Optimizing Microstructure and Property by Ausforming in a Medium-carbon Bainitic Steel

Abstract

The transformation behavior and microstructure in a medium-carbon bainitic steel were investigated by combination of metallography and dilatometry. The fine micro-structural units of carbide-free bainite in non-ausformed and ausformed materials were measured by a transmission electron microscope. Mechanical stabilization of austenite in deformed material and its effect on property were analyzed by nanoindentation and tensile tests. Ausforming with a strain of 0.2 at 573 K can not only accelerate bainite transformation, but also improve the comprehensive properties. The strength and ductility of nanostructured bainitic steel can be simultaneously enhanced by ausforming, which should be attributed to the refinement of bainite and the enhanced volume fraction of retained austenite. Compared to the non-deformed material, the mechanical stabilization of austenite can be optimized by ausforming, resulting in good transformation-induced plasticity effects. Also a very important advantage was that, the bainite transformation time could be minimized into practical scale by prior ausforming compared to traditional low-temperature austempering.

1. Introduction

Automobile industries are continuously challenged to reduce weight and improve fuel efficiency due to economic and environmental requirements. For that purpose, advanced high strength steels were paid great attention to, such as quenching and partitioning (Q&P) steel,¹⁾ medium Mn transformation-induced plasticity (TRIP) steel,²⁾ TRIP aided bainitic ferrite (TBF) steel,³⁾ etc. Carbide-free bainite usually formed in Si-rich steels accompanying with blocky and film-like retained austenite (RA), which can contribute to excellent combination of strength, ductility and toughness.⁴⁾

To obtain nano-scaled bainite, composition design with high-carbon high-alloy and very low temperature austempering were initially utilized.⁵⁾ However, it took several hours even a couple of days to complete transformation, which is unpractical from the viewpoint of production. Also the coarse blocky austenite with low stability usually transforms partially to brittle martensite during cooling after isothermal bainite transformation, resulting in deterioration of ductility and toughness.⁶⁾ Many efforts have been done to solve above problems, including optimal composition and processing design.^7,8,9,10) Decreasing carbon content can shorten the incubation of bainite nucleation, but raise martensite starting temperature (Ms) simultaneously. The theory indicates that a high level of substitutional solute was required at a low-carbon steel to produce nanostructured bainite where the difference between the bainite and the Ms then vanishes.¹¹⁾ Also at low carbon concentrations, experimentally observations show that the thin platelets of bainite tend to coalesce into coarse grains which are detrimental to strength and toughness.⁸⁾ The prospects therefore look promising for the design of a medium-carbon nanostructured bainite.

As for the control of RA, a multi-step bainitic austempering process was developed to reduce blocky microstructure and refine the subunits.¹²⁾ During the first step isothermal holding, a higher austempering temperature was normally adopted to partially form the relatively coarse bainitic ferrite plates. Extra carbon content can be rejected from newly formed bainite to surrounding austenite, leading to its chemical stabilization and a decrease in Ms. Thus, one can use a deep-cold isothermal bainite transformation at a lower temperature to obtain finer bainite. The secondly formed plates can also divide the untransformed austenite into pieces which further facilitates the stability of RA. In addition, an inverted multi-step bainitic austempering route was proposed,¹³⁾ in which the second-step temperature was higher than that in the first step. Compared to the previous strategy, a higher temperature in the second step can accelerate the bainite transformation. Although both above two multi-step methods contribute to the refinement of bainite and RA, leading to good comprehensive properties, the procedure is complicated and still takes more than two hours to complete transformation. Fortunately, a low temperature ausforming can be utilized to not only accelerating transformation rate but also promoting the amount of bainite.^14,15) Gong et al.¹⁶⁾ considered that planar dislocations remaining on the active slip planes at 573 K can assist bainite transformation, accompanied by strong variant selection. Anisotropic dilatation during bainite transformation also illustrates the general rule of variant selection, that is, in each austenite grain, the main variant of which domain the bainite should belong to the same Bain group.¹⁷⁾ Therefore, one can employ ausforming process to obtain sufficient nanostructured bainite in an expected transformation time. However, the control of RA by ausforming, which plays an important role in determining the property and performance, is still not clear. The relationship between ausformed bainite and property should be further clarified.

