The Effect of Quenching and Tempering Temperatures on the Microstructure and Properties of a New Low-Alloy Ultra-High-Strength Martensitic Steel

Abstract

This study systematically investigates the influence of quenching (850–910 °C) and tempering (160–280 °C) temperatures on the microstructural evolution and mechanical properties of a novel low-alloy ultra-high-strength martensitic steel (UHSMS). Comprehensive microstructural characterization combined with mechanical testing demonstrates that quenching at 880 °C results in the finest martensitic laths and the highest dislocation density, leading to an excellent strength–toughness balance. Subsequent tempering treatments reveal that the specimen tempered at 200 °C achieves an optimal combination of properties, with a yield strength of 1517 MPa, ultimate tensile strength of 2017 MPa, elongation of 10.4%, and impact toughness of 80.3 J/cm². This optimum is mechanistically linked to a cooperative effect where the fine tempered martensitic structure and stable film-like retained austenite (RA) enhance toughness and ductility, while the nano-scale precipitates (forming during the ε→θ carbide transition) simultaneously provide substantial precipitation strengthening, thereby minimizing the strength sacrifice typically associated with improved toughness. Furthermore, the 200 °C tempered specimen exhibits the largest shear lip on the tensile fracture surface and the maximum dimple size on the impact fracture surface, indicative of a high plastic strain capacity and excellent crack propagation resistance.

1. Introduction

Ultra-high-strength martensitic steel (UHSMS) serve as critical structural materials in aerospace and defense applications due to their exceptional mechanical properties [1,2]. While commercial grades such as 300M, AF1410, and Aermet 340 offer high strength, their fracture toughness remains inadequate for damage-tolerant design requirements [3,4]. Recent studies have demonstrated that microalloying with elements such as Nb, V, Ni, and Mo, combined with optimized heat treatment, can significantly enhance their strength–toughness balance [5,6,7]. For example, Hojo et al. [8] reported that the addition of 1.0% Cr to TRIP-assisted martensitic steel promotes the stabilization of retained austenite (RA) films along martensite laths, thereby suppressing crack propagation and improving fracture resistance.

Drawing on the composition of 30Si2MnCrMoVE—a Chinese-developed medium-carbon low-alloy ultra-high-strength steel used in solid rocket motor casings [9,10]—this study proposes a new alloy design strategy that is deliberately differentiated from this reference grade. Specifically, we intentionally increased the Cr content to 2.0 wt.% to further improve hardenability and promote the refinement of martensitic substructures. Additionally, a small but controlled addition of Nb (0.02 wt.%) was introduced to exploit synergistic precipitation strengthening, particularly during low-temperature tempering. The Ni content was also adjusted to enhance RA stability and toughness [11,12,13].

The mechanical behavior of UHSMSs is governed by the combined influence of multiple microstructural features, including prior-austenite grain size, multi-scale martensitic constituents (packets, blocks, and laths), precipitation characteristics, and the volume fraction and thermal stability of retained austenite [14,15]. Through systematic optimization of the heat-treatment parameters, precise microstructural control can be achieved. Quenching temperature primarily governs prior-austenite grain growth and martensite substructure formation, while tempering temperature regulates carbide precipitation and dislocation recovery, thereby establishing an optimal strength–ductility–toughness synergy [16,17]. In Nb–V–Ti microalloyed steels, Dong et al. [18] reported that elevated tempering temperatures promote martensite lath coarsening, yielding a marginal strength sacrifice for substantially enhanced impact toughness. Wang et al. [19] demonstrated that increasing quenching temperatures for 1.5Cr-0.3Mo-0.3Ni steel accelerates prior-austenite grain growth and compromises microstructural homogeneity, consequently reducing both strength and wear resistance. Their work further revealed that high-temperature tempering induces martensite lath boundary annihilation and carbide coarsening. Shi et al. [20] identified a critical tempering threshold at 482 °C for AerMet100 steel, beyond which rapid deterioration of mechanical properties occurs, correlated with the reduction in dislocation density within martensite laths. Collectively, these studies confirm that heat-treatment optimization is a fundamental strategy for enhancing the performance of martensitic steels. An optimal microstructure achieving a superior strength–toughness synergy comprises refined tempered martensite laths, nano-scale inter-lath RA films, and homogeneously distributed fine carbides [21].

