Figure 10 40 Shows the Continuous Cooling

1 Introduction

Over the last several decades, nonquenched and tempered steels, obtained by alloy design and thermomechanically controlled processing, have been applied gradually in the automotive field, such as crankshafts, spindles and connecting rods, and in other fields, owing to the advantages of low cost, good performance and being environmental friendly. Along with the demand for higher strength steels in automotive parts making, traditional ferrite and pearlite-quenched and tempered steels are no longer to meet the requirements [1,2]. Hence, the bainite and martensite microstructures are considered to be developed in a nonquenched and tempered way. As the bainitic steels are possibly superior to martensitic steels in toughness and plasticity, the research on nonquenched and tempered bainitic steels has received a lot of attention recently.

Clearly, the morphology of bainite has an apparent influence on mechanical properties. As Caballero and Bhadeshia [3,4,5] claimed, the lath bainite, especially carbide-free lath bainite, exhibited a higher impact toughness value than that of granular bainite due to the easy nucleation of microcracks in the martensite/austenite islands. Typically, the carbide-free bainite is generally obtained by adding Si or Al to hinder the precipitation of brittle carbides such as cementite to preserve the carbon-rich retained austenite and to achieve better mechanical properties [6], which benefit the future wider applications of the carbide-free bainite steels [7,8,9,10]. The technology widely applied to obtain the carbide-free bainite microstructure is to use the austempering heat treatment method in the bainitic temperature range generally [11,12,13], which is conducive to low production efficiency and high energy consumption. Therefore, it is necessary to combine the environment friendly process and excellent microstructure to facilitate the industrial production effectively.

Hence, in this study, the chemical composition was designed to promote the bainite transformation kinetics and to retard the carbide precipitation, to obtain the carbide-free bainite under nonquenched and tempered conditions, such as air cooling or controlled cooling after hot deformation. More specifically, C was added to ensure the high strength by solid solution strengthening. Si was added to promote carbon partitioning in the retained austenite to make it more stable [14,15,16,17]. Mn reduces the bainitic transformation temperature by significantly reducing the ∆G ^γ→α , explored by Shah et al. [16]. Furthermore, the addition of Mo effectively suppresses the ferrite transformation and decreases the transformation temperature of bainite, which helps to obtain ultrafine carbide-free bainite [17,18,19,20,21]. Besides, a trace amount of B inhibits the nucleation of ferrite and promotes the formation of bainite [22,23] in the experimental steel. In addition, Nb improves the strength and toughness by the grain refining and the precipitation strengthening of microalloying carbonitride [24]. Ti is used to refrain from the formation of BN.

It is generally known that both the chemical composition and the controlled rolling and cooling processes affect the microstructure and mechanical properties significantly. It is unclear how the cooling rate affects the microstructure, including the evolution of retained austenite and mechanical properties of the designed medium-carbon experimental steels in this study. The continuous cooling process was conducted on the thermal simulator to figure out the most suitable cooling rate for obtaining high strength and toughness microstructure in this medium-carbon alloy system and to verify that based on the semi-industry experimental work. Overall, the effects of the cooling rate on the transformation microstructures and mechanical properties of carbide-free bainitic steels are studied to provide references for the application of this new type of carbide-free bainitic steels in terms of the alloy composition and process designs.

2 Materials and experimental procedure

2.1 Materials preparation

The investigation was carried out on self-designed steel prepared in an intermediate frequency vacuum induction furnace. The chemical composition is listed in Table 1. A total of 0.25 wt% C was added to ensure enough strength by solid solution strengthening. High contents of Si and Mn were designed to inhibit the precipitation of brittle carbides, especially cementite (Fe₃C). In addition, the mole fractions of the phases at equilibrium were calculated by Thermo-Calc software package and TCFE8 database to design the proper contents of Si and Mn. As shown in Figure 1, the types of secondary phases of the experimental steel and the steel containing 0.3 wt% Si and 1.0 wt% Mn (general content) did not differ significantly, except for the obvious difference in the precipitation behavior of Fe₃C. The start and finish precipitation temperatures of Fe₃C are 728 and 535°C for the experimental steel and 730 and 450°C for the 0.3Si–1.0Mn steel, respectively. Hence, from a thermodynamic point of view, Fe₃C is less likely to form in steel after adding more Si and Mn. Meanwhile, the appropriate control of kinetic conditions will be further beneficial to inhibit the precipitation of Fe₃C and the formation of bainite in the experimental steels containing high contents of Si and Mn to obtain carbide-free bainite finally.

