The Effect of Annealing and Aging Temperature on the Microstructure and Properties of an 800 MPa Grade Dual-Phase Steel with High Formability

Abstract

To develop automotive steel with a higher strength–ductility balance, an 800 MPa grade Nb-Ti microalloyed dual-phase steel with high elongation was designed. Continuous annealing tests were conducted using a continuous annealing simulation testing machine. The effects of annealing temperature and aging temperature on the microstructure and mechanical properties were investigated through SEM, EBSD, TEM, and tensile testing. The results indicate that the microstructure of the steel primarily consists of ferrite, martensite, and a small amount of retained austenite. As the annealing temperature rises, the martensite content increases, and the retained austenite content first increases and then decreases. Therefore, the tensile strength and elongation initially increase and then decrease, while the yield strength gradually decreases. As the aging temperature rises, the martensite content decreases, while the tempering degree of martensite increases and the retained austenite content rises. Therefore, the tensile strength gradually decreases, and the yield strength and elongation gradually increase. The optimal comprehensive performance was achieved with an annealing temperature of 830 °C and an aging temperature of 330 °C, resulting in a tensile strength of 915 MPa, a yield strength of 470 MPa, and an elongation of 24.4%. This represents a 28.4% increase in elongation compared to conventional dual-phase steels of the same strength grade. The strength–ductility product reaches 22.3 GPa%.

1. Introduction

With the growing demand for lightweighting and safety in the automotive industry, the research and development of advanced high-strength steels (AHSS) have become a critical direction in materials science. As the first generation of high-strength steels, dual-phase steels have been widely applied in the automotive field [1]. The microstructure of traditional dual-phase steel (DP steel) is mainly composed of ferrite and martensite, which has the characteristics of high strength, good ductility, and high initial work hardening rate. Its mechanical properties are closely related to the relative proportions, distribution, and microstructural characteristics of the two phases [2,3,4]. However, with the increase in strength, the plasticity of DP steel decreases rapidly, and it is difficult to form on many parts with high tensile ductility [5]. Some scholars [6,7] suggest that appropriately reducing the yield strength or using an ultra-low yield ratio (less than 0.5) is beneficial for stamping, while others [8] propose that introducing a certain amount of retained austenite into DP steels is an effective method. The dual-phase steel with high formability (DH steel) is achieved by introducing a small amount of metastable retained austenite based on a two-phase structure. Under stress deformation conditions, the retained austenite undergoes phase transformation-induced plasticity (the TRIP effect), which improves the strength and ductility of the material. For example, Zhou Li et al. [9] introduced 5.1% retained austenite into the test steel by increasing the carbon content in DP780, resulting in a 50% increase in elongation. The improvement in the formability of DH steel has not increased the difficulty of production, and it possesses good weldability and resistance to edge shear sensitivity, making it more suitable for the processing and forming of complex structural components and safety parts. As a new type of advanced high-strength steel, it has great market application prospects.

A large number of studies have shown that the continuous annealing process parameters, including annealing temperature and aging temperature, significantly influence the microstructure and properties of DH steel [10]. Xiaoyue Ma et al. [11] studied the effect of aging temperature on 1200 MPa grade DH steel. It was found that with the increase in aging temperature, the content of retained austenite increased, the tensile strength and yield strength decreased, and the elongation increased. The optimum properties of 1209 MPa tensile strength and 14.6% elongation were obtained when the test steel was annealed at 860 °C and aged at 310 °C. Liu Pengfei et al. [12] studied the effect of annealing temperature on the microstructure and properties of 980 MPa grade cold-rolled dual-phase steel. It was found that with the increase in annealing temperature, the proportion of ferrite in the microstructure decreased and the yield strength increased. When the annealing temperature is 840 °C and the aging temperature is 250 °C, the strength–ductility product (the product of tensile strength and elongation) is 15.6 GPa%.