Despite these advances, systematic investigations that simultaneously decouple the individual and interactive effects of quenching and low-temperature tempering on microstructural evolution and mechanical properties—particularly within the low-tempering regime (160–280 °C)—remain limited. Furthermore, quantitative linkages between specific microstructural features (e.g., martensite block size, dislocation density, ε→θ carbide transition kinetics) and the corresponding strengthening and toughening mechanisms are still insufficiently established for Cr-containing medium-carbon low-alloy UHSMSs. To address these gaps, the present study systematically investigates the effects of the quenching (850–910 °C) and low-temperature tempering (160–280 °C) parameters on the microstructural evolution and mechanical properties of a newly designed Cr-containing medium-carbon low-alloy martensitic steel. The objectives are to: (i) identify the optimal heat-treatment regime for achieving the best strength–toughness combination; (ii) elucidate the underlying microstructural mechanisms governing the strength–toughness synergy; and (iii) provide theoretical guidance for the development and application of low-alloy ultra-high-strength martensitic steels.

2. Experimental Section

2.1. Materials

The investigated steels were prepared using a 25 kg vacuum arc melting furnace (Jinzhou Electric Furnace Co., Ltd., Jinzhou, China). To ensure high cleanliness of the experimental steels, raw materials with low impurity contents were melted in a MgO crucible at 1600 °C. After melting, the liquid steel was cast into a mold manufactured using gray cast iron to produce a shell ingot. Its chemical compositions are provided in

Table 1. Chemical composition design of novel low-alloy steel (wt.%).

C	Mn	Si	Cr	Mo	Cu	Ni	O	Nb	Fe
0.42	1.02	1.04	2.00	0.31	0.003	1.06	0.0008	0.02	Bal.

. The critical phase transformation temperatures were determined by laser confocal scanning microscopy (VL2000DX-SVF17SP, Lasertec Co., Yokohama, Japan) and dilatometry (DIL402, Netzsch, Selb, Germany). The austenite-start (Ac₁) and austenite-finish (Ac₃) temperatures were measured as 730 °C and 800 °C, respectively. Guided by these findings, samples were heated to 850 °C, 880 °C, and 910 °C and held for one hour to ensure full austenitization, then rapidly cooled in oil to obtain entirely martensitic structures. Microstructural characterization and mechanical testing were performed to identify the optimum quenching condition. Subsequently, specimens subjected to the selected quenching treatment were tempered at 160 °C, 200 °C, 240 °C, and 280 °C for 2 h to systematically evaluate the tempering response. The complete heat-treatment schedule is shown in Figure 1. The resulting specimens are labeled as ‘Q850...’, where ‘Q’ denotes quenching, and the number indicates the austenitizing temperature. The tempered specimens are labeled as ‘T160...’, where ‘T’ denotes tempering and the number indicates the tempering temperature.

2.2. Microstructure Characterization

Samples etched with 4% nital for 10–15 s were seen under a Zeiss Axiolab 5 microscope for optical microscopy (OM, Zeiss, Oberkochen, Germany). The microstructural morphology was revealed using backscattered electron (BSE) imaging on a Phenom Pro scanning electron microscope (SEM, Phenom-World, Eindhoven, The Netherlands) running at 15 kV. Before being examined in an SEM equipped with an Oxford Nordlys max3 detector (Oxford Instruments, Abingdon, UK), the specimens prepared for electron backscatter diffraction (EBSD) were electrolytically polished using a 5 vol% perchloric acid in ethanol solution. Dislocation morphology was observed using a Thermo Fisher Scientific Themis Z spherical aberration-corrected transmission electron microscopy (TEM, Talos F200X, Thermo Fisher Scientific, Waltham, MA, USA) analysis. The equivalent circle diameter was calculated from the measured grain area. To ensure the representativeness of the microstructural analysis, characterization specimens were sectioned from the central region of the heat-treated billets. For each condition, observations were conducted across multiple fields of view using each characterization technique. Phase analysis was performed using Co-Kα radiation and X-ray diffraction (XRD, Panalytical, Almelo, The Netherlands) at a scan rate of 5°min⁻¹. Using the formulas provided in Equations (1) and (2) [22,23], the volume percent of RA was computed from the integrated intensities of the (200)α, (211)α, and (200)γ diffraction peaks.

