Effects of Tempering Temperature on the Microstructure and Mechanical Properties of Vanadium-Microalloyed Medium-Carbon Bainitic Steel

Abstract

This study examined the impact of tempering temperature on the microstructure and properties of vanadium (V)-microalloyed medium-carbon bainitic steel. A series of heat treatments were performed on the steel, and the microstructural evolution and mechanical properties were systematically investigated through X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and mechanical testing systems (MTS). The findings revealed that tempering temperature has a significant influence on microstructural changes. Specifically, at 350–450 °C, retained austenite begins to decompose and carbides start to precipitate. At 550–600 °C, bainitic ferrite laths undergo coarsening. Regarding mechanical properties, both tensile strength and yield strength initially increase with tempering temperature before decreasing as the temperature continues to rise. The diffusion and redistribution of carbon atoms during tempering enhance the elongation of all tempered samples compared to their untempered counterparts. Optimal comprehensive mechanical properties are achieved at 450 °C, where precipitation strengthening from vanadium, enhanced stability of retained austenite, and synergistic strengthening effects of decomposition products are most pronounced. This research provides a theoretical foundation for optimizing the heat treatment process of such steels and offers insights into the synergistic effects of V-microalloying and tempering.

1. Introduction

In the field of materials science, extensive research has been dedicated to optimizing steel properties for engineering applications [1,2,3,4,5]. Vanadium microalloying and tempering processes of bainitic steel have remained focal points of investigation. Vanadium microalloying has been widely employed in medium-carbon bainitic steels to significantly modify their microstructure and mechanical properties. Feng et al. demonstrated that vanadium microalloying enhances strength of the steel through dispersion strengthening and austenite grain refinement via the formation of nano-sized V(C,N) precipitates [6]. Similar conclusions were also reported in studies by Zhang et al. and Lu et al. [7,8]. However, in bainitic steels, the V(C,N) precipitation is influenced by multiple factors including cooling rate and bainitic transformation temperature, leading to variations in precipitation quantity and strengthening effectiveness. As revealed in Wang et al.’s study, while vanadium addition can refine M/A constituents and inhibit proeutectoid ferrite formation, its precipitation-strengthening effect remains limited [9]. Although vanadium addition enhances steel strength to varying degrees [10], improvements in toughness still present challenges.

The tempering process of bainitic steel has similarly attracted significant attention. Tempering exerts crucial influences on both the microstructure and properties of bainitic steels. Variations in tempering temperature induce complex microstructural changes including decomposition of retained austenite, precipitation and coarsening of carbides, and modifications in bainitic ferrite laths, thereby significantly affecting mechanical properties such as strength, toughness, and hardness [11,12,13,14]. Kang et al. revealed that appropriate tempering processes can further enhance mechanical properties of bainitic steel, while excessive tempering temperatures deteriorate the strength–ductility balance [15]. Similar phenomena were demonstrated by Wang et al. and Guo et al. [16,17]. The stability evolution of retained austenite under different tempering temperatures differentially impacts comprehensive properties, while carbide precipitation behavior remains closely associated with steel strengthening mechanisms [18,19].

Although significant progress has been made in the study of V-microalloyed medium-carbon bainitic steel, several issues remain unresolved. The current understanding of microstructural evolution and property changes in V-microalloyed medium-carbon bainitic steel at different tempering temperatures is still inadequate [20]. In particular, the synergistic effects of V microalloying and tempering as well as the control of V precipitation and retained austenite transformation during tempering to achieve an optimal balance of strength, toughness, and hardness require further investigation. This study aimed to address this research gap by comprehensively examining the influence of tempering temperature on the microstructure and mechanical properties of V-microalloyed medium-carbon bainitic steel, thereby providing theoretical foundation and data support for optimizing the heat treatment process of such steels.

2. Materials and Methods

2.1. Experimental Materials

This work focused on a V-microalloyed medium-carbon bainitic steel with nominal composition Fe-0.42C-1.17Mn-1.15Si-0.96Cr-0.73Mo-0.56Al-0.12V. The steel was produced through vacuum induction melting, homogenized, and thermomechanical processing via forging into 60 × 60 mm cross-sectional bars.