Table 1

The chemical composition of the experimental steel (wt%)

C	Si	Mn	Mo	Nb	Ti	Cr	Cu	B	N	Fe
0.25	1.54	1.84	0.23	0.032	0.019	0.51	0.14	0.0025	0.0028	Bal.

Figure 1

The mole fractions of secondary phases calculated by Thermo-Calc: (a) 0.3Si–1.0Mn steel and (b) experimental steel.

The dynamic continuous cooling process was carried out using the Gleeble-3800 thermal simulator. The dilatometry specimens were machined as shown in Figure 2. As shown in Figure 3, the specimens were heated and austenitized at 1,150°C for 300 s first, cooled to 1,050°C at the rate of 10°C/s, then compressed twice with strains of 0.4 and 0.6 at 1,050 and 900°C, respectively, at a strain rate of 10 s⁻¹ and followed by controlled cooling to room temperature at the rate of 0.05–10°C/s. The specimens for microstructural investigation were sectioned parallel to the rolling direction (RD) after hot deformation.

Figure 2

Schematic of the sample for the thermo-simulation experiment on Gleeble3800.

Figure 3

Schematic illustration of the thermo-mechanical simulation schedule.

2.2 Microstructural characteristics

The deformed metallographic specimens were ground, polished and etched with 4 vol% nital before the observation by optical microscopy (OM, MX6R) and field emission scanning electron microscopy (FESEM, JSM-6701F). The phase composition of the microstructure at different cooling rates, especially the amount of retained austenite, was determined by X-ray diffractometer (XRD, Bruck™-D8 Advance) using Cu radiation by the continuous mode. The scanning range in 2θ was 47°–94° at a rate of 0.02°/s, covering the diffractions of (200) _γ , (200) _α , (220) _γ , (211) _α and (311) _γ . The amount of retained austenite was calculated from an average based on the diffraction peaks according to equation (1) [25].

(1) V i = 1 1 + G ( I α / I γ )

where V _i is the volume fraction of austenite for each peak and I α and I γ are the integrated intensities of ferrite and austenite peaks, respectively. G-values for each peak are as follows: 2.5 for I α ( 200 ) I γ ( 200 ) , 1.38 for I α ( 200 ) I γ ( 220 ) , 2.02 for I α ( 200 ) I γ ( 311 ) , 1.19 for I α ( 211 ) I γ ( 200 ) , 0.06 for I α ( 211 ) I γ ( 220 ) and 0.96 for I α ( 211 ) I γ ( 311 ) .

Besides, the microstructure units were further analyzed by the transmission electron microscopy (TEM) and electron back-scattering diffraction (EBSD) techniques. The TEM studies were conducted on 3-mm-diameter thin foils by FEI Tecnai G2 F30 S-TWIN TEM at an accelerating voltage of 300 kV. The EBSD maps were carried out on Auger Electron Spectrometer (AES, PHI710) equipped with an EBSD system and acquired using TSL OIM Analysis 7 software. The scanning step length was set as 0.04 μm, and the scanning area was 20 × 20 μm.

2.3 Mechanical properties

The hardness values of the samples obtained at different cooling rates were measured by the 430SH Vickers hardness tester under the load weight of 500 g. To ensure the reliability of the results, 10 points were measured, and hardness values were averaged for each sample.

The nanoindentation measurements were carried out on Nano Indenter XP by the quasi-static method based on the Oliver and Pharr method [26]. The indentation depth was 1,000 nm, and Poison's ratio was 0.25. The nanohardness and Young's modulus were tested and averaged by 10 points. In addition, the curves of displacement and load under different cooling rates were plotted. Compared with Vickers hardness, the test points of nanoindentation can be guaranteed inside a grain, generally, which is more conducive to the determination of micromechanical properties of the fine-grain microstructure.

