior austenite grains were divided to multiple regions by long bainite sheaves, and short bainite sheaves dispersed in these regions. Moreover, long bainite plates adjoined large blocky retained austenite and short bainite plates adjoined small ones.

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Fig. 2. Typical OM (left), SEM (middle), and TEM (right) images of the designed steel austempered at (a) 260 °C, (b) 350 °C, and (c) 400 °C. The insets in the SEM images represent the details of bainite sheaves. γb is blocky retained austenite, γf is filmy retained austenite, and αBF is bainite plates.

The average sizes and size distributions of bainitic ferrite and retained austenite were quantitatively measured and shown in Fig. 3. Average thicknesses of bainite plates and filmy retained austenite and the average size of blocky retained austenite increased slightly when the austempering temperature increased from 260 °C to 350 °C and then they increased substantially as the austempering temperature increased from 350 °C to 400 °C. Whereas, the increase in the average length of bainite plates exhibited the opposite trend. Obviously, all of the curves had a significant turning point at 350 °C (the nose-tip temperature of the TTT diagram). The average thicknesses of bainite plates and filmy retained austenite formed at 350 °C were ~80 nm and ~40 nm, respectively, similar to that at 260 °C. The average length of bainite plates formed at 350 °C was ~30 µm, which was approximately three times longer than that at 260 °C. The length of bainite plates formed at 350 °C ranged from 5 µm to 40 µm, and the size of blocky retained austenite ranged from 0.5 µm to 4.5 µm. These results shown that nanobainite plates and retained austenite formed at 350 °C displayed a multiscale size distribution. In this work, this kind of nanobainite structure was called the dualphase and multiscale structure. When the austempering temperature reached 400 °C, the microstructure coarsened considerably.

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—

—

Fig. 3. (a) The average length of bainite plates (LαBF), average size of blocky retained austenite (Dγb),and average —

—

thicknesses of bainite plates ( t αBF) and filmy retained austenite ( t γf) as a function of the austempering temperature. Distributions of (b) the length of bainite plates (LαBF) and (c) size of blocky retained austenite (Dγb) in the designed steel austempered at different temperatures.

The carbon content (Cγ) and the martensite start temperature (Msγ) of retained austenite were shown in Fig. 4. With increasing the austempering temperature from 260 °C to 400 °C, Cγ decreased monotonically. On the basis of the Cγ value, the Msγ temperatures were estimated by Equation (1) [21], as shown in Fig. 4b. All Msγ temperatures of retained austenite formed at different austempering temperatures were lower than room-temperature (25 °C). This result proved that retained austenite did not transform to martensite during cooling from austempering temperatures to room temperature. This is because once bainitic ferrite forms through shear mechanism, the excess carbon atoms in bainitic ferrite are quickly rejected into the surrounding austenite, reducing the Msγ temperatures of the retained austenite [28–29]. Msγ (°C) = 539 − 423 C − 30.4 Mn − 12.1 Cr − 7.5 Si 6

(1)

where the symbol of each element corresponds to its content (in wt.%) in retained austenite.

Fig. 4. (a) Carbon content and (b) martensite start temperature of retained austenite as a function of austempering temperature.

Meanwhile, the carbon in bainitic ferrite remained supersaturated and formed carbon clusters, based on the analysis by Caballero et al. [30, 31]. This was clearly confirmed by the carbon atom maps of the steel austempered at 350 °C, as shown in Fig. 5. Moreover, the amount of carbon clusters in long nanobainite plates (i.e. the C-poor area in Fig. 5a) was considerably less than that in short nanobainite plates (i.e. the C-poor area in Fig. 5b). According to the concentration profiles, the carbon content of long nanobainite plates was ~0.13 at.% (~0.03 wt.%), close to the paraequilibrium solubility levels (0.0218 wt.%). The carbon content of short nanobainite plates was considerably supersaturated with ~0.43 at.% (~0.10 wt.%) C. Carbon distribution in large blocky retained austenite adjoining long nanobainite plates was more inhomogeneous than that in small blocky retained austenite neighboring short ones. This result will be explained in the discussion section.