The phase-transformation points Ac₁ and Ac₃ were measured using a DIL 402 dilatometer (Netzsch, Germany) under controlled heating conditions (5 °C/min). Figure 1a shows the original dilatation curve during continuous heating, and Figure 1b shows the first derivative curve, i.e., d(ΔL)/dt-T curve, which better reflects the changes in dilatation. The phase-transformation points Ac₁ and Ac₃ were determined to be 740 °C and 840 °C, respectively.

The solved temperature for complete V dissolution in the austenite matrix in equilibrium state was calculated using the solubility product formula for V and C [21]:

Lg([V]r·[C]r) = 6.72 − 9500/T

(1)

where [V]r is the content of V dissolved in the matrix (wt.%), [C]r is the content of C dissolved in the matrix (wt.%), and T is the temperature in Kelvin (K). The austenitizing temperature was determined to be 925 °C.

The Ms temperature and bainitic isothermal transformation curve were measured using a DIL 805A/D dilatometer(Linseis, Selb, Germany). The samples were heated to 925 °C with a controlled heating rate of 10 °C/s, held for 30 min, and then cooled to room temperature at 30 °C/s. The temperature–dilatation curve is shown in Figure 1c. The Ms temperature was identified as 288 °C using the tangent method. Similarly, the bainitic isothermal transformation curve was obtained by cooling the sample to 300 °C at 30 °C/s after holding at 925 °C for 30 min, as shown in Figure 1d. The bainitic transformation was completed in approximately 2 h without an incubation period.

The bainitic microstructure was obtained through isothermal quenching. The heat treatment parameters were as follows: heating to 925 °C for complete austenitization, holding for 30 min, isothermal holding in a salt bath at 300 °C for 2 h, and tempering at 350 °C, 450 °C, 550 °C, and 600 °C for 1 h, followed by subsequent cooling to room temperature.

2.2. Microstructural and Mechanical Property Tested

The microstructures of the samples were observed using scanning electron microscopy (SEM), transmission electron microscopy (TEM), and energy-dispersive spectroscopy (EDS). SEM samples were mechanically ground and polished and then etched with 4% nitric alcohol. TEM samples were precision-sectioned to 0.5 mm thick slices, ground to ~50 μm, and thinned using a twin-jet electropolishing device with 7% perchloric acid alcohol. The phase composition was analyzed via a SmartLab X-ray diffractometer (Rigaku, Tokyo, Japan) with a Co target, at a scanning speed of 3°/min and a scanning angle (2θ) of [40°, 130°]. The content of retained austenite (RA) was quantified according to the multi-peak integrated intensities of the (110)α, (200)α, (211)α, (220)α, and (111)γ, (200)γ, (0)γ, and (311)γ peaks [22].

$$V_{\mathsf{\gamma }}={\displaystyle \frac{(\frac{1}{n}){\sum }_{j=1}^n\frac{I_{\mathsf{\gamma }}^j}{R_{\mathsf{\gamma }}^j}}{\left(\frac{1}{n}\right){\sum }_{j=1}^n\frac{I_{\mathsf{\gamma }}^j}{R_{\mathsf{\gamma }}^j}+(\frac{1}{n}){\sum }_{j=1}^n\frac{I_{\mathsf{\alpha }}^j}{R_{\mathsf{\alpha }}^j}}}$$

(2)

$$R=\left(1/{\mathsf{\nu }}^2\right)\left[{\left|F\right|}^{2}P\left(\left(1+{cos}^{2}2\theta \right)/sin\theta sin\theta \right)\right]e^{-2M}$$

(3)

where V_γ is the content of austenite, I is the diffraction peak intensity, n is the effective number of diffraction peaks for calculation, R is the material scattering factor, ν is the unit cell volume, F is the structure factor, P is the multiplicity factor, and e^−2M is the temperature factor.

Tensile tests were conducted on samples subjected to different tempering processes using an MTS810 universal testing machine (MTS Systems, Eden Prairie, MN, USA) with a certain strain rate (3 mm/min). The actual strain was acquired via an extensometer. The gauge length and diameter of the standard cylindrical specimen for tensile tests were 25 mm and 5 mm, respectively. Three tensile specimens were prepared for each tempering process to ensure accuracy.