$$V_i={\displaystyle \frac{1}{1+G(I_{\alpha }+I_{\gamma })}}$$

(1)

$$C_Y=\left({\displaystyle \frac{\lambda \sqrt{h^2+k^2+l^2}}{2\mathrm{s}\mathrm{i}\mathrm{n}\theta }}-3.578\right)/0.033$$

(2)

2.3. Mechanical Properties Testing

Instron 3369 universal testing equipment (Instron, Norwood, MA, USA) was used to determine the tensile parameters under the different processing circumstances, including yield strength (YS), ultimate tensile strength (UTS), and elongation (A). Standard U-notched specimens were subjected to Charpy impact tests at room temperature. Single-edge notched bend (SEB) specimens were used to assess fracture toughness (K_IC). Vickers microhardness was measured using a 1 kg force and a dwell duration of 15 s. Figure 2 shows the dimensions of the specimen.

3. Results

3.1. Mechanical Properties

Table 2. Mechanical properties at different quenching temperatures.

Sample ID	YS (MPa)	UTS (MPa)	A (%)	CUN (J/cm2)	Hardness (HV)
Q850	1211.6 ± 18.5	2224.3 ± 29.4	5.0 ± 0.2	20.8 ± 0.9	625.3 ± 8.6
Q880	1144.9 ± 18.2	2308.8 ± 33.4	6.2 ± 0.4	23.0 ± 1.3	645.8 ± 7.0
Q910	1080.4 ± 13.9	2201.2 ± 27.6	3.9 ± 0.2	22.9 ± 1.0	637.6 ± 10.4

compiles the mechanical performance of the steel after quenching across the applied temperature range. Relative to the conventional 30Si2MnCrMoVE steel, the developed alloy demonstrates a markedly superior strength–toughness synergy, attaining a UTS above 2200 MPa while retaining a room-temperature Charpy impact energy greater than 20 J/cm². A gradual decline in yield strength is observed when the quenching temperature increases from 850 °C to 910 °C. Conversely, quenching at 880 °C results in peak values for ultimate tensile strength, elongation, impact energy, and hardness.

Table 3. Mechanical properties at different tempering temperatures.

Sample ID	YS (MPa)	UTS (MPa)	A (%)	CUN (J/cm2)	KIC (KJ/m2)	Hardness (HV)
T160	1614.6 ± 17.2	2155.1 ± 26.0	10.3 ± 0.4	72.5 ± 2.6	57.2 ± 1.1	584.4 ± 6.6
T200	1517.1 ± 22.7	2017.5 ± 20.2	10.4 ± 0.3	80.3 ± 0.9	66.0 ± 0.6	583.6 ± 13.1
T240	1515.5 ± 19.1	1924.1 ± 18.9	9.9 ± 0.2	77.5 ± 1.5	66.7 ± 0.8	579.1 ± 7.6
T280	1486.7 ± 18.3	1859.5 ± 19.7	9.7 ± 0.2	68.6 ± 1.4	67.0 ± 1.5	572.7 ± 9.2

summarizes the mechanical performance of the specimens after quenching at 880 °C and tempering at various temperatures. Tempering greatly increases the YS, A, and impact energy in comparison to the as-quenched state, whereas the UTS is somewhat decreased. With an increase in tempering temperature from 160 °C to 280 °C, a progressive reduction in yield strength, tensile strength, and hardness is accompanied by a steady enhancement of fracture toughness. At 200 °C, tempering causes impact energy and A to reach their maximum values. In conclusion, tempering at 200 °C yields the best overall performance, maintaining a UTS above 2000 MPa while providing exceptional ductility and toughness.