Furthermore, the dislocation densities at different cooling rates were analyzed by XRD on Bruck™-D8 Advance diffractometer using Cu radiation by the step mode. The X-ray tube was operated at 40 kV and 40 mA, and the measurements were performed in the range of 35°(2θ) to 145°(2θ) with a step size of 0.02° and a counting time of 0.6 s, covering the diffractions of (110) _α , (200) _α , (211) _α , (220) _α , (310) _α and (222) _α . Then, the dislocation density was calculated by the following modified Williamson–Hall (MWH) equation [27,28,29,30]:

(2) Δ K = 0.9 D + π M b 2 ρ 1 / 2 2 K C 1 / 2 + O ( K 2 C )

where ΔK is the full-width-half-maximum (FWHM), D is the average particle size, C is the dislocation contrast factor, b is the magnitude of the Burgers vector and equals 0.245 and ρ is the dislocation density. K is the magnitude of diffraction vector and equals 2*sinθ/λ, where θ and λ are the diffraction angle and the wavelength, respectively. In the present test, λ equals CuKα, which is 0.15418 nm. M and O are the constants depending on the effective cut-off radius of the dislocations. M is considered to be 1.4 by using the value given by HajyAkbary et al. [30], and O is neglected due to its high order term.

As Ungár and Borbély [31] indicated, the average contrast factor of dislocations C ¯ is used instead of C, as defined by equation (3) for the specific (hkl) diffraction:

(3) C ¯ = C ¯ h 00 ( 1 − q H 2 )

where q is a constant obtained experimentally, and H ² is defined in equation (4):

(4) H 2 = h 2 k 2 + k 2 l 2 + l 2 h 2 ( h 2 + k 2 + l 2 )

Combining equation (2) and (4), the value of ρ can be calculated from equation (5) causing the intercept, and the slope of the line can be obtained with Origin software conveniently.

(5) ( Δ K − 0.9 D ) 2 K 2 = ( π M b 2 ρ 1 / 2 2 ) 2 0.285 ( 1 − q H 2 )

Finally, the experimental steel blank (50 × 25 × 150 mm) at 0.5°C/s was produced by the semi-industry process as shown in Figure 4. In particular, the steel blank was cooled slowly by burying in the sand. The cooling time from 900 to 200°C was 22.5 min, and the average cooling rate was approximately equal to 0.5°C/s. The strength was measured by a standard tensile tested at room temperature using a universal testing machine (CMT4105). The impact specimens with v-notch parallel were tested at room temperature using a pendulum impact testing machine (ZBC2452-B). The fracture surfaces were observed by JSM-6480LV SEM.

Figure 4

Schematic illustration of the semi-industry process.

3 Results

3.1 Transformation microstructures

The effect of cooling rates on the microstructure of experimental steel is shown in Figures 5 and 6, obtained by optical microscopy and scanning electron microscopy, respectively. Figures 5(a) and 6(a) show that the microstructure was composed of very small amount of polygonal ferrite (PF) and granular bainite (GB) at the cooling rate of 0.05°C/s. When the cooling rate increased to 0.1°C/s, the size of PF was reduced to less than 10 µm, and the microstructure was almost composed of GB, as shown in Figures 5(b) and 6(b). The microstructure changed from GB to lath bainite (LB), and the volume fraction of LB increased with the increase of the cooling rate from 0.3 to 1°C/s, as shown in Figures 5(c)–(e) and 6(c)–(e). Then, the martensite transformation occurred at the cooling rate of 1°C/s, and the microstructure was almost martensite (M) at the cooling rate exceeding 2°C/s as shown in Figures 5(e) and (f) and 6(e) and (f).

Figure 5

Representative OM micrographs of experimental steel with different cooling rates: (a) 0.05°C/s; (b) 0.1°C/s; (c) 0.3°C/s; (d) 0.5°C/s; (e) 1°C/s; (f) 2°C/s; (g) 5°C/s and (h) 10°C/s.

Figure 6

Representative FESEM micrographs of experimental steel with different cooling rates: (a) 0.05°C/s; (b) 0.1°C/s; (c) 0.3°C/s; (d) 0.5°C/s; (e) 1°C/s and (f) 2°C/s.

Besides, the Vickers hardness played an important role in identifying the types of microstructures. The variations in Vickers macrohardness as the function of the cooling rate for the investigated samples are shown in Figure 7. As has been widely known, with the increasing cooling rate and the microstructure evolving from GB to M, the Vickers macrohardness values turn out to gradually upward trends, which change from 370 to 522 HV.