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Fig. 5. Carbon atom maps (left) and the corresponding concentration profiles across the αBF/γb interfaces along the arrow z (right) in the regions of (a) long nanobainite plates and (b) short nanobainite plates obtained by a site-specific FIB lift-out technique in the designed steel austempered at 350 °C. Error bars represent the statistical error as calculated.

The engineering stress–strain curves and the summarized mechanical properties of the designed steel austempered at different temperatures are presented in Fig. 6a. With increasing austempering temperature from 260 °C to 400 °C, the tensile strength, yield strength, and area reduction decreased gradually. By contrast, the uniform elongation, total elongation, product of strength and elongation, and impact energy increased with increasing austempering temperature from 260 °C to 350 °C. Then, they decreased gradually with increasing austempering temperature from 350 °C to 400 °C. The uniform elongation of the sample austempered at 350 °C (the nose-tip temperature of the TTT diagram) reached a maximum value of 28.2%, which was almost equivalent to the total elongation of 31.5%. This result indicated that an excellent uniform deformation capability was obtained. At the same time, the high tensile strength of 1552 MPa and impact energy of 85 J were achieved in the sample austempered at 350 °C. Moreover, the mechanical properties of the sample austempered at 370 °C (approximately the nose-tip temperature of the TTT diagram) 8

were also excellent, as shown in Fig. 6a. Fig. 6b shown the comparison of tensile properties between the designed steel austempered at 350 °C and 370 °C and other advanced steels. The symbols of the designed steel were located in the upper right of the tensile property map, indicating that the designed steel possessed a better combination of strength and ductility than those advanced steels reported before. In addition, the designed steel had a high impact energy and low raw material cost (owing to the absence of expensive alloy elements, e.g. Ni, Mo and Nb), so it had a wide application potential as an alternative for traditional high-strength steels.

Fig. 6. (a) Engineering stress–strain curves and mechanical properties of the designed steel austempered at different temperatures. (b) Comparison of tensile properties between the designed steel and other advanced steels, including martensitic [32], nanobainitic [10, 17, 33], TRIP [34–36], TWIP [37, 38], quenching and partitioning [39], and high specific strength steels [40]. YS stands for yield strength, TS for tensile strength, UE for uniform elongation, TE for total elongation, RA for area reduction, AKU for Charpy impact energy, and PSE for product of strength and elongation.

Valuable information about the fracture process can be obtained directly by fractographic observation. Dimples and quasi-cleavage facets are indicators of ductile and brittle features, respectively. The photographs and SEM fractographs of the tensile samples austempered at 260 °C, 350 °C, and 400 °C were shown in Fig. 7. Ductile dimples, formed by microvoid coalescence, were predominant in the fractures of the samples austempered at 260 °C and 350 °C (Fig. 7c, e). According to the cup and cone fracture (Fig. 7b, d), the major fracture micromechanism of the samples austempered at 260 °C and 350 °C was ductile nucleation and growth of microvoids, and the final shearing of the samples. The depth of dimples in the fracture of the sample austempered 350 °C (Fig. 7e) was obviously greater than that in the sample austempered at 260 °C (Fig. 7c), indicating that the former had higher ductility. In the samples austempered at 350 °C and 400 °C, 9

hardly any necking was observed (Fig. 7a), suggesting all of the elongation was uniform. The sample austempered at 400 °C exhibited large and flat quasi-cleavage facets (Fig. 7g), which revealed the brittle features.

Fig. 7. (a) Tensile samples before and after deformation. SEM fractographs of the tensile samples austempered at (b, c) 260 °C, (d, e) 350 °C, and (f, g) 400 °C.

The variation in the instantaneous work hardening index (n*) with the true strain (εu) was shown in Fig. 8. The instantaneous work hardening index of the tensile sample austempered at 260 °C continuously decreased and rapidly reached true strain at a low-strain regime. This result indicated that plastic instability and necking occurred prematurely. The tensile sample austempered at 400 °C fractured prematurely before n* equal to εu. By contrast, sustained high work hardening and excellent uniform deformation capability occurred at the tensile sample austempered at 350 °C.

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Fig. 8. Instantaneous work hardening index (n*) as functions of the true strain (εu) during tensile deformation. Equation n* = εu is the instability criterion.