3.2. Microstructure

3.2.1. Microstructure After Quenching

Figure 3 illustrates the OM, SEM, and EBSD analyses of the experimental steel subjected to varying quenching temperatures. The microstructure predominantly consists of quenched martensite, and its size progressively increases with higher quenching temperatures. This is primarily due to the coarsening of prior-austenite grains at elevated quenching temperatures (as shown in Section 4.1), which provides greater accommodation space for the formation of martensite variants within each austenite grain. The corresponding effective grain areas are measured as 2.06 μm², 4.95 μm², and 5.43 μm², with equivalent circle diameters of 1.33 μm, 1.89 μm, and 2.66 μm, respectively. Figure 4a displays the XRD patterns of specimens at different quenching temperatures. All patterns exhibit diffraction peaks corresponding to body-centered cubic (BCC) martensite/ferrite—(110)α, (200)α, and (211)α—accompanied by relatively weak peaks from face-centered cubic (FCC) RA, notably the (220)γ peak. No diffraction peaks corresponding to impurity phases such as carbides are detected. The RA content and its carbon concentration were calculated using Equations (1) and (2) based on diffraction peak positions and integrated intensities measured with Jade software (version 9.0). As shown in Figure 4b, the RA content increases marginally with quenching temperature, reaching a maximum of 5.9% at 910 °C, while the carbon concentration in RA remains relatively stable. During quenching, the carbon concentration in RA remains relatively stable because within this temperature range the dominant effect is an increase in the volume of stabilized austenite, not a significant shift in the equilibrium carbon partitioning level dictated by the fixed bulk alloy composition.

Figure 5 presents a systematic TEM analysis of specimens quenched at different temperatures, revealing a clear dependence of martensitic morphology on processing conditions. The dimensions of the martensite packets and blocks scale with increasing quenching temperature. At 850 °C, the microstructure exhibits martensite laths measuring approximately 0.23 μm in width, accompanied by discernible undissolved carbide particles (highlighted by yellow circles in Figure 5a). Quenching at 880 °C produces refined laths (0.14 μm wide) containing dense dislocation tangles (Figure 5b). When the temperature is raised further to 910 °C, the martensite laths widen (to approximately 0.36 μm) and their interfacial definition becomes progressively less distinct. It is observed that the specimen quenched at 880 °C exhibits the most refined lath martensite structure and simultaneously the highest dislocation density among all conditions studied.

Under all of the settings investigated, the inverse pole figure (IPF) maps in Figure 6a–c, which correspond to the three quenching temperatures, show no strong crystallographic texture. The low-angle grain boundaries (LAGBs, 2–15°, red) and high-angle grain boundaries (HAGBs, 15–180°, blue) are distinguished on related grain boundary maps (Figure 6a–c). Based on quantitative analysis, the fractions of HAGBs measured in specimens quenched at 850 °C, 880 °C, and 910 °C are 73.1%, 69.1%, and 74.8%, respectively; the complementary LAGB proportions are 26.9%, 30.9%, and 25.2%. Interestingly, the specimen that was quenched at 880 °C has the largest percentage of LAGBs. Prior-austenite grain boundaries and martensite packet boundaries are examples of HAGBs that effectively impede crack propagation in martensitic steels, increasing fracture toughness [24,25]. In contrast, LAGBs (such as the boundaries between martensite laths) possess a comparatively simple structure and a lower interfacial energy, generally providing less resistance to crack advance, although an appropriate population of such boundaries can contribute to strength. The LAGB content initially rises with increasing quenching temperature, reaching a maximum at 880 °C before declining. This non-monotonic trend is attributed primarily to the coarsening of prior-austenite grains and martensite packets at higher temperatures, together with the finest martensite lath size obtained specifically at 880 °C, which collectively promote a higher density of low-angle interfaces.

3.2.2. Microstructure After Tempering

Figure 7 displays the OM, SEM, and EBSD analyses of the steel tempered at various temperatures. The microstructure primarily consists of tempered martensite along with a minor fraction of RA. Based on the measurements from Figure 7c–l, the effective grain areas of the tempered steel are 5.97 μm², 6.00 μm², 6.17 μm², and 8.34 μm², corresponding to equivalent circle diameters of 1.90 μm, 1.93 μm, 1.98 μm, and 2.27 μm, respectively. These data indicate a modest coarsening of the tempered martensite blocks with rising tempering temperature. At 160 °C, the martensite lath boundaries appear well-resolved; with further elevation of the tempering temperature, these interfaces become increasingly blurred. Figure 8 displays the X-ray diffraction profiles obtained from the steel tempered across the selected temperature range. The diffraction patterns confirm that tempered martensite constitutes the primary phase in all samples. The persistence of the austenite diffraction peaks, which closely match those in the as-quenched state, indicates the high thermal stability of the RA throughout the tempering process. As tempering proceeds to higher temperatures, partial transformation of austenite to martensite results in a modest decrease in its volume fraction relative to the quenched condition, although the carbon concentration within the RA shows little variation. The distribution of grain boundaries after tempering at each temperature is illustrated in Figure 9, where a continual decline in the proportion of LAGBs is observed with rising tempering temperature.