In the present work, the ~0.43 wt.% C bainitic steel was designed to investigate the effect of ausforming on bainite transformation and property. The morphology of bainite and RA in the deformed material is also discussed in details, and efforts are also made to build a link between property and ausformed bainite.

2. Experimental Method

The chemical composition of the steel is 0.43C-1.90Si-2.83Mn-0.57Al-0.06Cu (wt.%) with Fe balance. Si was added to suppress cementite precipitation during bainite austempering. Mn was added to increase the hardenability to obtain bainite and Al is for accelerating transformation. The material was refined in vacuum induction furnace and casted into a 50 kg ingot. The ingot is homogenized at 1523 K for 5 hours, followed by 8 passes hot rolling into a 12.0 mm slab. After hot rolling with finishing temperature of 1173 K, the steel plate was air-cooled to ambient temperature. The bainite starting temperature (Bs) and Ms of the tested steel were calculated as 707 K and 516 K, respectively, according to the following equations:¹⁸⁾   

$$B_S=839-{\sum }_i{P}_i{x}_i-270\times \left[1-exp\left(-1.33x_C\right)\right]$$(1)

  

$$M_S=565-{\sum }_i{K}_i{x}_i-600\times \left[1-exp\left(-0.96x_C\right)\right]$$(2)

where i = Mn, Si, Cr, Ni, and Mo, and the concentration x in wt.%.   

$${\sum }_i{P}_i{x}_i=86x_{Mn}+23x_{Si}+67x_{Cr}+67x_{Ni}+75x_{Mo}$$(3)

  

$${\sum }_i{K}_i{x}_i=31x_{Mn}+13x_{Si}+10x_{Cr}+18x_{Ni}+12x_{Mo}$$(4)

Samples for thermomechanical simulation tests were machined to a dumbbell shape with the central cylinder of 8.0 mm diameter and 12.0 mm height. The surface of samples was conventionally polished to keep the measurement face level and minimize the effect of surface roughness. Ausforming and austempering tests were conducted on a Gleeble-3500 thermal simulator according to the processing schedules shown in Fig. 1. The specimens were heated to 1273 K at 5 K s⁻¹ and kept for 15 min, and then cooled to 573 K. A cooling rate of 5 K s⁻¹ was utilized to avoid ferrite and pearlite transformation. Subsequently, two routes were designed, i.e. one was applied compressive deformation with a strain of 0.2 at the rate of 1 s⁻¹ and the other was free of strain. The load was promptly removed after deformation to keep the sample free of external compressive stress during austempering. The deformed sample was then isothermally kept at 573 K for 60 min, while the non-deformed one was held at 573 K for 90 min shown as the blue line in Fig. 1. After isothermal holding, both two cases were quenched to ambient temperature. The dilatation data during the whole process were recorded by a laser dilatometer. To investigate the tensile property of ausformed bainite, hot rolling and consequent salt bath treatment tests were carried out using 140 mm × 20 mm × 10 mm blocks. The schedule was same as that in the thermomechanical simulation test. After ausforming, the hot-rolled sample was put into a salt-bath furnace and isothermally held at 573 K for certain time.

Fig. 1.

Heat treatment and ausforming routes. (Online version in color.)