This may be related to the deformation-induced martensite transformation from retained austenite during tensile deformation [41, 42]. This transformation implies a relaxation of the local stress concentration and stress redistribution during plastic deformation, which delays the necking up to the high strain region. Thus, the uniform elongation increases and better ductility is achieved. The volume fraction of retained austenite (Vγ), as shown in Fig. 9a, increased from ~10.4 vol.% to ~33.2 vol.% with increasing the austempering temperature from 260 °C to 400 °C. Fig. 9b shown the decrease in the volume fraction of the retained austenite (Vγ) during tensile deformation. The retained austenite in the samples austempered at 260 °C and 350 °C was completely transformed to martensite. For the former, the contribution of the TRIP effect on the mechanical properties was limited by low V (~10.4 vol.%). Less TRIP effect led to the early onset of necking. For the sample austempered at 400 °C, although Vγ was very high (~33.2 vol.%), the tensile sample fractured prematurely. This was because the overly coarse hard phase (martensite) from the coarse retained austenite accelerated the initiation and propagation of microcracks. According to the XRD result, there was still 8.8 vol.% of untransformed retained austenite close to the fracture surface.

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Fig. 9. (a) Volume fraction of retained austenite (Vγ) as a function of austempering temperature. (b) The decrease in the volume fraction of the retained austenite (Vγ) as a function of the true strain (εu) during tensile deformation.

Because martensite was not formed during the cooling after austempering (Fig. 4b), all of the martensite in the tensile samples was due to deformation induced transformation during tensile deformation. To observing martensite transformation during tensile deformation in the samples austempered at 350 °C, the microstructures in custom-made deformation stage were shown by TEM (Fig. 10). Under a true strain of 0.05 (Fig. 10a), the large retained austenite blocks partly transformed to α martensite, as was confirmed by the diffraction pattern in the inset. Under a true strain of 0.10 (Fig. 10b), multiple small martensitic grains were formed in one large retained austenite block. This result will be discussed in the discussion section. Close to the fracture surface, filmy retained austenite also transformed to martensite (Fig. 10c), which was consistent with the analysis in Ref. [2]. They reported that morphology was an important factor to be considered on the mechanical stability of retained austenite. Thin films of retained austenite were the most stable, so that it transformed usually in the later stage of tensile deformation. In addition, the TEM result close to the fracture surface (Fig. 10d) shown that long nanobainite sheaves undergone significant deformation without the appearance of microcracks, which will be also discussed later.

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Fig. 10. TEM images of the designed steel austempered at 350 °C after the tensile test: (a) at true strain of 0.05, (b) at true strain of 0.10, and (c, d) close to the fracture surface. Insets in a and c are the selected area electron diffraction (SAED) patterns of martensite within the red dashed circles. Yellow circles in the SAED patterns represent the diffraction beam selected for the dark-field image. αM′ represents α martensite.

4. Discussion Ravi et al. [43] show that bainite plates are initially nucleate at the prior austenite grain boundaries (named grain-boundary nucleation), which leads to an increase in the number density of nucleation sites. Afterwards, bainite plates are nucleate at these newly created nucleation sites (named autocatalytic nucleation). The nucleation driving force is the greatest possible reduction in free energy that can be achieved during formation of a ferrite nucleus such that the composition of surrounding austenite matrix remains unaffected. It is calculated by MUCG83 thermodynamic model [44, 45], as shown in Fig. 11. With increasing austempering temperature, the nucleation driving force decreases, leading to low nucleation rate, as reported by others [2]. Therefore, at high austempering temperature, less bainite plates are nucleated at the prior austenite grain boundaries. They, then, seldom interfere with each other during growth and can traverse the prior austenite grains, which is determined by the shear mechanism of bainite transformation [46]. According to the microstructure images (Fig. 2b, c), when the austempering temperature reaches 350 °C and above, bainite plates nucleated at the prior austenite grain boundaries can, indeed, traverse the prior austenite grains. At low austempering temperature, the nucleation driving force is large and therefore the nucleation rate is high. In this work, when the austempering temperature is as low as 260 °C, a large number of bainite plates are nucleate at the prior austenite grain boundaries. They interfere with each other during growth and therefore their length is short (Fig. 2a). The growth of secondary bainite plates (nucleated via autocatalytic nucleation) is blocked by the preformed bainite plates (nucleated via grain-boundary nucleation), resulting in a reduction in the length of secondary bainite plates.