C, Si and Mn are the main alloying elements in the design of cold-rolled dual-phase steel in China, and their content directly determines the final microstructure and mechanical properties. Al contributes to deoxidizing molten steel and inhibits the decomposition of retained austenite as well as the precipitation of carbides during the aging process. Its function is similar to Si. Cr element can increase the hardenability of the steel to ensure its strength and stabilize the retained austenite. The addition of microalloying elements such as Nb and Ti to DH steel can inhibit the recrystallization of austenite and prevent the growth of ferrite grains. During the aging stage, the carbonitrides of Nb and Ti precipitate, and these precipitated particles can hinder dislocation movement, thereby achieving the effect of precipitation strengthening [13,14,15]. Wu et al. [16] studied the effect of the Nb element on hot-dip galvanized DH780 steel. It was found that the addition of trace Nb not only refined the grains of the steel and improved the uniformity of the galvanized layer but also increased the ferrite content and decreased the martensite/bainite content. The tensile strength decreased from 860 MPa to 815 MPa, and the elongation increased from 15% to 17%. Although Nb addition improved elongation, the retained austenite content was only 6%, which hindered the full activation of the TRIP effect of retained austenite.

At present, the strength of dual-phase steel has been greatly improved, but its elongation still cannot meet the corresponding requirements. Many scholars have studied the introduction of retained austenite into dual-phase steels to improve elongation. However, the dual-phase steels they prepared contained relatively little retained austenite (about 5%), resulting in a limited improvement in elongation and insufficient utilization of the TRIP effect. Thus, significant potential for further elongation improvement remained. In order to obtain DH steel with higher elongation at the same level, the composition was reasonably designed in this paper, and more than 8% retained austenite was introduced into the test steel through appropriate heat treatment process. Eventually, DH steel with excellent comprehensive performance was obtained, which provided an important reference for industrial production.

2. Materials and Methods

In this study, a C-Si-Mn-Al-Cr-based chemical composition was designed, and Nb and Ti elements were added appropriately to facilitate grain refinement and ensure strength. C is a martensite-strengthening element. Excessive content will increase the brittleness of the experimental steel, while insufficient content is unfavorable for the stability of retained austenite. Generally, the C content of dual-phase steel with a tensile strength of 800 MPa grade should not exceed 0.2%; therefore, the C content is set at approximately 0.17%. Mn can expand the γ-phase region, stabilize austenite, and improve hardenability, but excessive Mn content tends to cause segregation. Typically, the Mn content of dual-phase steel with a tensile strength of 800 MPa grade should not exceed 2.5%; thus, the Mn content is set at around 2.0%. Both Al and Si can inhibit the precipitation of carbides. Controlling the contents of Al and Si within the range of 1.0–1.5% enables the synergistic effect of Si and Al, thereby improving the toughness and ductility of the steel. Cr can increase the hardenability of steel and stabilize retained austenite, but excessive content raises costs. Therefore, its content is controlled below 0.2%. Both Nb and Ti contribute to grain refinement strengthening and precipitation strengthening. To save costs, the contents of both elements are controlled at approximately 0.02%. The raw materials were melted into steel ingots using a vacuum melting furnace at Benxi Iron & Steel Group Co., Ltd. (Benxi, China). The raw material information used is as follows: electrolytic manganese flakes (Erachem Comilog, New Johnsonville, TN, USA; Mn ≥ 99.9%), low-carbon ferrochromium (Outokumpu, Helsinki, Finland; Cr ≥ 65%), aluminum granules (KBM Affilips, Oss, The Netherlands; Al ≥ 99.7%), ferrosilicon (Elkem, Oslo, Norway; Si: 75%), ferroniobium (CBMM, Araxá, Brazil; Nb: 65–67%), and ferrotitanium (VSMPO-AVISMA, Verkhnyaya Salda, Russia; Ti: 68–72%), and high-purity graphite recarburizer (Fangda Carbon New Material Co., Ltd., Lanzhou, China; C ≥ 99.9%). The actual chemical composition of the test steel was determined using an inductively coupled plasma optical emission spectrometer (5900 ICP-OES, Agilent Technologies Inc., Santa Clara, CA, USA), as shown in

Table 1. Chemical composition of test steel (wt.%).