Microstructures were characterized on Keyence optical microscope (OM) and a Nova 400 Nano field emission scanning electron microscope (FE-SEM) with an acceleration voltage of 20 kV. The metallographical specimens were etched with 4% nital. The fine microstructures were observed using a JEM-2100F transmission electron microscope (TEM). The samples for TEM examination were mechanically ground down to 30 μm thickness and then electrolytic thinned to perforation using an electrolyte composed of 5% perchloric acid and 95% glacial acetic acid at ambient temperature, and voltage of 40 V. The volume fractions of RA in the ausformed and non-ausformed samples after isothermal transformation were determined using an X’Pert diffractometer with CoKα radiation under the following conditions: acceleration voltage, 40 kV; current, 150 mA; and step, 0.06°. Tensile specimens were prepared according to ASTM standard and the strain rate was ~4×10⁻³ s⁻¹. Nanoindentation was performed on the lightly etched specimens using 2000 μN load for 30 s in a triboindenter TI-900 equipped with scanning probe microscope. Each test involved a 5 × 5 array of indents. Vickers hardness tests were performed on a HV1000A micro-hardness tester (0.2 kg-1960 mN). The average value of at least ten individual measurements was calculated, including several martensite bands and bainite blocks of the microstructure.

3. Results and Discussion

3.1. Original Microstructure and Property

A small metallographic specimen was selected from the center of hot-rolled slab and the original SEM micrograph of the steel was shown in Fig. 2(a). The microstructure consists of granular bainitic ferrite and carbides, which shows microstructure characteristic of medium-temperature transition. Figure 2(b) shows the phase transition under equilibrium conditions calculated by Thermal-calc with the database of TCFE-6. The starting temperature of γ→α transformation is 1081 K, and cementite begins to precipitate at 1004 K. After hot rolling, the steel was air cooled to ambient temperature, leading to intermediate phase transformation. However, the present work aims to produce nanostructured bainite. Thus a larger cooling rate should be considered. The tensile strength and total elongation of the test steel are 1016 MPa and 23.3%, respectively, which can be taken as reference.

Fig. 2.

(a) SEM micrograph of initial materials before heat treatment; (b) phase diagram of equilibrium calculation, γ refers to face-centered cubic (fcc) austenite and α is body-centered cubic (bcc) ferrite. (Online version in color.)

3.2. Microstructure after Ausforming and Austempering

Figure 3 shows the microstructures of the non-ausformed and ausformed samples after isothermal transformation for different times. Both the samples comprise of needle-like bainite and matrix of martensite, as shown by arrows in Figs. 3(a) and 3(b). The deformed sample seems to contain more bainite compared to the non-deformed one in spite of the shorter isothermal holding time. From the high-resolution SEM micrographs in Figs. 3(c) and 3(d), a sheave of bainitic ferrite (BF) consists of several subunits with film-like RA between them. Blocky RA also can be seen in both cases. Large blocks of martensite (M, dotted box in Fig. 3(c)) still exist in the non-deformed materials even undergoing 90 min isothermal transformation, while the microstructure of the deformed one was divided into pieces. Carbon diffusion from BF into untransformed austenite during the progression of bainite transformation has been proved by many researchers. Base on T₀ theory, bainite transformation ceases when carbon concentration of untransformed austenite reaches the critical value where the free energy of austenite is equal to that of ferrite.¹⁹⁾ In addition, both partition local equilibrium (PLE) and negligible partition local equilibrium (NPLE) modes proved that solute drag effects can bring isothermal bainite transformation into stable stage,^20,21) leading to remaining of austenite after long isothermal holding. Therefore, blocky M was obtained during the cooling process after bainite transformation. The growth of bainite can section residual austenite, thus one can see that the marked M block was cut into four parts by three sheaves. Ausforming can stimulate dense nucleation of bainite, resulting in multiple-crossed morphology in Fig. 3(d).

Fig. 3.

(a) and (c) OM and SEM micrographs of the non-deformed sample after isothermal holding at 573 K for 90 min; (b) and (d) OM and SEM micrographs of the ausformed sample after isothermal holding at 573 K for 60 min. (Online version in color.)