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Fig. 11. Calculated nucleation driving force as a function of austempering temperature in steels with different Si contents using the MUCG83 code. The Si content of the designed steel (2.6 wt.%) slightly exceeds the content range of Si (0–2.5 wt.%) specified in the MUCG83 code. Therefore, we calculated curves of steels with different Si contents (0–2.5 wt.%). The corresponding curves of steels with different Si contents are almost the same, indicating that the model can be applied to the designed steel.

The arrangements of bainite plates and retained austenite in a given prior austenite grain at different stages of bainite transformation are shown schematically in Fig. 12. αBF-g and αBF-a represent the bainite plates nucleated via grain-boundary nucleation and autocatalytic nucleation, respectively. γut represents the untransformed austenite at the early bainite transformation stage. PAGB is the prior austenite grain boundaries. Long nanobainite sheaves at 350 °C divide the prior austenite grains to separated γut and reject carbon atoms into γut (Fig. 12b). From the viewpoint of the diffusionless school [47], the bainite transformation stops when the carbon-enriched retained austenite reaches approximately a composition defined by the T0 curve. The T0 curve defines the locus of all points on a temperature versus carbon content where austenite and ferrite of the same chemical composition have identical free energy [48]. Some γut reaches the composition defined by the T0 curve, so it is retained at room temperature, forming large blocky retained austenite and adjoining long αBF-g (Figs. 2b and 12b). Meanwhile, other γut below the composition defined by the T0 curve is divided to small blocky retained austenite by αBF-a. Therefore, The small blocky retained austenite adjoins short αBF-a (Figs. 2b and 12b). Careful Calculation [17] in the 0.4C–2.0Mn–1.7Si–0.4Cr (wt.%) steel shows that it takes more than 4 h to accomplish carbon homogenization in the blocky austenite (with the halfwidth of 2.5 µm) during austempering at 360 °C, due to the low diffusion coefficient of carbon in austenite. In this work, therefore, the carbon distribution in large blocky retained austenite with the halfwidth above 2.5 µm (Fig. 3c) is inhomogeneous (Fig. 5a) under austempered at 350 °C for 2h. In the early stage of bainite transformation at 350 °C, a carbon concentration gradient occurs in large blocky retained austenite. The carbon content of the large blocky retained austenite near the interfaces is higher than that of the interior. With increasing austempering time, the carbon content in the large blocky retained austenite near the interfaces slightly decreases due to diffusion into the interior [17, 49]. the carbon in the surrounding long αBF-g diffuses into large blocky retained austenite to compensate for the carbon loss of the large blocky retained austenite near the interfaces, i.e., uphill diffusion, because the chemical potential of carbon in austenite is higher than that in ferrite. Consequently, the carbon content of long nanobainite plates is as low as ~0.13 at.% (Fig. 14

5a).

Fig. 12. Schematics of bainitic ferrite and retained austenite arrangement in a given prior austenite grain at the early (up) and last stages (down) of bainite transformation in the steel austempered at (a) 260 °C, (b) 350 °C, and (c) 400 °C. αBF-g and αBF-a represent the bainite plates nucleated via grain-boundary nucleation and autocatalytic nucleation, respectively. γut represents the untransformed austenite at the early bainite transformation stage. PAGB is the prior austenite grain boundaries.

As suggested by previous studies [22–24], the mechanical stability of austenite (i.e. its capability to transform to martensite under strain) is related to its morphology, grain size, and carbon content. Large blocky retained austenite is the most unstable [2] to transform by TRIP effect, and several are the reasons. First, because of the absence of the constraint to transformation exerted by the surrounding ferritic plates. Second, because large blocky retained austenite contains a large number of potential nucleation sites for the transformation to martensite therefore requiring lower driving force for martensite nucleation. And finally, because of its lower carbon concentration. In this work, the deformation-induced martensite transformation from large blocky retained austenite occurs in the early stage of tensile deformation (Fig. 10a, b), consistent with the observations presented in Refs. [25, 26]. Furthermore, it is found that the carbon distribution in large blocky retained austenite is the most inhomogeneous (Fig. 5a). Large blocky retained austenite partly transforms to martensite at low strain (Fig. 10a) and multiple small martensitic grains can be formed in one large retained austenite block as the strain increases (Fig. 10b). This is because the chemical composition is an 15

