The Effect of Ultrafast Heating on the Microstructure and Mechanical Properties of the 2.2 GPa Grade Hot Forming Steel

Abstract

The aim of the present work is to evaluate the effect of ultrafast heating on the microstructure and mechanical properties of hot forming steel. The initial microstructure utilized in this study was a cold-rolled microstructure, and the test steel was heated to full austenitization at a rate of 200 °C/s, followed by water quenching. It was observed that the ultrafast heating process significantly refines both the prior austenite grains and martensite laths while inheriting high-density dislocations from the initial cold-rolled microstructure. Consequently, the coupling mechanism between dislocation strengthening and grain refinement strengthening remarkably enhanced both the yield strength and ultimate tensile strength of the test steel. Eventually, the yield strength of the hot forming steel reached 1524 MPa, along with an ultimate tensile strength of 2221 MPa and uniform elongation of 5.2%.

1. Introduction

Hot forming steel has become one of the most prevalent advanced high-strength steels (AHSS) due to its exceptional formability advantages, negligible springback, low alloy element addition, and ultra-high strength. It is primarily utilized for the production of load-bearing and energy-absorbing structural components. The conventional hot forming process typically involves cold rolling and long-term annealing of the strip based on a hot rolled plate, followed by austenitizing in a roller hearth furnace and subsequent transfer to die for forming and quenching. The integrated forming–cooling process enables the manufacturing of safety-related car body components with high strength in a single step [1,2]. Currently, 22MnB5 steel with an ultimate tensile strength (UTS) of 1500 MPa and yield strength (YS) of 1000 MPa represents one of the most mature commercially hot forming steels. However, as more stringent standards regarding gas emissions and fuel consumption are imposed for automotive lightweight purposes, there is an urgent demand for higher-strength hot forming steels [3,4,5]. Optimized methods to meet such requirements include increasing carbon content [6], incorporating microalloying elements (Ti/Nb/V/Mo) [7,8,9], optimizing production processes [10,11,12], and replacing AHSS with much lighter non-ferrous alloys, such as 7xxx aluminum alloys [13].

In recent years, the ultrafast heating process, also termed the flash process, has garnered significant attention in academia and the automotive industry due to the intricate nature of phase transformation behavior during ultrafast heating and the challenges associated with comprehending and controlling this complexity. Several studies [14,15,16,17,18] have reported that the drastic reduction in heating time caused by ultrafast heating impedes the dissolution of cementite in the pearlite region within the initial microstructure, thereby hindering carbon diffusion, element migration, and austenite grain interface movement. Consequently, this results in an uneven element distribution, the refinement of austenite grains, and an enhancement in the strength and toughness of steel. Dai [19] employed ultrafast heating (300 K/s) to produce high-performance Q&P steel. The non-uniform distribution of elements coupled with rapid austenite transformation kinetics led to ultrafine and carbon-rich austenite formation after full austenitization. Compared to austenitization at a slow heating rate of 4 K/s, there was a significant increase in the product of strength and plasticity from 19.5 GPa% to 25.7 GPa%. Li [20] discovered that an ultrafast heating rate (500 °C/s) dramatically influences the interaction between the recrystallization of deformed ferrite and the austenitization process, resulting in a transformation from net-like martensitic morphology to a chain-like structure, which partly enhances the strength of dual-phase steel. Wen [21] investigated the impact of the ultrafast heating process on the microstructure and mechanical properties of cold-rolled 7 wt.% Mn steel, which revealed the formation of bimodal austenite grains, along with an increase in the volume fraction of retained austenite. Such factors combined with high-density dislocations inherited from cold-rolled microstructure resulted in excellent mechanical properties in which the product of strength and plasticity reached as high as 55 GPa%.