3.3. TEM Results

The fine microstructure of BF was confirmed by examining the microstructure at higher magnifications using TEM. Figure 4 shows a typical morphology of bainitic plates in the non-ausformed steel, in which most of them are in the form of clusters of parallel plates, emanating from the prior austenite grain boundaries shown by yellow arrows. Thin film-like RA is trapped between adjacent BF plates with longer and straight feature, as shown in Fig. 5. The dark features lying within the plates and at the angle to the axis of these plates in Fig. 5(a) are similar in appearance to typical low bainitic carbides. However, they are actually regions of RA, as evidenced by dark field micrograph of Fig. 5(b). The bainitic plate is composed of sub-units and the films of austenite have sometimes a typical wavy morphology characteristic in high Si steels, which are dispersed between the sub-units of BF. The sub-units do not appear to be entirely separated by RA. The extent of transformation in some plates can be so large as to make the individual plate indistinct, for example, Fig. 4. In addition, different crystallographic orientation of plates shown by red line can be seen in one prior austenite grain.

Fig. 4.

Typical TEM micrographs showing full morphology of bainite in the non-ausformed sample transformed at 573 K for 90 min. (Online version in color.)

Fig. 5.

Plate-shaped bainite formed in the non-ausformed sample, (a) bright field image and (b) corresponding dark field image taken from an austenite reflection, α refers to bainite plate and γ is RA. (Online version in color.)

Figure 6 presents a characteristic microstructure that was achieved when prior-deformation of 20% was applied followed by isothermal transformation at 573 K for 60 min. The general morphology is different from that in the non-deformed sample, that is, the less crystallographic orientations of sheaf appear. In local area, BF sheaves align more or less in a common direction. It may due to the variant selection of bainite in ausformed austenite, which has been theoretically and experimentally proved by many works.^16,17,22,23) BF exhibits the orientation relationships with respect to the parent austenite, scattered around Kurdjumov-Sachs (K-S), Nishiyama-Wassermann (N-W) and Greninger-Troiano (G-T) relationships.²⁴⁾ There are different crystallographic bainitic variants that can form from a given orientation of the parent austenite. Ausforming can increase the probability of certain variants to form since the energy that is needed to transform from austenite to that crystallographic variant is lower than for others. When only several of the variants have grown from each prior austenite grain, it is said that there is variant selection, resulting in the micro-structural alignment.

Fig. 6.

TEM Micrograph of BF sheaves align more or less in a common direction in ausformed sample. (Online version in color.)

The width and length of bainite unit were quantitatively measured and averaged on different fields of TEM observations, where hundreds of units were measured for each thermomechanical and isothermal history. The variation of the average width and length of bainitic unit, aspect ratio and respective standard deviation for the different thermo-mechanical conditions are summarized in Fig. 7. For the non-ausformed material, the average width and length is ~0.724 μm and ~4.398 μm, respectively, while ~0.151 μm in width and ~0.830 μm in length are presented for the ausformed sample. For both width and length of unit, it can be seen that the value decrease significantly in the sample with a strain of 0.2 when compared with the non-deformed one. The results demonstrate that microstructures formed during prior deformation limit the growth of bainite unit. However, it is of interest to note that the aspect ratio (width/length) remains unobvious changed for different conditions.

Fig. 7.

Average width and length of bainite plate unit in different thermo-mechanical conditions. (Online version in color.)

It is clearly seen that the grain boundaries serve as the nucleation site of growth of bainitic plate in non-deformed sample, while in prior-deformed sample to produce the fine intersected platelet structure characteristic of BF sheaves, two conditions are necessary to be satisfied. One is a number of intragranular nucleation events are required, another is that nucleation gives rise to multiply oriented platelet variants so that the chaotic interlocking structure is evident in Figs. 3(d) and 6. Additional nucleation sites within grain can be induced by dislocations formed by the plastic deformation of the austenite.¹⁰⁾ The nucleation site is increased in deformed sample but the growth space is restrictive, resulting in that each nucleus then transforms a smaller size due to interaction of neighbouring platelets.