3. Results and Discussion

3.1. Microstructure

Figure 2 presents the SEM images of samples tempered at various temperatures. No martensitic structure was observed in any of the tempered samples, indicating that no martensitic transformation occurred during air cooling after tempering. The microstructure of the untempered sample (Figure 2a) comprised bainitic ferrite (BF) and RA, with no carbides detected. At 350 °C and 450 °C (Figure 2b,c), the retained austenite began to decompose, and carbides started to precipitate. At 550 °C and 600 °C (Figure 2d,e), the BF laths fragmented, and their boundaries became indistinct. The characteristic bainitic structure gradually disappeared, and a significant amount of fine granular carbides precipitated and coarsened. Within the tempering temperature range of 350–600 °C, the microstructural evolution involved the progressive decomposition of RA and BF, accompanied by the precipitation and coarsening of carbides as the tempering temperature increased. In addition, Figure 2f shows the EDS analysis of the 550 °C-tempered sample. A large number of carbides were observed, and point analysis confirmed the presence of V-containing particles at the nanoscale.

Figure 3 shows the TEM images of samples tempered at 450 °C and 550 °C. The microstructure consisted of lath-like BF and RA, with a high density of dislocations in the BF laths. The RA was mostly film-like and distributed between the laths.

Figure 4a shows the microstructure of the 450 °C-tempered sample. EDS mapping (Figure 4b) confirmed the presence of V-containing particles, approximately 40 nm in size and rhombus-shaped. The precipitation of these particles contributed to second-phase strengthening, enhancing the tensile strength of the steel.

Figure 5a shows the XRD patterns of samples tempered at different temperatures. The content of RA was calculated using Equations (2) and (3), as shown in

Table 1. Volume fraction of RA after tempering at different temperatures.

Temperature/°C	Untempered	350	450	550	600
VRA/%	12.95	10.37	12.87	3.75	5.02

. At lower tempering temperatures, the content of RA changed little, as carbon atoms in BF continued to diffuse into RA, enhancing its thermal stability. At higher tempering temperatures, the RA content decreased significantly due to the carbides precipitation, which consumed carbon in the RA, reducing its thermal stability and leading to its decomposition into cementite and ferrite. Figure 5b exhibits the (200)γ peak of austenite, with an inset showing an enlarged view. The (200)γ peak shifted to the left at 350 °C and 450 °C, indicating an increase in carbon content and stability of RA. At 550 °C and 600 °C, the (200)γ peak shifted to the right, indicating a decrease in carbon content and stability of RA.

3.2. Mechanical Properties

Figure 6 gives the engineering stress–strain curves and relationship between mechanical property parameters and tempering temperature. None of the tensile curves displayed a distinct yield point, and the most significant work hardening rate was observed during the early stage of uniaxial deformation. Detailed mechanical properties are summarized in

Table 2. Mechanical properties of samples subjected to different heat treatments.

Heat Treatment °C	σ0.2, MPa	σb, MPa	δgt %	δ %	PSE, GPa%
Untempered	1289	1680	3.3	10.8	18.1
350	1389	1752	6.2	12.7	22.3
450	1431	1765	4.0	12.0	21.2
550	1153	1643	8.6	13.1	21.5
600	1176	1615	7.7	12.1	19.5

. The elongation of tempered specimens was notably higher compared to the untempered sample. During tempering, carbon atoms diffused from BF into RA, reducing dislocation density and relieving residual stresses within the BF. These mechanisms collectively contributed to the enhanced elongation of the tempered samples. As tempering temperature increased, both tensile strength and yield strength initially rose before subsequently declining. At 350 °C, the carbon content in RA increased, enhancing its mechanical stability and leading to elevated yield and tensile strengths. The peak tensile and yield strengths were achieved at 450 °C, reaching 1765 MPa and 1431 MPa, respectively. This can be attributed to the gradual precipitation of nano-sized V-containing particles (as shown in Figure 4), which increased carbon content and stability of RA. At 550 °C, the tensile strength approximated that of the untempered sample, while at 600 °C, both tensile and yield strengths fell below those of the untempered sample. In summary, the 450 °C-tempered sample demonstrated the optimal balance of mechanical properties.

Figure 7 shows the Rockwell hardness of samples tempered at different temperatures. The hardness of untempered sample increased gradually up to 450 °C, reaching a maximum HRC of 50. At higher tempering temperatures, the hardness decreased significantly, with the lowest hardness of HRC 47.9 observed at 600 °C. The increase in hardness at 350–450 °C was primarily due to the precipitation of V-containing particles, which contributed to second-phase strengthening. At 550 °C, owing to the decomposition of filmy and blocky RA, the hardness decreased significantly, forming a mixture of cementite and ferrite, which reduced the mechanical stability of RA. Additionally, the dislocation density in BF decreased with increasing tempering temperature, further reducing hardness.