The impact of ultrafast heating on the microstructure and mechanical properties of hot forming steel has also been documented. Kolleck [22] introduced an induction heating system for the hot forming process of 22MnB5 steel, achieving a maximum heating rate of 200 K/s. Compared to the long-term austenitizing process in a traditional roller hearth furnace, the YS of 22MnB5 steel only increased by 30 MPa, while the UTS significantly increased by 178 MPa. Wen [23] successfully achieved rapid austenitization of 22MnB5 steel without holding through resistance heating, which resulted in the refinement of austenite grains and the increase of UTS to 1703 MPa. Hou [24] researched a novel hot forming steel that incorporated a higher content of Cr and Si elements based on 22MnB5 steel. The cold-rolled test steel austenized to 930 °C at a heating rate of 100 °C/s and held for 120 s. Compared to 22MnB5 steel, the UTS of test steel rises from 1583 MPa to 1722 MPa. Grydin [25] found that contact heating can realize rapid austenitization of 22MnB5 steel, leading to the refinement of austenite grains and the martensite substructure. Such an improvement contributed to an increase in YS and UTS of approximately 100 MPa. However, previous studies have mainly focused on the effect of the ultrafast heating process on the UTS of hot forming steel, with limited attention given to YS. Additionally, most research on the ultrafast heating process has been conducted on 1500 MPa grade hot forming steels. At the same time, studies need to investigate the effect of the ultrafast heating process on the microstructure evolution and mechanical properties of higher-strength hot forming steels, such as those with UTS reaching 1800 MPa or even 2000 MPa.

The primary objective of this study is to evaluate the feasibility of producing ultra-high-strength hot-formed steel through the ultrafast heating process and investigate its impact on the microstructure evolution and mechanical properties of test steels compared with the traditional hot forming process. The electron backscatter diffraction technique (EBSD) was used to reconstruct the PAGs and count their grain size distribution, while transmission electron microscopy (TEM) was employed to further characterize the martensite and measure the average lath width. Additionally, this study quantitatively discussed the contribution of dislocation strengthening and grain refinement strengthening mechanisms to the YS of ultra-high-strength hot forming steel.

2. Materials and Methods

Table 1. The chemical compositions of the investigated steels (wt.%).

C	Si	Mn	B	Cr	V	N	Fe
0.39	0.34	1.68	0.004	0.31	0.044	0.0040	Bal.

shows the chemical compositions of the test steel in this study. The forging ingot was heated to 1200 °C and held for 1 h for solution treatment, followed by hot rolling. The thickness of the ingot was gradually reduced from 35 mm to 5 mm with a reduction ratio of approximately 20% for each pass, and the total reduction rate was 85.7%. The finish rolling temperature (FRT) was 950 °C. The steel plate was held at 610 °C in the rolling furnace for 2 h before being air cooled and cold rolled to a thickness of 1.2 mm with a total reduction rate of 76%. The cold rolling was carried out using a four-roll cold rolling stand with a tension simulator, where the diameter of the working roll was 175 mm, and that of the support roll was 380 mm.

Rectangular specimens measuring 200 mm × 30 mm × 1.2 mm were cut from cold rolled sheets and underwent hot forming treatment, with the length direction and rolling direction (RD) of the sheet maintained parallel. The phase transition temperatures of austenite and martensite were measured using a DIL-805 A type dilatometer using 10 × 4 × 1.2 mm³ rectangular samples. The ultrafast heating tests were conducted on the Gleeble 3500 thermal simulator at heating rates of 10 and 200 °C/s, with an austenitizing temperature set as 900 °C, a holding time of approximately 20 s, and a cooling rate for quenching maintained around 100 °C/s to ensure the formation of a complete martensitic structure in the test steels. The ultrafast-heated specimens were named UFH-10 and UFH-200. For comparison, the traditional heating treatment was carried out on reference samples where the cold rolled sheet underwent recrystallization annealing at 680 °C for 4 h, followed by the same hot forming treatment as that applied to ultrafast heating with a rate of 10 °C/s. This reference sample was named THF. Figure 1 illustrates the schematic diagram depicting the hot forming process for all three test steels.