3.4. Dilatometry Analysis

Figure 8 shows the dilatation results of the non-ausformed and ausformed samples along the diameter direction during isothermal bainite transformation at 573 K. There is a distinct difference between the final amounts of dilation in two cases, shown in Fig. 8(a). For the non-deformed sample, the dilatation can represent the real amount of bainite transformation, while it is inappropriate for the deformed one due to variant selection. According to the micro-structural result, the amount of bainite in deformed austenite is lightly larger than that in the non-deformed sample. However, the dilatation result shows a noticeable increase. Anisotropic dilatation during bainite transformation in ausformed austenite in nanostructured bainitic steel can be caused by variant selection. He et al.¹⁷⁾ also reported that the ausformed sample shrunk in the axial direction but expanded in the radial direction. Nevertheless, the transformation rates can be calculated based on dilatation data using tangent method, and the result was given in Fig. 8(b). One can assign the approximately horizontal dilatation at the final stage of each case in Fig. 8(a) as one hundred percent of individual bainite reaction. The incubation period (the time to complete 5% bainite transformation) for the bainite transformation in the ausformed sample is only about 287 seconds, which is obviously shorter than the non-ausformed sample (about 1120 seconds). Also the time to complete bainite transformation is much shorter in the ausformed case. The accelerated transformation kinetics can be attributed to the enhanced nucleation efficiency and facilitated variants by deformation.

Fig. 8.

(a) Dilatation change versus time during isothermal holding reflecting bainite transformation, and (b) transformation rates in non-ausformed and ausformed samples, showing a much faster kinetics of bainite transformation after prior deformation.

3.5. Volume Fractions of Different Phases

The XRD diffractograms for the non-deformed and deformed materials are given in Fig. 9. The volume fraction of RA was calculated based on integrated intensities of (200)α, (211)α, (200)γ, (220)γ, and (311)γ diffraction peaks, according to Eq. (5):²⁵⁾   

$$V_i=\frac{1}{1+G\left(I_{\alpha }/I_{\gamma }\right)}$$(5)

where V_i is the amount of austenite for each peak, I_α and I_γ are the integrated intensity of ferrite and austenite peaks, respectively, and G is constant for each peak. The carbon content in RA was estimated using Eq. (6):²⁶⁾   

$$\begin{array}{l}{C}_{\gamma }=({\alpha }_{\gamma }-3.5780-0.00095M{n}_{\gamma }\\ \hspace{1em}\hspace{1em}-0.0056A{l}_{\gamma }-0.0015C{u}_{\gamma })/0.033\end{array}$$(6)

where C_γ is the carbon content in RA and α_γ is the lattice parameter of RA, which is determined by the position of three austenite peaks using Cohen’s method, and Mn_γ, Al_γ and Cu_γ represent the concentration of individual elements (mass-%) in RA. Based on microstructure results, the volume fractions of bainite were calculated using Image Pro-Plus software, and the detail was described in Ref.²⁷⁾. After multiple measurements, the average results are shown in Table 1. The amount of bainite can be promoted by ausforming with a strain of 0.2 at 573 K, although the isothermal holding time is shorter than the non-ausformed sample. Moreover, the volume fraction of RA was increased to ~23.2% by deformation, indicating that the stability of austenite can be enhanced by ausforming. The higher carbon content of RA in the non-deformed sample is mainly due to the long bainite transformation time, which allows more carbon diffusion into adjacent austenite.

Fig. 9.

X-ray diffraction diagrams, (a) non-ausformed sample and (b) ausformed with a strain of 0.2. (Online version in color.)

Table 1. Volume fractions of different phases (%) and carbon content in RA (wt.%).

Sample	BF	M	RA	Cγ
Non-ausformed	53.8±2.4	38.4±1.8	7.8±0.8	1.18±0.08
Ausformed	63.3±2.5	13.5±1.3	23.2±1.5	0.96±0.08

3.6. Mechanical Properties

Vickers micro-hardness tests were carried out on different phases, illustrated in Fig. 10(a), in which BF and M were located in different areas. It is obvious that the area of the collapsed rhombus of M is small than that of BF. The average results were plotted in Fig. 10(b). The hardness of bainite in the ausformed sample is 563±34 HV_0.2, larger than that in the non-ausformed one (532±29 HV_0.2). On the contrary, the hardness of M decreases slightly after deformation with a strain of 0.2, which should be attributed to the lower carbon content in the residual austenite ~0.96 wt.%.