important factor controlling the mechanical stability of austenite. Elements such as C, Mn and Si [50] significantly enhance the mechanical stability of austenite, among them C is the element that exhibits the strongest influence. It is possible that the low-C part of large blocky retained austenite tends to transform to martensite earlier than the high-C part. Based on the analysis in Ref. [51], in dual phase ferrite–martensite steels, ductile proeutectoid ferrite in the vicinity of martensite has stress relief effect just ahead of the current microcrack. This is because ductile proeutectoid ferrite deforms as a result of plastic constraining of martensite and thus excessive load is not applied to the proeutectoid ferrite during plastic deformation. Although bainitic ferrite in bainitic steels is plate-like different from equiaxed proeutectoid ferrite in dual phase ferrite–martensite steels, bainitic ferrite is also a ferrite. Bainitic ferrite will be extremely ductile when its carbon content is close to the paraequilibrium solubility levels of 0.0218 wt.% C, such as long nanobainite plates with ~0.03 wt.% C (Fig. 5a). At the last stage of tensile deformation, almost all of the retained austenite has been transformed to martensite (Fig. 9b). Due to plastic constraining of deformation-induced martensite, large plastic deformation occurs to long nanobainite sheaves with ductile bainite plates (Fig. 10d), which can relax local stress concentration of crack tip, therefore blunting the crack tip and retarding the crack propagation. To better understand the evolution of dualphase and multiscale carbide-free nanobainite structure during tensile straining, a schematic model is built, as shown in Fig. 13. This unique structure before tensile deformation is shown in Fig. 13a. At the early stage of tensile deformation, large blocky retained austenite is partly and selectively transformed to martensite grains (Fig. 13b, c). With increasing strain, small blocky retained austenite and filmy retained austenite are also transformed gradually to martensite grains (Fig. 13c, d). The sustained martensitic transformation provides the extra work hardening (Fig. 8) and therefore ensures the high uniform elongation (Fig. 6a). During tensile straining, ductile long nanobainite plates (with ~0.03 wt.% C) undergo significant deformation because of plastic constraining of martensite. At the last stage of tensile deformation, the propagation of microcracks is effectively suppressed by ductile long nanobainite plates (Fig. 13e).

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Fig. 13. Evolution of dualphase and multiscale carbide-free nanobainite structure during tensile straining.

The strength and ductility of materials are two conflicting mechanical properties. Improving strength or ductility often leads to the degradation of the other property, which is known as the strength-ductility trade-off [52]. The pursuit of improved strength and ductility has always been the goal of material researchers. In principle, PSE (i.e., the product of strength and elongation) can be used to characterize the combination of strength and ductility of materials. The designed steel austempered around the nose-tip temperature of the TTT diagram exhibits a better combination of strength and elongation, e.g. PSE (Fig. 6a). Similar observations are also shown in Refs. [22, 53]. They report that the PSE of C-Si-Mn-Cr steels austempered near the nose-tip temperature of the TTT diagram was higher than that at other temperatures. These results should not be accidental. Future research should be directed toward the bainitic microstructure and mechanical properties in other materials obtained by austempering around the nose-tip temperature of the TTT diagram. 5. Conclusions In the current study, a dualphase and multiscale nanobainite structure was obtained in a high-C high-Si steel, through austempering around the nose-tip temperature of the TTT diagram. The main results can be summarized as follows. 1. This dualphase and multiscale nanobainite structure offered an excellent combination of tensile strength ~1550 MPa, uniform elongation of ~30%, and impact energy ~80 J. 2. In this nanobainite structure, multiple lengths of nanobainite plates and various sizes and shapes of retained austenite were obtained. Large blocks of retained austenite were neighbored by 17