The microstructure of the investigated steels was examined using scanning electron microscopy (SEM, ZEISS, ULTRA 55, Carl Zeiss AG, Oberkochen, Germany), EBSD (FEI Quanta FEG 450, FEI Company, Hillsboro, OR, USA), and TEM (JEOL JEM 2010F, JEOL, Tokyo, Japan). The characterization was performed on the TD plane (normal to the transverse direction) at the center of the heat-treated zone, where the controlling thermocouple was welded. Samples for SEM were standardly polished and etched in 4% nital for 10 s. EBSD measurements were performed at 20 kV with a step size of 0.12 μm and an angular accuracy of 0.1° on the sample that had been electropolished in an electrolyte (perchloric acid/acetic acid = 1:9). TEM samples were prepared using twin-jet polishing (5% perchloric acid alcohol at −30 °C, voltage: 50 V) after thinning the thickness of samples from 400 μm to 70 μm through mechanical grinding. X-ray diffraction (XRD) measurements using Cu-Kα radiation of 0.15406 nm wavelength were performed at 40 kV and 150 mA with a step size of 0.02° (2θ) and a counting time of 1.2 s. The diffraction peaks of the BCC structure were used to estimate the dislocation densities in test steels according to the literature [26,27].

Mechanical properties were identified using the uniaxial tensile with the sub-size tensile specimens (gauge width: 4 mm and gauge length: 10 mm), and the tensile direction was parallel to the rolling direction. The tensile test was performed with a strain rate of 3.33 × 10⁻⁴ s⁻¹ (crosshead speed was 1 mm/min) at room temperature. The strain was measured using an extensometer. The YS was determined as the 0.2% offset, whereas the uniform elongation (UE) was determined at the UTS. Due to the non-standard tensile specimens, the total elongation (TE) was only used as reference data within this study. All the tensile tests were repeated three times.

3. Results

3.1. Microstructure Evolution before Austenitization and Quenching

The hot-rolled (HR) microstructure and V-containing precipitates of the test steels are characterized in Figure 2. After hot rolling and coiling, the HR microstructure consists of polygonal ferrite and pearlite islands containing lamellar ferrite and lamellar cementite carbide. These pearlite islands are randomly distributed without indicating banded distribution, suggesting an even distribution of C/Mn in the matrix. The volume fraction of pearlite islands is estimated to be approximately 35%. Proeutectoid ferrite grains include equiaxed ferrite and long-strip ferrite. By measuring the equivalent circular grain size of equiaxed ferrite, it was determined that the average size of ferrite grains is 4 μm, while long strip ferrites have an average width of about 3.5 μm with varying lengths. The non-recrystallized zone temperature (T_nr) for the test steel was calculated using Equation (1) [28] to be 1010 °C, which exceeds the finish rolling temperature of 950 °C. Consequently, most proeutectoid ferrite grains retained their morphology from non-recrystallized austenite as elongated strips, with only a tiny portion undergoing recrystallization during coiling to form equiaxed grains.

T_nr = 887 + 446 C + (6445 Nb − 644 Nb1/2) + (732 V − 230 V1/2) + 890 Ti + 363 Al − 357 Si

(1)

where C, Nb, V, Ti, Al, and Si represent the mass fractions of corresponding solid solution elements in wt.%.

Most cementite carbides are distributed in the ferrite grain boundaries as spherical and short strips, with only a few located within the ferrite grains formed during the proeutectoid ferrite transformation process. Due to the slow cooling rate, newly formed ferrite surrounds carbon-rich austenite, eventually transforming into single or chain-like cementite carbides. These particles in pearlite exhibit a chain-like morphology composed of small spherical particles rather than a perfect layer. A TEM analysis of carbon extraction replicas (Figure 2b) reveals the presence of V-containing precipitates in the HR microstructure. Yellow dotted circles have been used to highlight the precipitates to illustrate their distribution and morphology. The average diameter of these precipitates is estimated as 10 nm by counting over 300 particles. The EDS composition analysis in Figure 2c confirms that these precipitates are V (C, N).