C	Mn	Si + Al	Cr	Nb	Ti	P	S	N
0.17	2.04	1.20	0.17	0.022	0.016	≤0.015	≤0.008	≤0.004

. The measurement accuracy of C, Mn, Si, Al, and Cr is ±0.01 wt.%, and the measurement accuracy of other elements is ±0.001 wt.%.

The ingots were heated to 1200 °C for 30 min and then forged into forging billets. These billets were then cut into dimensions of 30 mm × 80 mm × 100 mm for rolling. The cut billets were held at 1200 °C for 1 h, then hot-rolled from 30 mm down to 3.7 mm, achieving a total reduction of 88%. The start rolling temperature was 1150 °C and the finish rolling temperature was 900 °C. After finishing rolling, the steels were water-cooled to 550 °C, followed by a simulated coiling process at 550 °C for 1 h, and then cooled with the furnace. The hot-rolled plates were pickled to remove surface oxide scales, and then the pickled plates were cold-rolled along the hot-rolling direction. The cold-rolling reduction rate was 60%. Rectangular specimens of 50 mm × 190 mm were cut from cold-rolled plate, and continuous annealing experiments were carried out on a CCT-AY-II (Ulvac-Riko Inc., Tokyo, Japan) continuous annealing machine.

The continuous cooling transformation (CCT) curve of the test steel was calculated by using JMatPro (v7.0, Sente Software Ltd., Surrey, UK, 2021, Proprietary license) software, and the critical point temperatures were as follows: A_c1 = 707.1 °C, A_c3 = 896.7 °C, M_s = 383.4 °C, M_f = 271.1 °C (the CCT curve is shown in Figure S1 of the Supplementary Materials). In order to obtain sufficient initial austenite in the annealing holding stage to form a strengthened martensite in the subsequent cooling, and to retain part of the ferrite to ensure plasticity, annealing at a higher temperature within the range of A_c1 and A_c3 was selected. The aging process aims to induce moderate tempering of the martensite formed after rapid cooling, resulting in a certain volume fraction of tempered martensite, thereby ensuring a good strength–ductility balance in the experimental steel. Therefore, the aging temperature was set between M_f and M_s. Accordingly, the annealing temperatures were set at 810 °C, 830 °C, 850 °C, and 870 °C, with the aging temperature fixed at 350 °C. To investigate the effect of aging temperature on microstructure and properties, the aging temperatures were set at 310 °C, 330 °C, 350 °C, and 370 °C, with the annealing temperature fixed at 830 °C. The specific process route is shown in Figure 1.

Subsequently, microstructural observation was conducted on the annealed specimens. The microstructure of the steel annealed at 810–870 °C and aged at 310–370 °C was observed by Sigma 360 field emission scanning electron microscope (FE-SEM, Sigma 360, Oberkochen, Germany). Five regions were randomly selected for observation in each group of temperatures and the most representative was selected. The specimens were prepared through mechanical polishing followed by etching in a 4 vol.% nital solution for 2 s. The morphology and distribution of retained austenite were characterized using electron backscatter diffraction (EBSD, Symmetry S3, Abingdon, UK) and transmission electron microscopy (TEM, JEM2100, Tokyo, Japan). The scanning voltage of EBSD test was 20 kV and the step size was 0.08 μm. The content of retained austenite was quantified using Oxford Instruments AZtecCrystal (v2.1, Oxford Instruments NanoAnalysis, Buckinghamshire, UK, 2021) software, with an accuracy of ±0.01%. Three separate areas of each EBSD specimen were analyzed, and the most representative result was selected. The EBSD specimens were electropolished in a 10 vol.% perchloric acid solution. The voltage of electrolytic polishing was 13 V, the current was 2 A, and the electrolytic polishing time was 18 s. The TEM specimens were prepared by ion thinning and carbon extraction replica. The ion thinning specimens were used to observe the retained austenite, and the carbon extraction replica specimens were used to observe the precipitated phase.