3.3. Relationship Between Tempering Temperature and Microstructure

During tempering, microstructural evolution primarily involves the decomposition of RA, the precipitation and growth of V-containing carbides, and changes in BF laths. According to XRD analysis (Figure 5), at lower tempering temperatures (350 °C and 450 °C), the (200)γ peak of RA shifted to the left, indicating an increase in carbon content and stability [23]. This is because carbon diffusion in RA is relatively slow at these temperatures, and the high Si content in the steel inhibits the precipitation of cementite, resulting in a greater driving force for bainitic transformation and preferential decomposition of RA into bainite [24]. At this stage, the decomposition of BF and RA is not significant, and carbides begin to precipitate gradually.

At higher tempering temperatures (550 °C and 600 °C), the (200)γ peak shifted to the right, indicating a decrease in carbon content and stability of RA. This is because carbon diffusion in RA accelerates at higher temperatures, and the driving force for cementite precipitation exceeds that for bainite formation, leading to increased decomposition of RA. SEM images (Figure 2) show that at 550 °C and 600 °C, BF laths begin to fragment, and the boundaries become blurred. The typical bainitic structure gradually disappears, and a large number of fine granular carbides precipitate and coarsen. During this process, film-like austenite decomposes into ferrite with the same orientation as the surrounding bainite and discrete cementite, while blocky RA decomposes into a mixture of cementite and equiaxed ferrite.

3.4. Relationship Between Microstructure and Mechanical Properties

Precipitation strengthening: TEM and EDS analysis revealed the precipitation of nano-sized V-containing particles at 450 °C. V can form fine carbonitride particles in steel, contributing to precipitation strengthening. The precipitation of these particles enhances second-phase strengthening, improving the tensile strength of the steel [13]. At 550 °C, although V-containing particles also precipitate, the decomposition of RA and microstructural changes have a more significant impact on mechanical properties, leading to complex strength variations.

Microstructure composition and stability of retained austenite: The stability of RA changes with tempering temperature, affecting mechanical properties. At lower tempering temperatures, RA exhibits higher stability. During tensile deformation, blocky RA with low mechanical stability undergoes the TRIP effect and transforms into martensite at small strains, while film-like austenite with high stability transforms into martensite at large strains, delaying necking and increasing elongation compared to untempered samples [25,26,27].

At 450 °C, the carbon content in RA increases, enhancing its stability. The precipitation of V-containing particles and the increased stability of RA synergistically improve tensile and yield strengths, achieving a good balance of strength and toughness. At 550 °C and 600 °C, the decomposition of RA increases, forming a mixture of cementite and ferrite. The formation of cementite alters the stress distribution and deformation behavior of the microstructure. On one hand, the ability of the sample to coordinate deformation increases, improving elongation. On the other hand, the carbon content in RA decreases, reducing its stability, and bainitic laths coarsen, leading to a decrease in strength. For example, at 600 °C, both tensile and yield strengths are lower than those of the untempered sample despite the increased precipitation of V-containing particles, as the strength loss due to microstructural coarsening outweighs the strength gain from precipitation strengthening.

4. Conclusions

• At 350–450 °C, RA begins to decompose, and carbides start to precipitate, while the BF lath structure remains relatively stable. At 550–600 °C, BF laths coarsen, and the boundaries become blurred. The typical bainitic structure gradually disappears, and a large number of fine granular carbides precipitate and coarsen. Higher tempering temperatures promote the decomposition of RA and the precipitation of carbides, reducing the stability of the bainitic structure;
• The tensile and yield strengths of tempered samples initially increase and then decrease with increasing tempering temperature. At 350–450 °C, both yield and tensile strengths increase due to the precipitation of nano-sized V-containing particles and the increased carbon content and stability of RA. At 550–600 °C, although V-containing particles continue to precipitate, the decomposition of RA and other factors lead to a decrease in strength. The elongation of tempered samples is higher than that of untempered samples due to the diffusion of carbon atoms from BF to RA and the reduction in dislocation density and retained stress in BF;
• The V-microalloyed medium carbon bainitic steel tempered at 450 °C exhibits the best combination of mechanical properties. At this temperature, the precipitation strengthening of V-containing particles, the increased stability of RA, and the synergistic effects of decomposition products achieve an ideal balance of strength, toughness, and hardness.