Figure 3 depicts the cold-rolled and batch-annealed microstructures of the test steels. The cold-rolled microstructure consists of deformed ferrite and deformed pearlite, exhibiting a flat and elongated morphology. The average widths of ferrite and pearlite are approximately 3.1 μm and 3.7 μm, respectively, while the length data exhibit a relatively scattered distribution. The layered cementite in pearlite is not completely fragmented, and some carbides still maintain their morphology. Additionally, scattered spherical cementite carbides can be observed within the ferrite matrix. After batch annealing, a fully recrystallized microstructure comprising equiaxed ferrite grains and spherical carbides is obtained. The ferrite grain size distribution is not uniform, and the average value is about 4.9 μm. There were too many carbide particles to calculate their average size and density.

3.2. Microstructures after Austenitization and Quenching

Figure 4 shows the SEM micrographs of the microstructures in UFH-200, UFH-10, and THF steels after austenitization and quenching. A fully martensitic microstructure was obtained at a cooling rate of 100 °C/s during the quenching process following austenitizing isothermal treatment, which is attributed to the high hardenability of medium carbon steel containing boron. Compared to UFH-10 (Figure 4c) and THF (Figure 4f) steels, the SEM micrographs with a higher magnification reveal that the martensite packet and block boundaries are less distinct in UFH-200 (Figure 4b) steel. Meanwhile, the PAG boundaries are also difficult to distinguish accurately. The average PAG size of the THF steel is larger, since only several PAGs are present in a limited region. Although there is a non-uniform distribution in PAG size for UFH-10 steel, more PAGs are observed at the same magnification. Additionally, the martensite substructure size of THF steel is relatively large, indicating the significant influence of PAG size on it. Furthermore, a few undissolved carbides were detected within the as-quenched microstructures of both UFH-10 and THF steels.

The IPF diagrams in Figure 5a₁–c₁ illustrate the martensite structures of the test steels. Notably, the UHF-200 steel exhibits a significantly refined martensite structure. By analyzing the IPF color of the EBSD micrograph, the average sizes of martensite blocks in UFH-200, UHF-10, and THF steels were determined to be 1.98 ± 0.12 μm, 2.61 ± 0.19 μm, and 2.92 ± 0.23 μm, respectively. Furthermore, the PAGs were reconstructed based on the Kurdjumov–Sacks (K-S) crystal orientation relationship between the parent phase (austenite) and child phase (martensite) utilizing AZtec Crystal 2.1 software [29], as shown in Figure 5a₂–c₂. Compared to the fully recrystallized initial microstructure (THF steel), a significant refinement in the average size of PAGs was obtained after austenitization and quenching when the initial microstructure consisted of deformed pearlite and ferrite (UFH-200 and UFH-10 steels).

It should also be noted that the heating rate plays a crucial role in determining the average PAG size. As the heating rate increased from 10 °C/s to 200 °C/s, there was a significant refinement in the average size of PAGs quantified by an equivalent circle diameter (ECD). In the case of THF steel produced through the traditional hot forming process, the average PAG size was 14.1 μm (Figure 5a₃). However, for the ultrafast heating process with the cold-rolled initial microstructure, when the heating rate reached 10 °C/s, the average PAG size reduced to 9.9 μm (Figure 5a₂) and further refined to 7.8 μm (Figure 5a₁) at the heating rate of 200 °C/s. Regarding grain size distribution, UFH-200 steel exhibited approximately 60% grains with PAG sizes less than 15 μm and no grains larger than 25 μm, which differs from UFH-10 and THF steels, where grains larger than 25 μm were present in specific proportions. Additionally, the deformed ferrite and pearlite morphology in the initial microstructures of UFH-200 and UFH-10 steels were eliminated after full austenitization and replaced by equiaxed austenite and refined martensite.