According to GB/T228.1-2021 standard [17], the tensile specimens were cut from the continuous annealing plate with a gauge distance of 50 mm. The room temperature tensile test of the tensile specimens was carried out by using the (MTS Systems Corporation, Eden Prairie, MN, USA) tensile testing machine. Three tensile specimens were taken from each temperature group for tensile testing and the final results were taken as the average of the three tests. The tensile speed was 2 mm/min, and the tensile direction was parallel to the rolling direction.

3. Results and Discussion

3.1. Effect of Annealing Temperature on Microstructure

Figure 2 shows the microstructure of the test steel annealed at 810–870 °C and aged at 350 °C. The microstructure is primarily composed of ferrite matrix and island martensite. When the annealing temperature is 810 °C, ferrite and martensite exhibit a banded distribution, and the martensite is mainly quenched martensite. As the annealing temperature increases, the volume fraction of martensite gradually rises, and the banding characteristics gradually diminish. When the annealing temperature rises to 850 °C, the banded microstructure is almost invisible, as shown in Figure 2c. When the annealing temperature is 870 °C, a small amount of bainite and tempered martensite appear in the microstructure, and the tempered martensite shows a fuzzy martensite lath. This is due to the increase in austenitization at higher annealing temperature, which induces a reduction in the average alloy content solvated in austenite, diminished austenite stability, thereby inducing the formation of a small amount of bainite during the subsequent cooling process [18]. Moreover, the depletion of alloy content within austenite undermines the steel’s temper resistance, leading to the formation of tempered martensite [19].

Due to the challenge of accurately identifying retained austenite in SEM test, EBSD was used to characterize the content and distribution of retained austenite in the microstructure, and the EBSD maps are shown in Figure 3. In the band contrast (BC) map, quenched martensite with a higher density of dislocations appears darkest in contrast, while ferrite exhibits a lighter contrast. The red region represents the retained austenite (RA), which is mainly distributed in the ferrite/martensite grain boundaries, and ferrite grains. Statistical analysis reveals that the RA content initially increases and then decreases (the RA content was marked in the upper right corner of the BC maps). This is because as the annealing temperature increases, the initial austenite content rises, and thus the content of retained austenite also increases accordingly. However, when the annealing temperature continues to rise, the initial austenite content further increases, resulting in a decrease in the average alloy content within the austenite. This reduces the stability of the austenite, causing a significant transformation during the subsequent cooling process. Thus, the final retained austenite content decreases. When the annealing temperature is 830 °C, the retained austenite content reaches the highest level of 8.98%. The retained austenite will produce a TRIP effect during the deformation process, and the increase in its content is beneficial to the improvement of the plasticity and toughness of the material [20,21,22].

In order to study the morphology and distribution of precipitates at different annealing temperatures, the precipitates of the specimens were observed by transmission electron microscopy and analyzed by energy spectrum. The results are shown in Figure 4. Figure 4(a1,b1) show the morphology of the precipitated phases at 830 °C and 850 °C annealing temperatures, and the large-sized square precipitated phase and the small-sized circular precipitated phase can be observed at both annealing temperatures. At 830 °C, the small-sized precipitates are in the majority, and at 850 °C, the large-sized precipitates are in the majority. This may be because higher annealing temperature enhance atomic diffusion capacity, making it easier for second-phase particles to aggregate and grow. This process follows the Ostwald ripening mechanism [23], where small-sized particles dissolve and large-sized particles continue to grow. Statistical analysis revealed that the average precipitate sizes were 8.59 nm at 830 °C and 10.66 nm at 850 °C, as shown in Figure 4(a3,b3). The EDS analysis was performed on the marked places in Figure 4(a1,b1), and the results are shown in Figure 4(a2,b2). Both precipitated phases are Nb-Ti composite precipitated phases, tentatively identified as (Nb,Ti)C particles. The small-sized particles have a higher proportion of Nb, while the large-sized particles contain a higher proportion of Ti, indicating that Ti-rich precipitates tend to form larger particles. Coarsening of precipitates reduces their obstruction to dislocations, potentially leading to a decrease in strength [24]. The copper peaks observed in Figure 4(a2,b2) originate from the copper grid used as the specimen carrier during the TEM test.