Figure 7

Vickers macrohardness variation under different cooling rates.

3.2 Transformation behavior

Transformation temperatures were determined from slope change in the curve that denoted the length changes of the dilatometry samples, as shown in Figure 8(a). Under the cooling rate of 0.3°C/s, the start transformation temperature of bainite (B_s) is 500°C and the finish transformation temperature of bainite (B_f) is 331°C. The results of B_s and B_f under other four different cooling rates are presented in Table 2. Both the B_s and B_f points decreased with the increase in the cooling rate. Specifically, the B_s point was decreased from 505 to 429°C, and the B_f point was decreased from 351 to 288°C. Besides, within the cooling rate from 0.1 to 2°C/s, different amounts of lath bainite can be obtained, which are the focus of this study. Combined with expansion curves and microstructures, the dynamic continuous cooling transformation (CCT) diagram of the experimental steel was obtained and presented in Figure 8(b).

Figure 8

(a) Expansion curve and transformation temperatures of the sample under the cooling rate of 0.3°C/s and (b) dynamic continuous cooling transformation curves of the experimental steel.

Table 2

B _s and B _f transformation temperatures of the samples

Transformation	Cooling rate (°C−1 s−1)
Temperature/°C	0.05	0.1	0.3	0.5	1
B s	505	500	500	481	429
B f	351	345	331	294	288

3.3 Mechanical properties

The results of nanohardness and Young's modulus at three different slow cooling rates are demonstrated in Figure 9. From Figure 9(a), the nanoindentation point measurement was conducted in one LB or one GB, which could avoid the strengthening effect of grain boundaries. Figure 9(b) shows the depth of nanoindentation decreased with the increase of cooling rate even under the larger load. Finally, Figure 9(c) shows the results of nanohardness and Young's modulus. At the cooling rate of 0.05°C/s, the average values of nanohardness and Young's modulus were 4.26 ± 0.46 GPa and 214.7 ± 12.4 GPa, respectively. When the cooling rate increased to 0.5°C/s, the corresponding average values were 4.76 ± 0.39 and 216.8 ± 8.1 GPa. Then, the cooling rate increased to 1°C/s, and the average values were 5.36 ± 0.9 GPa and 210.9 ± 19.8 GPa.

Figure 9

Characteristics of nanoindentation measurement: (a) SEM morphology of nanoindentation in lath and granular bainites at 0.5°C/s cooling rate; (b) load–displacement curves; (c) variation of nanohardness and Young's modulus of the microstructure at different cooling rates.

With consideration of the results of microstructure characterization, the cooling rate of 0.5°C/s was selected to prepare the samples for the mechanical properties test, as shown in Figure 4. Table 3 presents the yield strength (R _el), tensile strength (R _m) and Charpy impact energy (A _kv) at room temperature of the experimental steel, which are all average values of three standard tensile and impact specimens.

Table 3

The mechanical properties of the experimental steel

R el/MPa	R m/MPa	A/%	A kv/J
847	1,298	18	38

4 Discussion

As has been known for the transformation evolution, the morphology of bainite will change from granular to lath as the cooling rate increases. The morphologies of GB and LB were observed by TEM, as shown in Figure 10. The bright-field micrograph of the GB at the cooling rate of 0.05°C/s and that of the LB at the cooling rate of 0.5°C/s are shown in Figure 10(a) and (d), respectively. Figure 10(b) and (c) show the SAD patterns of ferrite and austenite of GB, and Figure 10(g) and (h) show those of LB. Meanwhile, Figure 10(e) and (f) present the magnified bright-field and dark-field micrographs of LB. The results show that with the increase of cooling rate, the microstructure was transformed from granular bainitic ferrite with a high dislocation density and block M/A islands to parallel bainitic ferrite laths with a higher dislocation density and retained austenite films distributed between bainitic ferrite laths. The widths of the LB ferrite lath and the retained austenite film shown in Figure 10(d) measured by IPP (Image Pro-Plus software) are 268.1 and 89.5 nm, respectively, which means that the microstructure of the LB is clearly finer than that of the GB.