The TEM micrographs in Figure 6 present the martensite microstructures of the test steels and provide further characterization of the martensite laths. In addition to high-density dislocations within the laths, two intriguing phenomena were observed in the UFH-200 steel: (i) ultra-fine martensite laths with an average width of only 35 nm, as shown in Figure 6a₁ and (ii) martensite laths perpendicular to the main ones within a small area region as depicted in Figure 6a₂. The selected area electron diffraction (SAED) analysis confirms that these two parts are martensite rather than retained austenite. These phenomena, as mentioned above, will be thoroughly analyzed in Section 4.2. The average widths of martensite laths for UFH-200, UFH-10, and THF steels were measured as 121 ± 19 nm, 229 ± 32 nm, and 511 ± 55 nm, respectively.

The dislocation densities of the UFH-200, UFH-10, and THF steels were estimated using the XRD pattern. Figure 7a presents the 2θ-scan profiles of the three test steels and Figure 7b shows the normalized α(110) peaks extracted from the XRD patterns. Compared to the UFH-10 and THF steels, the width of the (110) peak in UFH-200 steel is broader, which corresponds to a higher dislocation density in UFH-200 steels. The estimated dislocation densities of UFH-200, UFH-10, and THF steels using XRD patterns are 6.3 × 10¹⁵ m⁻², 5.0 × 10¹⁵ m⁻², and 4.1 × 10¹⁵ m⁻², respectively.

3.3. Mechanical Properties of Hot Forming Steels

The engineering stress–strain curves of the test steels are presented in Figure 8a, accompanied by a summary table of their mechanical properties. It can be observed from the table that two factors influence the mechanical properties of the test steels, namely the initial structure prior to hot forming and the heating rate. Among these, the THF steel with the fully recrystallized initial microstructure exhibits the lowest mechanical properties. Conversely, the UFH-10 steel with the cold-rolled initial microstructure undergoes full austenitization at a heating rate of 10 °C/s, demonstrating an increase in YS and UTS by approximately 17%, as well as an enhancement in UE and TE by over 50%. For UFH-200 and UFH-10 steels with identical cold-rolled initial microstructures, it is evident that the heating rate significantly affects their mechanical properties. As the heating rate during austenitization increases from 10 °C/s to 200 °C/s, YS substantially rises from 1294 MPa to 1524 MPa, an increase exceeding 17%. Although UTS only slightly improves by 100 MPa, it still reaches a strength level equivalent to 2.2 GPa. Compared to the UFH-10 steel, the UE and TE of the UFH-200 steel only experienced a 5% decrease, indicating that the plasticity of hot forming steel was not significantly reduced despite its extreme increase in strength. In general, when replacing the initial microstructure before austenitization and quenching with the cold-rolled microstructure and increasing the heating rate to 200 °C/s compared to the traditional hot forming process (THF), the test steel exhibited an increase in YS by 424 MPa, UTS by 308 MPa, and UE by approximately 50%. Such an improvement in strength is remarkable for a low alloy content hot forming steel with a whole martensitic structure.

Figure 8b illustrates the true stress–strain curves and strain hardening rate curves of the test steels. The strain hardening rate of the UFH-200 and UFH-10 steel is significantly higher than that of the THF steel since the latter’s UTS is noticeably lower. The strain hardening rate of UFH-200 steel exhibits a significant increase compared to that of UFH-10 steel within the range of true strain ≤ 0.01. However, beyond this threshold, the strain hardening curves for both test steels overlap, indicating that the UFH-200 test steel only demonstrates relatively high strain hardening ability during the initial deformation stage. As deformation progresses, its strain hardening ability becomes equivalent to the UFH-10 steel, explaining the mere 100 MPa difference in their UTS. Furthermore, it is noteworthy that the strain hardening rates for all three test steels do not decay to zero, implying that tensile failure occurs after necking.