long bainite plates, and small blocks were adjoined by short bainite plates. Due to the low diffusion coefficient of carbon in austenite, large blocks of retained austenite with the halfwidth of above 2.5 µm typically had an inhomogeneous distribution of carbon atoms. The carbon content of long nanobainite plates was as low as ~0.13 at.%, because the carbon in long nanobainite plates diffused into large blocky retained austenite to compensate for the carbon loss of the austenite near the interfaces. 3. Affected by their inhomogeneous carbon distribution, large blocks of retained austenite transformed partly and gradually to multiple martensite grains as strain increases. This was because C exhibits the strongest influence in mechanical stability of austenite. 4. Long nanobainite plates with the carbon content close to the paraequilibrium solubility levels are extremely ductile. The large plastic deformation of the long nanobainite sheaves with ductile bainite plates can relaxed local stress concentration of crack tip, therefore blunting the crack tip and retarding the crack propagation. Acknowledgements This work was supported by the National Natural Science Foundation of China [Nos. 51831008 and 51471146] and the Innovation Funding Project for Postgraduate of Heibei Province [No. CXZZBS2018050]. References [1] A. Leiro, E. Vuorinen, K.G. Sundin, B. Prakash, T. Sourmail, V. Smanio, F.G. Caballero, C. Garcia-Mateo, Roberto Elvira, Wear of nano-structured carbide-free bainitic steels under dry rolling–sliding conditions, Wear 298–299 (2013) 42–47. [2] C. Garcia-Mateo, F.G. Caballero, T. Sourmail, M. Kuntz, J. Cornide, V. Smanio, R. Elvira, Tensile behaviour of a nanocrystalline bainitic steel containing 3 wt% silicon, Mater. Sci. Eng. A 549 (2012) 185–192. [3] Y.H. Wang, F.C. Zhang, T.S. Wang, A novel bainitic steel comparable to maraging steel in mechanical properties, Scripta Mater. 68 (2013) 763–766. [4] K. Hase, C. Garcia-Mateo, H.K.D.H. Bhadeshia, Bimodal size-distribution of bainite plates, Mater. Sci. Eng. A 438–440 (2006) 145–148. [5] W. Solano-Alvarez, E.J. Pickering, H.K.D.H. Bhadeshia, Degradation of nanostructured bainitic steel under rolling contact fatigue, Mater. Sci. Eng. A 617 (2014) 156–164. [6] W. Solano-Alvarez, E.J. Pickering, M.J. Peet, K.L. Moore, J. Jaiswal, A. Bevan, H.K.D.H. Bhadeshia, Soft novel form of white-etching matter and ductile failure of carbidefree bainitic steels under rolling contact stresses, Acta Mater. 121 (2016) 215–226. [7] K.K. Wang, Z.L. Tan, G.H. Gao, X.L. Gui, R.DK. Misra, B.Z. Bai, Ultrahigh strength-toughness combination in Bainitic rail steel: The determining role of austenite stability during tempering, Mater. Sci. Eng. A 662 (2016) 162–168. [8] M.N. Yoozbashi, S. Yazdani, Mechanical properties of nanostructured, low temperature bainitic steel designed using a thermodynamic model, Mater. Sci. Eng. A 527 (2010) 3200–3205. [9] G.H. Gao, H. Zhang, X.L. Gui, P. Luo, Z.L. Tan, B.Z. Bai, Enhanced ductility and toughness in an ultrahigh-strength Mn-Si-Cr-C steel: The great potential of ultrafine filmy retained austenite, Acta Mater. 76 (2014) 425–433. [10] Y. Wang, K. Zhang, Z.H. Guo, N.L. Chen, Y.H. Rong, A new effect of retained austenite on ductility enhancement in high strength bainitic steel, Mater. Sci. Eng. A 522 (2012) 288–294 . [11] K. Zhang, M.H. Zhang, Z.H. Guo, N.L. Chen, Y.H. Rong, A new effect of retained austenite on ductility enhancement in high-strength quenching-partitioning-tempering martensitic steel, Mater. Sci. Eng. A 528 (2011) 8486–8491. [12] Z.N. Yang, Y.L. Ji, F.C. Zhang, M. Zhang, B. Nawaz, C.L. Zheng, Microstructural evolution and 18

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