3.2. Effect of Annealing Temperature on Mechanical Properties

Table 2. Mechanical properties of the test steel at different annealing temperatures.

Annealing Temperature/°C	Yield Strength/MPa	Tensile Strength/MPa	Total Elongation/%	Yield Ratio	Strength–Ductility Product/GPa·%
810	521 ± 14	816 ± 15	24.0 ± 0.5	0.64	19.6
830	499 ± 15	833 ± 12	24.9 ± 0.7	0.60	20.7
850	507 ± 11	810 ± 8	24.5 ± 0.2	0.63	19.8
870	490 ± 8	807 ± 7	22.2 ± 0.9	0.61	17.9

lists the mechanical properties of the test steel from room-temperature tensile tests, and Figure 5a gives the engineering stress–engineering strain curves of the test steels at different annealing temperatures. As the annealing temperature increases, the tensile strength and elongation first increase and then decrease, and the yield strength decreases from 521 MPa to 490 MPa. When the annealing temperature increases from 810 °C to 830 °C, the tensile strength increases from 816 MPa to 833 MPa due to the increase in the volume fraction of hard-phase martensite in the microstructure. The variation in elongation with annealing temperature is related to the retained austenite content and the tempering degree of martensite [25]. At an annealing temperature of 830 °C, on one hand, the increase in retained austenite content enhances plasticity through the TRIP effect during tensile deformation; on the other hand, the microstructure becomes more homogeneous, and the tempering degree of martensite increases, which enhances the ductility of the martensite phase, resulting in an elongation as high as 24.9%. At this temperature, a good matching of strength and plasticity was achieved between the phases, and the strength–ductility product was as high as 20.7 GPa%. When the annealing temperature rises from 830 °C to 870 °C, the tempering degree of martensite increases, causing carbon dissolved in martensite to precipitate progressively. Since the hardness of martensite primarily depends on its carbon content [26], the tensile strength of the specimen gradually decreases. Additionally, at an annealing temperature of 870 °C, the partial replacement of martensite by bainite also contributes to the reduction in tensile strength. Meanwhile, the retained austenite content decreases, the martensite volume fraction increases, and martensite coarsening occurs, resulting in embrittlement and a subsequent reduction in elongation. The decrease in yield strength may be due to the reduced dislocation density in ferrite with increasing annealing temperature, making it easier to yield during deformation [27]. It may also be because the coarsening of the precipitated phase weakens the hindrance to dislocation, thus leading to a decrease in yield strength.

The strain hardening rate n is closely related to the formability. The uniform plastic deformation stage of the tensile process is generally fitted by the Hollomon equation, as shown in Equation (1).

$$\mathsf{\sigma }=\mathrm{k}{\mathsf{\epsilon }}^{\mathrm{n}}$$

(1)

where σ is the true stress, ε is the true strain, k is the strength coefficient.

The instantaneous value of n can be obtained by the logarithmic derivative of Equation (1), as shown in Equation (2).

$$\mathrm{n}={\displaystyle \frac{\mathrm{d}(\mathrm{I}\mathrm{n}\mathsf{\sigma })}{\mathrm{d}(\mathrm{In}\mathsf{\epsilon })}}$$

(2)

Figure 5b shows the strain hardening rate–true strain curves of the test steel at different annealing temperatures. It can be seen that the strain hardening rate shows the same trend. The curves can be divided into three stages: first, it drops sharply from a high initial value; then the strain hardening rate decreases slowly with increasing strain; and finally, a rapid decrease due to necking. This result is consistent with the findings of Wu Qiuyun et al. [28]. During uniform plastic deformation, strain hardening and softening compete, leading to the generation and annihilation of dislocations. The essence of work hardening is actually dislocation movement [29]. In the first stage of deformation, the content of martensite is lower at 810 °C, and the hindrance of martensite to dislocation movement in ferrite is weaker, so the strain hardening rate is lower. In the middle and late stages of deformation, martensite with lower carbon content at 870 °C is prone to plastic deformation under lower stress, which weakens the constraint of martensite on dislocation motion in ferrite; that is, the coordinated deformation between ferrite and martensite is improved [30,31], so the strain hardening rate is lower.