Figure 10

Representative TEM micrographs and selected area diffraction patterns for experimental steel under different cooling rates: (a) bright-field micrograph of the GB at the cooling rate of 0.05°C/s; (b) SAD pattern for zone 1 in (a); (c) SAD pattern for zone 2 in (a); (d) bright-field micrograph of the LB phase at the cooling rate of 0.5°C/s; (e) magnified bright-field micrograph of (d); (f) dark-field micrograph of LB; (g) SAD pattern for zone 3 in (e) and (h) SAD pattern for zone 4 in (e).

To recognize the phase constituent of the samples in the range of the cooling rates for the formation of bainite, the amount of the retained austenite was determined by XRD, as shown in Figure 11. Figure 11(a) shows that the microstructures of all the investigated samples consist of ferrite and austenite phases, and there is no obvious carbide diffraction peak, which further shows that the experimental steels are carbide-free bainitic steels. The relatively high amount of retained austenite is due to the strong retardation of carbide precipitation by Si, and the carbon is dissolved in the austenite. From the calculation results by equation (6) [32,33], the carbon concentrations of retained austenite were much higher than those of ferrite matrixes, which is 1.38 wt% with the cooling rate of 0.05°C/s and 1.02 wt% even with the cooling rate of 2°C/s as shown in Figure 10(b).

(6) [ C ] = { a 0 0.35780.000096 [ Mn ] + 0.00152 [ Si ] 0.00031 [ Mo ] } / 0.0033

where a ₀ is the actual lattice parameter of retained austenite, obtained by the XRD patterns, for example, the values of a ₀ are equal to 0.3603 nm with the cooling rate of 0.05°C/s; [Mn], [Si] and [Mo] stand for the weight percentages of elements Mn, Si and Mo in austenite, respectively. The MS of the retained austenite can be calculated by equation (7) [34]:

(7) M S ( ° C ) = M S 0 − 564 ( X C − X ¯ )

where M S 0 is the transformation start point of martensite (414.8°C calculated by J Mat Pro), with the average carbon content. With the cooling rate of 0.05°C/s, MS of the retained austenite is −222.5°C, which means the retained austenite obtained with the low cooling rate is stable enough to some extent. X _C is the carbon concentration of the retained austenite (wt%). X ¯ is the average carbon concentration (wt%) in the material.

Figure 11

(a) The X-ray diffraction profile and (b) amount of retained austenite under different cooling rates.

Figure 11(b) shows that the amount of the retained austenite decreases with the increase of the cooling rate. In particular, the amount was decreased from 11.41 to 4.67% when the cooling rate increased from 0.05 to 2°C/s. To recognize the distribution and the morphology of the retained austenite further, EBSD was adopted, and the results are shown in Figure 12. Figure 12(m) and (n) show the image quality maps of GB and LB microstructures under the cooling rates of 0.05 and 0.3°C/s, respectively. Figure 12(o) and (p) show the phase distribution maps for ferrite and austenite in Figure 12(m) and (n). In Figure 12(o) and (p), the red block represents the austenite phase (FCC) and the green part represents the ferrite phase (BCC). These results further prove that, for GB at the cooling rate of 0.05°C/s, the retained austenite was dispersed in the matrix of bainite ferrite in the block and most of the retained austenite block or M/A islands formed along the grain boundaries. Meanwhile, for LB at the cooling rate of 0.3°C/s, the retained austenite was distributed between two adjacent parallel laths, along with the obvious reduction of the retained austenite (FCC phase). Consequently, the decrease of the retained austenite as the increasing cooling rate, as shown in Figure 11, can be explained from two aspects. On the one hand, the volume fraction of the retained austenite decreased because of the substantial reduction of GB with lots of M/A islands during the microstructure transformation from GB to LB with the faster cooling rate. On the other hand, the phase transformation driving force of austenite into martensite is raised due to the increased subcooling degree with the increase of cooling rate, and then, more martensite formed and less retained austenite was reserved relatively. In addition, from Figure 11(b), the amount of the retained austenite declined sharply during the cooling rate of 0.05–1°C/s and dropped slowly with the cooling rate of 1–2°C/s, which is thought to be related to the transition from GB to LB during low cooling rates and the simple martensite phase transformation during faster cooling rates.