Fig. 10.

(a) Image showing hardness tests on BF and M, and (b) Vickers micro-hardness results of non-ausformed and ausformed samples. (Online version in color.)

The nominal engineering stress-strain curves were given in Fig. 11. The ultimate tensile strengths (UTS) of the non-deformed and deformed samples are ~1521 MPa and ~1755 MPa, respectively. It indicates that ausforming with a strain of 0.2 at 573 K can increase the strength of bainitic steel. Furthermore, the total elongation (TE) of the ausformed material ~18.1% is better than that of the non-ausformed one ~14.7%. Compared with the original material, both non-ausforming and ausforming processes can significantly increase the strength by microstructure strengthening due to isothermal bainite transformation.

Fig. 11.

Engineering stress-strain curves of non-ausformed and ausformed samples. (Online version in color.)

According to Bhadeshia and Young,^28,29) the strength of bainite and martensite can be calculated as follows in Eq. (7):   

$$\sigma ={\sigma }_{Fe}+{\sum }_i{\sigma }_{SS}^i+{\sigma }_C+k_{ϵ}{\left({\overline{L}}_3\right)}^{-1}+{\sigma }_{ppt}+K_D{\rho }_d^{0.5}$$(7)

Where σ_Fe is strengthening contribution of pure annealed bcc iron, ${\sigma }_{SS}^i$ is solid solution strengthening due to substitutional solutes, σ_C is solid solution strengthening due to carbon, σ_ppt is precipitation strengthening from carbide particles and $k_{ϵ}{\left({\overline{L}}_3\right)}^{-1}$ is the lath size strengthening component including the constant k_ϵ, and ${\overline{L}}_3$ is mean linear intercept of laths. $K_D{\rho }_d^{0.5}$ is the dislocation strengthening component of the model, where the constant K_D = 0.38 μb for bcc metals, μ is the shear modulus and b is the burgers vector (~0.25 nm). Compared to the non-ausformed sample, the bainite laths in the ausformed sample were much finer, leading to a high strength. The dislocation density is calculated based on the TEM measurements, 2.51 × 10¹⁵ m⁻² and 5.32 × 10¹⁵ m⁻² for non-deformed and deformed samples, respectively. Based on Eq. (7), the calculated contribution to strength due to refinement of BF by ausforming is about 603 MPa, and that due to enhanced dislocation is ~174 MPa. One may note that the difference between the UTSs is just ~234 MPa, smaller than the combined calculation. The reason should be attributed to the larger amount of hard martensite in non-deformed sample. From Table 1, the carbon content of non-deformed RA is the higher as ~1.18 wt.%, resulting in enhanced hardness of martensite as well as strength.

3.7. Mechanical Stabilization

In low-temperature banitic steel, the strength is mainly related to the fine scale of the BF plates and martensite (hard phases), ductility is mostly controlled by the RA (soft phase).³⁰⁾ In order to take full advantage of TRIP effect, the stability of RA, i.e., its capability to transform to martensite under stress, must be neither too low nor excessively high. Before that, the stabilization of austenite during isothermal transformation should be considered first, which directly determine the volume fraction of RA at ambient temperature. The stabilization of ausformed austenite depends on carbon content and dislocation caused by deformation, in which the latter plays a more important role,³¹⁾ leading to the larger amount of RA in the deformed sample.