3.3. Effect of Aging Temperature on Microstructure

Figure 6 shows the microstructure of the test steel annealed at 830 °C and aged at 310–370 °C. The test steel is mainly composed of ferrite and martensite. When the aging temperature is 310 °C, a large amount of martensite has been formed before the isothermal process. This part of martensite is tempered during the isothermal process and transformed into tempered martensite, so a large amount of tempered martensite can be seen in the microstructure. As the aging temperature increases, the amount of austenite transformed into martensite during the rapid cooling stage decreases, leading to a reduction in tempered martensite content but an increase in tempering degree [32]. When the aging temperature rises to 370 °C, a large amount of bainite appears in the microstructure, presumably due to the aging temperature being in proximity to the bainite transformation zone.

Figure 7 shows the EBSD maps of the test steel at different aging temperatures. The retained austenite is distributed in the ferrite/ferrite grain boundary, ferrite/martensite phase interface, or ferrite grain in the form of a block or film. With the increase in aging temperature, the content of retained austenite in the test steel gradually increases (the RA content was also marked in the upper right corner of the BC maps). This is because the increase in aging temperature will promote the diffusion of carbon atoms from supersaturated martensite to austenite, allowing more carbon to enter the austenite and improving the stability of retained austenite [33]. Therefore, the retained austenite content increases with the increase in aging temperature. When the aging temperature rises to 370 °C, the retained austenite content reaches the highest level of 9.18%.

In order to further explore the morphology and distribution of retained austenite, the specimens aged at 330 °C and 350 °C were characterized by TEM, as shown in Figure 8. Figure 8a,d display the microstructural morphology of ferrite (F) and martensite-austenite (M/A) islands. Ferrite appears white, while M/A islands appear black, and numerous dislocation lines can be observed within ferrite grains. Through calibration of diffraction patterns, the black blocky structures at ferrite grain boundaries in Figure 8b,e are identified as M/A constituents, and M/A is a structure formed by the incomplete transformation of carbon-rich austenite and martensite during the cooling process [34]. Figure 8c,f show the strip-like retained austenite in ferrite grains with thicknesses of about 0.3–0.4 μm and 0.3–0.6 μm, respectively. A row of second-phase particles precipitated along the dislocation line can also be observed in the ferrite grains (as shown in Figure 8g). The energy spectrum analysis of the second phase particles shows that the precipitated phase contains significant amounts of Nb and Ti microalloying elements, which may be a Nb-Ti composite precipitated phase.

3.4. Effect of Aging Temperature on Mechanical Properties

Table 3. Mechanical properties of the test steel at different aging temperatures.

Aging Temperatures/°C	Yield Strength/MPa	Tensile Strength/MPa	Total Elongation/%	Yield Ratio	Strength–Ductility Product/GPa·%
310	399 ± 11	919 ± 7	21.5 ± 0.8	0.43	19.7
330	470 ± 13	915 ± 18	24.4 ± 0.4	0.51	22.3
350	499 ± 15	833 ± 12	24.9 ± 0.7	0.60	20.7
370	591 ± 10	793 ± 15	25.9 ± 0.5	0.75	20.5