Figure 12

Characteristics of the experimental steel at the cooling rates of 0.05, 0.3, 0.5 and 1°C/s by EBSD: (a)–(d) EBSD maps in the inverse pole figure coloring (see the inset); (e)–(h) grain boundary misorientation maps; (i)–(l) misorientation angle distribution histograms; (m)–(n) image quality maps at the cooling rates of 0.05°C/s and 0.3°C/s, respectively and (o)–(p) phase maps for ferrite and austenite of (m) and (n).

It is generally known that at very low cooling rates, the phases formed are closed to those corresponding to equilibrium, that is, ferrite and pearlite. However, from the results shown in Figures 5(a) and 8(b), the microstructure contained a very small amount of polygonal ferrite even at the cooling rate of 0.05°C/s. The retardation of the proeutectoid ferrite transformation is considered to be caused by the addition of microalloying elements such as Mn, Mo and Nb. Furthermore, as shown in Figure 8(b) and Table 2, both the B_s and B_f points were decreased with the increase in the cooling rate, which is owing to the inhibition of carbon diffusion at the faster cooling rate, and a larger subcooling is required for phase transformation.

Besides the transformation microstructure and behavior, the Vickers hardness and nanohardness were influenced by the cooling rate, as shown in Figures 7 and 9(c). Conversely, in the range of 0.05–2°C/s, the Vickers hardness increased rapidly and grew very slowly or almost unchanged within a cooling rate of more than 2°C/s. As known to all, both the increase of dislocation density and the change of morphology from granular bainite to lath bainite or martensite are beneficial to enhance strength. Hence, focusing on the cooling rate range of bainite transformation, the microstructure characteristics under the cooling rates of 0.05, 0.3, 0.5 and 1°C/s were investigated by EBSD as shown in Figure 12. From Figure 12(a) and (e), the granular bainite under the cooling rate of 0.05°C/s was composed of large blocky bainite ferrites separated by high angle grain boundaries with M/A islands, which have the similar orientation with the ferrite matrix and enveloped by high angle grain boundaries. Figure 12(b) and (f) shows that the lath bainite under the cooling rate of 0.3°C/s was composed of packets of nearly parallel laths (or plates) with different orientations, and the bainitic ferrites were separated by high angle grain boundaries. Besides, Figure 12(p) shows that there are a small amount of elongated retained austenite films between two bainitic laths. Apparently, the sizes of both bainite ferrite and the retained austenite of LB with the cooling rate of 0.3°C/s are less than those of GB with the cooling rate of 0.05°C/s. When the cooling rate increased to 0.5°C/s, both the length and width of the bainite ferrite became small, and the microstructure was further refined as shown in Figure 12(c) and (g). Finally, at the cooling rate of 1°C/s, the martensite appeared in the microstructure.

As shown in Figure 12(d) and (h), a parent austenite grain group of martensite could be divided into several packets with different crystallographic orientations and separated by high misorientation angle grain boundaries, and each packet could be further subdivided into blocks, which have a certain amount of martensitic laths with similar orientations and separated by low misorientation angle grain boundaries [35]. The refinement of microstructure and phase transformation strengthening are beneficial to the improvement of hardness. In addition, the dislocation density under these four different cooling rates is presented in Figure 13. Figure 13(a) shows X-ray diffraction patterns. Then, the amplified (200) diffraction peaks (ferrite) indicated that with increasing the cooling rate from 0.05 to 0.5°C/s, the diffraction peaks were broadened and enhanced. But when the cooling rate increased to 1°C/s, the diffraction peak narrowed and weakened, compared with that under 0.05°C/s. As shown in Figure 13(c), the dislocation density was increased in the range of the cooling rate from 0.05 to 0.5°C/s. However, from 0.5 to 1°C/s, the dislocation density did not increase, instead, slightly reduced. To be specific, the results are 4.01(±0.14) × 1,016, 5.47(±0.12) × 1,016, 6.02(±0.16) × 1,016 and 5.98(±0.27) × 1,016 in sequence. It is generally known that the increase of the dislocation density is conducive to strength enhancement. But the reason for the higher hardness of the microstructure under 1°C/s could be related to the solid solution strengthening by the reserved carbon in matrix under the faster cooling rate and enough high dislocation density.