4. Discussion

4.1. Influence of the Initial Microstructures

The difference between the cold-rolled initial microstructure of UFH-10 steel and the fully recrystallized initial microstructure of THF steel is described in Section 3.1. When the heating rate remains constant, different initial microstructures also result in distinct austenite transformation behaviors during hot forming, leading to diverse martensite structures after quenching. In the cold-rolled microstructure, recrystallization of the deformed ferrite and spheroidization of layer cementite initially occur during the heating process [30]. This process significantly refines the ferrite grains. Austenite nucleation begins at cementite–ferrite interfaces and ferrite grain boundaries once the temperature reaches the starting point for austenite transformation [31]. As the temperature rises, austenitic nuclei gradually grow while the remaining cementite carbides dissolve. Although numerous cementite carbides provide ample nucleation sites for phase transformation, austenite grains are constrained within ferrites. The average recrystallized ferrite grain size determines austenite grain sizes [32]. Therefore, when the initial microstructure of the test steel is fully recrystallized, despite numerous isolated spherical cementite particles, annealing at 680 °C for an extended period results in a large and uneven ferrite grain size. Consequently, this phenomenon leads to a larger average PAG size in the THF steel. Compared to the UFH-10 steel, there is a significant increase of 42.4% in the average PAG size of the THF steel. Furthermore, it should be noted that the average PAG size plays a crucial role in determining martensitic substructure sizes [33]. Specifically, the average martensitic lath width in the THF steel exceeds that in the UFH-10 steel by more than twice, thus significantly influencing its mechanical properties. This aspect will be further discussed in subsequent sections.

4.2. Influence of the Heating Rate on the Microstructure Evolution

The above results highlight the significant influence of heating rate on the PAG sizes and martensite structure for UFH-200 and UFH-10 steels with a cold-rolled microstructure. Increasing the heating rate from 10 °C/s to 200 °C/s results in a reduction of 21% and 24% in average PAG size and martensite block size, respectively. Additionally, there is a substantial refinement of approximately 47% in the average width of the martensite lath. An analysis was conducted on the dilatation curve (Figure 9a) of the test steels after fully austenitizing at heating rates of 10 °C/s and 200 °C/s without isothermal holding to gain further insights into the impact of the ultrafast heating process on austenite transformation and microstructure refinement. The relative change in length is defined as ΔL/L₀, where ΔL represents length change and L₀ denotes the initial length of the specimen. The temperature point at which a slight slope shift within the dilatation curve defines the start (Ac₁) and finish temperature (Ac₃) for austenite transformation. It can be observed that when the heating rate is increased to 200 °C/s (Figure 9a upper), the Ac₁ and Ac₃ temperatures are elevated compared to a heating rate of 10 °C/s (Figure 9a down), resulting in a broadened temperature range for austenite transformation. This phenomenon aligns with previously reported research [19,34,35]. Figure 9b,c illustrate the calculation of the transformed volume fraction of austenite during continuous heating for UFH-200 and UFH-10 steels, respectively, based on the dilatation curves using the lever law (solid line). The corresponding formation rates are obtained by deriving the volume fraction of transformed austenite as a function of time (dashed line). Notably, the total austenitizing term for UFH-10 steel lasts 8.7 s, while that for UFH-200 steel is only 0.48 s. When the heating rate is 10 °C/s, there is sufficient time for recovery and recrystallization of the cold-rolled microstructure before austenite transformation. Thus, the formation rate for austenite reaches a maximum after approximately 4 s. When the heating rate is increased to 200 °C/s, the cold-rolled microstructure undergoes minimal recrystallization due to the limited time available. During this process, a strong interaction occurs between ferrite recrystallization and austenite transformation, with higher heating rates resulting in faster phase transformation kinetics [36]. The formation rate of austenite not only reaches its maximum rapidly but also exhibits double peaks. The first peak corresponds to the nucleation sites formed by numerous undissolved cementite particles, which quickly grow into pearlite islands and form a carbon-rich austenite. The second peak corresponds to the diffusion-controlled growth of austenite into ferrite grains, forming a carbon-depleted austenite and ultimately forming ultra-fine austenite with a non-uniform carbon distribution [37].