lists the room-temperature tensile mechanical properties of the test steel at different aging temperatures, and Figure 9a shows the engineering stress–engineering strain curves. It can be seen that when the aging temperature is 310–350 °C, the tensile curve tends to yield continuously, while the tensile curve of the specimen aged at 370 °C has an obvious yield platform. The presence of a yield plateau in the tensile curve is closely related to the tempering degree of martensite [35]. Higher aging temperature leads to greater tempering of martensite, resulting in more carbides (cementite) precipitating during tempering [36]. These carbides pin mobile dislocations, reducing their availability during subsequent deformation, thereby increasing yield strength and forming a yield plateau [37]. When the aging temperature is 310 °C, the tensile strength is as high as 919 MPa but the elongation is only 21.5% due to the formation of more bulk martensite. With the increase in aging temperature, the tempering of martensite is more sufficient, the carbon content in martensite decreases continuously, and softening occurs gradually [38]. And the quenched martensite in the structure gradually decreases, which is replaced by retained austenite and softer bainite, so the tensile strength of the test steel gradually decreases and the elongation gradually increases. When the aging temperature is 330 °C, the microstructure is composed of ferrite, martensite, and a small amount of retained austenite, achieving an excellent strength–ductility balance between hard and soft phases. At this time, the comprehensive performance of the test steel is the best. The tensile strength, yield strength, and elongation are 915 MPa, 470 MPa, and 24.4%, respectively, and the strength–ductility product is as high as 22.3 GPa%. When the aging temperature increases to 370 °C, the elongation is as high as 25.9%, but the tensile strength is only 793 MPa, and the yield ratio is relatively high at 0.75.

Studies have shown [39] that the yield ratio of 800 MPa grade dual-phase steel should be controlled below 0.75. In practical applications, a too high yield ratio will deteriorate the workability of dual-phase steel, weaken the work hardening capacity, and easily cause problems such as cracking and wrinkling during stamping of parts, reducing the pass rate and increasing costs. Figure 9b shows that as the aging temperature increases, the yield ratio gradually rises, while the strain hardening rate of the specimen gradually decreases. This indicates a significant negative correlation between the yield ratio and strain hardening capacity. The lower yield ratio indicates that the steel still has a significant stress increase space after yielding. These dislocations continue to accumulate and proliferate at the interface of the soft and hard phases, which makes the flow stress increase continuously, and the strain hardening rate is high macroscopically [40]. Conversely, when the yield ratio is high, the steel approaches its ultimate strength before yielding, leaving almost no room for further strengthening after yielding. Dislocations are difficult to move and further proliferate in the initial stage, so the strain hardening capability is weak. When the aging temperature is 370 °C, the specimen exhibits the highest yield ratio and the lowest strain hardening capacity. Therefore, the aging temperature of 370 °C is not an optimal temperature.

To verify whether the retained austenite exerted the TRIP effect during tensile deformation, EBSD characterization was performed on the fracture surfaces of tensile specimen annealed at 830 °C and aged at 330 °C (Figure 10). It can be found that the retained austenite content of the specimen after stretching is significantly reduced, from 7.59% to 0.21%, indicating that approximately 7.38% of the retained austenite transformed into martensite via the TRIP mechanism. Only a trace amount of retained austenite remains after fracture, and this untransformed retained austenite is distributed on the ferrite/martensite grain boundaries, indicating that the retained austenite distributed in the ferrite grains transformed more completely.

4. Conclusions

This paper designs the chemical composition of an 800 MPa grade DH steel and investigates the effects of annealing/aging temperature on the microstructure and properties under a continuous annealing process. The following conclusions were drawn:

(1)

Increasing the annealing temperature leads to a rise in martensite content, while the retained austenite content initially increases and then decreases. Consequently, both the tensile strength and elongation first increase and then decrease. When the aging temperature was held constant at 350 °C, the specimen annealed at 830 °C attained the highest retained austenite content (8.98%), effectively activating the TRIP effect and leading to optimal comprehensive mechanical properties.

(2)

Increasing the aging temperature reduces the content of tempered martensite but increases the content of retained austenite. Consequently, the elongation gradually improves, while the tensile strength gradually decreases. When the annealing temperature was maintained at 830 °C and the aging temperature was set at 330 °C, the retained austenite content reached 7.59%, achieving an excellent balance among ferrite, martensite, and retained austenite. At this point, the specimen exhibited optimal comprehensive properties: a tensile strength of 915 MPa, an elongation of 24.4%, and a product of strength and elongation as high as 22.3 GPa%.

This study reveals the internal relationship between the continuous annealing process parameters and the content of retained austenite in DH steel, which provides a theoretical reference for improving the performance by regulating the content of retained austenite through heat treatment, and provides an important reference for the industrial production of DH steel with good strength and plasticity of 800 MPa grade.