Figure 13

Dislocation density characterization of experimental steel under different cooling rates: (a) X-ray diffraction profile; (b) comparison of (200) diffraction peak (ferrite) and (c) variation of dislocation density by different cooling rates.

Moreover, Figure 12(i)–(l) present the fractions of high misorientation angle grain boundaries, which are 75.9, 78.0, 78.8 and 67.7% under the cooling rate of 0.05, 0.3, 0.5 and 1°C/s, respectively. Since high misorientation angle grain boundaries can improve the toughness by preventing crack propagation effectively, the toughness of the tested samples was increased when the cooling rate is slower than 1°C/s according to the results of EBSD, and the best is the sample under the cooling rate of 0.5°C/s. In addition to misorientation angle grain boundaries, the values of Young's modulus of microstructure under the cooling rates of 0.05, 0.5 and 1°C/s were measured by nanoindentation, as shown in Figure 9(c). As Hahn and Rosenfield suggested, the elastic modulus is a key parameter in toughness, and they are positively correlated, as expressed in equation (8) [36,37].

(8) K Ic ≃ [ 2 3 E σ y ε f ( 0.0005 + n 2 ) ] 1 / 2

where E is Young's modulus, σ _y is the tensile yield stress, n is the strain hardening coefficient and ε _f is the critical true strain for coalescing voids.

With the combination of Figure 9(c) and equation (8), the toughness under the cooling rate of 0.5°C/s is the best, which is considered that the microstructure under the cooling rate of 0.5°C/s has a good matching of strength and toughness.

From the results of semi-industry, the Charpy impact energy was 38 J as presented in Table 3, and the fracture surface was observed by SEM, as shown in Figure 14. Presented in Figure 14(a), the impact fracture surface could be divided into three zones, shear lips (Figure 14(c)), ductile region (Figure 14(e)) and brittle region (Figure 14(g)). The shear lips are composed of large and flat dimples. The ductile region is composed of deep dimples of different sizes and some voids, which represents the typical ductile fracture characteristic. The brittle region is mainly composed of a large cleavage river pattern coupled with some tear ridges, which is positive for the enhancement of impact energy [38].

Figure 14

SEM macrographs of (a) tensile and (b) impact fracture surface for experimental steel; (c) SEM micrograph of shear lip in (a); (d) SEM micrograph of shear lip in (b); (e) SEM micrograph of ductile region in (a); (f) SEM micrograph of ductile region in (b) and (g) SEM micrograph of brittle region in (a).

As presented in Table 3, the yield strength was 847 MPa and the tensile strength was 1,298 MPa. Figure 14(b) shows that the tensile fracture surface could be divided into two zones, shear lips and ductile region. The shear lips are presented by lots of dimples in Figure 14(d). The ductile region is presented by deep dimples of different sizes and voids in Figure 14(f). Whether it is a tensile fracture surface or impact fracture surface, both exhibit the characteristics of the ductile fracture.

5 Conclusions

For nonquenched and tempered carbide-free bainitic steels designed in the present study, the effects of the cooling rate after hot deformation on phase transformation, microstructure and mechanical properties have been investigated by the Gleeble-3800 thermal simulator and the semi-industry test. According to the results of TEM, XRD, EBSD and so on, the main conclusions are as follows:

1. It is feasible to obtain high strength and toughness carbide-free bainite steel by nonaustempering through reasonable alloy composition design with high Si and Mo and the controlled cooling rate after hot deformation.

2. With the increase of the cooling rate, the microstructure changed from granular bainite, lath bainite to martensite, and the hardness was enhanced by the phase transformation strengthening of LB/M. The amount of the retained austenite decreased with the increase in the cooling rate. The morphology of the retained austenite changed from blocky to film, which is advantageous for improving toughness.

3. The microstructure with the cooling rate of 0.5°C/s had the best combination of high strength (R _m: 1,298 MPa and R _el: 847 MPa) and good toughness (A _kv: 38 J) in the studied range of the cooling rate, which is a profitable trial to improve the strength based on a considerable toughness in a nonquenched and tempered way.

Acknowledgments

The authors gratefully acknowledged the financial support provided by National Natural Science Foundation of China (Grant No. 51674020) and National Key Research and Development Program of China (Grant No. 2016YFB0300102).

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