Under ultrafast heating (200 °C/s), the extremely short phase transition duration is insufficient to fully transform the ferrite into austenite. Consequently, a residual volume fraction of approximately 4.9% ferrite was observed in the direct quenching sample after non-isothermal austenitization, accompanied by the presence of a minor quantity of undissolved cementite carbides within the ferrite grains (Figure 10). This unexpected phenomenon serves as a foundation for explaining the formation of martensitic laths in localized regions perpendicular to the primary lath depicted in Figure 6a₂.

The dilatation curves of UFH-200 and UFH-10 steels heated to 900 °C and isothermally held for 20 s are depicted in Figure 11a and Figure 11c, respectively. When the heating rate is 200 °C/s, the specimen exhibits a decrease in dilatation during the holding period, indicating the ongoing transformation of residual ferrite into carbon-depleted austenite. Conversely, when the heating rate is 10 °C/s, full austenitization occurs during the heating stage, and there is no change in dilatation during the isothermal holding stage. The decrease in dilatation of UFH-200 steel suggests that the small amount of carbon-depleted austenite formed during the holding period is surrounded by carbon-rich austenite transformed during the heating period. The austenite transformation during ultrafast heating is reported to be a process of the simultaneous dissolution of ferrite and cementite, controlled by the non-partition local equilibrium (NPLE) mode. In contrast, the transformation of residual ferrite to austenite during the isothermal holding stage is controlled by the partition local equilibrium (PLE) mode [38]. The disparity in the phase transformation modes results in distinct orientations of lath martensite formed after quenching between the austenite transformed during the holding stage and that formed during the heating stage. Consequently, this leads to the formation of martensitic laths perpendicular to the primary lath direction within localized regions.

For the UFH-200 steel, despite full austenitization with the 20 s isothermal holding stage, there is still a lack of time for sufficient diffusion and uniform distribution of carbon within the austenite. Figure 11b,d illustrate martensite volume fractions as a function of time (solid line) calculated based on dilatation curves and lever law during the quenching stage for both UFH-200 and UFH-10 steels. The corresponding formation rates are obtained by deriving the volume fraction of transformed martensite as a function of time (dashed line). Compared to the UFH-10 steel, the martensite formation rate curve of the UFH-200 steel exhibits two peaks, indicating that as quenching progresses, the carbon-depleted austenite with a higher Ms temperature preferentially undergoes phase transformation into martensite before carbon-rich austenite with a lower Ms temperature. These findings confirm our previous analysis regarding uneven C distribution within the austenite of UFH-200 steel even after an isothermal holding treatment at 900 °C for 20 s. Figure 12 illustrates the schematic diagram of phase transformation in UFH-200 and UFH-10 steels, considering the cold-rolled microstructure as the initial one.

4.3. Evolution of Mechanical Properties

The variation in YS and UTS of the test steels with heating rate and different initial microstructures is illustrated in Figure 8a. Generally, the UFH-200 steel exhibits a significantly higher YS (1524 MPa) compared to UFH-10 (1294 MPa) and THF (1100 MPa) steels, primarily attributed to the synergistic effect of dislocation strengthening and grain refinement strengthening mechanisms. The dislocation strengthening increments of the test steels, according to the Taylor hardening law (${\mathsf{\sigma }}_{\mathrm{d}}=\mathrm{M}\mathsf{\alpha }\mathrm{Gb}\sqrt{\mathsf{\rho }}$), are estimated as 1089.7 MPa, 968.4 MPa, and 877.2 MPa, respectively, where M is the Taylor factor, taken as three; α is a constant, α = 0.24 [39]; b is the magnitude of the Burgers vector, taken as 0.25 nm [40]; G is the shear modulus, taken as 80 GPa; and ρ is the density of the dislocations.