Figure 12 shows the fine structures of BF and RA in the ausformed samples. From Fig. 12(a), the phase of γ randomly distributes among α-plates, which mostly seems to be in a coarse size than that in non-ausformed sample (Fig. 5(b)). Actually, film-like RA also exists between ausformed BF, which can be clarified by dark field image taken from an austenite reflection in Fig. 12(c). With a strain of 0.2 at 573 K, high-density dislocation was introduced into the parent austenite, which then inherited by bainitic plate. The dislocation-rich BFs were observed in dark field image taken from a ferrite reflection, as shown by arrows in Fig. 12(b). It was reported by He et al.³²⁾ that intensive mobile dislocations can not only increase strength but also improve ductility, which was realized by deformed and partitioned (D&P) process. In the present work, BF and RA with high dislocation density were simultaneously produced by ausforming and austempering processes, resulting in a high product of strength and elongation ~31.8 GPa% (UTS×TE).

Fig. 12.

TEM micrographs showing bainite plates formed in ausformed sample, (a) bright field image, (b) corresponding dark field image taken from a ferrite reflection, (c) dark field image taken from an austenite reflection, and (d) corresponding diffraction pattern. (Online version in color.)

Finally, nanoidentation tests were performed to investigate the mechanical stabilization of RA in non-ausformed and ausformed samples. The typical load-displacement curves were plotted in Fig. 13, in which the terraces pointed by arrows signify the transformation from RA to M. It is interesting to find that, the loads to induce M transformation in non-deformed sample are all above ~1500 μN (Fig. 13(a)), while it mostly ranges 250~1000 μN for the ausformed one (Fig. 13(b)). For the very rare indent, the load increases to ~1780 μN shown by red line in Fig. 13(b). It indicates that the RA in non-ausformed sample is too stable due to high carbon content and mostly film-like morphology. The tensile fracture happened before the TRIP effect was produced, leading to a limited ductility. Another important reason is that the volume fraction of RA is just ~7.8% which is very small. Bhadeshia³³⁾ et al. reported that tensile failure in nanostructured bainite occurred when the RA content is diminished to about 10%. In this case, only the ausformed sample can have TRIP effect, which seems corresponding to the nanoidentation result. In total, ausforming with a strain of 0.2 at 573 K can not only increase the strength but ductility as well, which can be attributed to refinement of BF and large volume fraction of RA. Also a very important advantage was that, the bainite transformation time could be optimized into a practical scale.

Fig. 13.

Nanoindentation results of (a) non-ausformed and (b) ausformed samples, and five typical load-displacement curves were selected corresponding to five indents in upper-left image, terrace means martensite transformation induced by stress. (Online version in color.)

4. Conclusions

The transformation behavior and microstructure evolution in designate medium-carbon nanostructured bainitic steel were investigated by metallographic and dilatometry. Mechanical stabilization of austenite in deformed material and its effect on property were analyzed by nanoindentation and tensile tests. The conclusions can be drawn as follows:

(1) Ausforming with a strain of 0.2 at 573 K can not only accelerate bainite transformation, but also improve the comprehensive property which shows enhanced combination of strength and ductility, i.e. ultimate strength ~1755 MPa and total elongation ~18.1%.

(2) Low temperature bainite can be refined by ausforming due to additional nucleation sites induced by dislocation. However, the aspect ratio (width/length) remains unobvious changed for different conditions. Micro-structural alignment of bainite sheaves can form in deformed austenite due to variant selection.

(3) The simultaneous increase of strength and ductility can be attributed to refinement of BF and more volume fraction of RA. Compared to the non-deformed material, the mechanical stabilization of austenite can be optimized by ausforming, resulting in better TRIP effects.

Acknowledgement

The authors gratefully acknowledge the financial supports from the National Nature Science Foundation of China (Nos. 51704217, 51874216), China Postdoctoral Science Foundation (2017M622533), the Major Projects of Technological Innovation in Hubei (No. 2017AAA116), Hebei Joint Research Fund for Iron and Steel (E2018318013). The authors also thank Prof. Hatem Zurob at McMaster University, Canada for the suggestions and corrections on the present work. TEM work was conducted at Canadian Centre for Electron Microscopy (CCEM), Hamilton, ON, Canada.

Data Availability

Both the raw data and the processed data required to reproduce these findings can be provided by the authors.

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