Meanwhile, the YS contributed by the grain refinement strengthening mechanism is calculated according to the Hall–Patch formula (${\sigma }_g=k_{y}d$^−1/2), where k_y is the Hall–Patch coefficient, which is 120 MPa·μm^−1/2 for martensitic steel [41] and d is the average martensite lath width. The fine grain strengthening increments of the test steels are 345.0 MPa, 250.8 MPa, and 167.9 MPa, respectively.

It can be observed that the transition of the initial microstructure from a fully recrystallized structure to a cold-rolled structure and the increase in the heating rate to 200 °C/s resulted in the high-density dislocations in martensite after the hot forming process. Furthermore, significant refinement is also observed in the martensite lath, leading to a substantial increase in YS for the UFH-200 and UFH-10 steels compared to THF steels. Combining the analysis of microstructure evolution discussed in the above sections with the quantitative calculation of strengthening mechanisms clarifies that an increase in heating rate up to 200 °C/s significantly enhances the YS of hot forming steel, reaching strength levels as high as 1.5 GPa.

Compared to the initial microstructure, the effect of the heating rate on the UTS of the test steels is less significant, which is primarily due to the significantly lower average PAG sizes in UFH-200 and UFH-10 steels compared to THF steel, indicating that refining PAG greatly improves the UTS of martensitic steel [42]. The strain hardening rate curves in Figure 8b also demonstrate that THF steel exhibits a lower strain hardening ability during the overall stage of tensile deformation, which aligns with its lower UTS. Previous studies [11,43] have suggested that synergistic strengthening between micro-alloyed precipitates and high-density dislocations effectively enhances the strain hardening ability of ultra-high-strength steels. However, in this study, due to the lower solution temperature of V-containing carbonitrides [44], they dissolved entirely during the austenitizing process. Therefore, no synergistic strengthening exists between precipitates and high-density dislocations in UFH-200 and UFH-10 steels, which is confirmed by the overlap of the strain hardening curve for UFH-200 and UFH-10 steels. Consequently, increasing the heating rate from 10 °C/s to 200 °C/s only leads to a UTS increase of 100 MPa, which is solely attributed to the relative refinement of average PAG size.

5. Conclusions

In this study, ultrafast heating was utilized in the hot forming process and the impact of the ultrafast heating process on the microstructure evolution and mechanical properties of the test steel was thoroughly analyzed while accurately quantifying the contributions of dislocation strengthening and grain refinement strengthening mechanisms to the YS, which substantiates an alternative promising method for producing hot forming steel with excellent mechanical properties besides the traditional process. The conclusions are as follows:

1. The ultrafast heating process (200 °C/s) significantly enhanced the kinetics of austenite transformation compared to the traditional hot forming process (10 °C/s), which resulted in the formation of ultrafine PAGs with a heterogeneous distribution of carbon, leading to a reduction in the average width of martensite lath by approximately 47%. Consequently, there was an increase of about 100 MPa in the contribution of the grain refinement strengthening mechanism to yield strength.

2. Compared to the fully recrystallized initial microstructure, the strong interaction between austenite nucleation and ferrite recrystallization existed in the cold-rolled initial microstructure during the ultrafast heating process, thereby preserving high-density dislocations in PAG and as-quenched martensite. The contribution of the dislocation strengthening mechanism to the YS of UFH-200 steel exceeded 1000 MPa.

3. The ultrafast heating process significantly enhanced the YS and UTS of hot forming steel while also ensuring a certain level of plasticity. Compared with the traditional hot forming process, ultrafast heating significantly enhanced the YS and UTS of hot forming steel by 38.5% and 22.5%, respectively. Ultimately, this results in excellent mechanical properties with YS of 1524 MPa, UTS of 2221 MPa, and UE of 5.2%.