Figure 10 presents TEM micrographs of the steel at different tempering temperatures, revealing characteristic microstructural features including dislocations, RA films between lamellae, and elongated carbide precipitates within the tempered martensite matrix. Over the tempering interval of 160 °C to 240 °C, a refined lamellar morphology and a high density of dislocations are maintained, with measured lath widths of 0.15 μm, 0.17 μm, and 0.18 μm for the respective temperatures. However, tempering at 280 °C results in a significant increase in plate width to 0.28 μm, accompanied by a marked reduction in dislocation density.

Figure 11 presents a detailed TEM analysis of carbide precipitation evolution in the experimental steel under different tempering conditions. At 160 °C, needle-shaped ε-carbides (transition carbides) are densely dispersed within martensite laths and along dislocation lines. As tempering progresses to 200 °C, θ-carbides (cementite) begin to form through heterogeneous nucleation at ε-carbide interfaces, maintaining the Pitsch–Petch orientation relationship with the ferritic matrix. This transition continues at 240 °C, where θ-carbide nucleation and growth progressively consume the metastable ε-carbides. Upon reaching 280 °C, the transformation is essentially complete, with nearly full dissolution of ε-carbides accompanied by significant coarsening of θ-carbides.

3.2.3. Fracture Analysis

The tensile fracture morphologies are shown in Figure 12. The fractures exhibit no radial zone, consisting only of fibrous zones and shear lips. The macroscopic tensile fracture morphologies of the specimens tempered at 160 °C and 200 °C display irregular fibrous zones and shear lips with significant undulations, showing distinct ductile fracture characteristics. The fibrous zones are relatively large in size and contain numerous dimples, whose presence is an important indication of improved material toughness. Additionally, microvoids (without inclusions) are observed within the dimples, indicating a microvoid coalescence fracture mode. The microscopic morphology of the shear lips is similar to that of the fibrous zones. For specimens tempered at 240 °C and 280 °C, in addition to microvoid coalescence dimples, some small cleavage facets appear in the fibrous zones of the tensile fractures. The shear lips are also composed of small dimples and small cleavage facets.

Figure 13 displays the macro- and micro-scale features of the impact fracture surfaces. The shear lip widths measured after tempering at 160 °C, 200 °C, 240 °C, and 280 °C are 398 μm, 514 μm, 323 μm, and 341 μm, respectively. Quasi-cleavage characteristics, including river patterns and tear ridges along with secondary cracks, are observed on the fracture surface of the sample tempered at 160 °C. The 200 °C tempered fracture surface displayed the widest shear lip width, characterized by large ductile pits surrounded by smaller ones, reflecting a high plastic strain capacity. The 240 °C tempered condition reveals predominantly brittle fracture characteristics, evidenced by numerous cleavage facets, secondary cracks, and transgranular fracture features. Meanwhile, the 280 °C tempered specimen shows a mixed morphology comprising intergranular fracture regions with partial fibrous structures containing shallow dimples, suggesting constrained plasticity and corresponding to the observed strength reduction.

Figure 14 presents the macro- and micro-scale features of the fracture surfaces. Three separate areas are always present on the fracture surfaces: a shear lip zone, a fibrous zone, and a fatigue pre-crack zone. While the fibrous zone has distinctive ductile fracture characteristics with noticeable surface undulations, including both dimples and cleavage facets that indicate plastic deformation during crack propagation, the pre-crack zone has a smooth, featureless appearance that resembles large cleavage facets. Macroscopic analysis reveals that the 160 °C tempered specimen possesses the narrowest shear lip among all conditions, consistent with its predominant brittle fracture behavior, limited crack propagation resistance, and inferior fracture toughness. Detailed microscopic examination of specific regions—fatigue pre-crack zone, transition zone, and fibrous zone—reveals temperature-dependent fracture mechanisms. The 160 °C tempered condition (Figure 14A) exhibits predominantly cleavage features throughout all zones, with numerous cleavage facets and limited secondary cracking in the pre-crack region. This uniform brittle fracture morphology results in obscured transition boundaries, indicating minimal plastic deformation during crack propagation. In contrast, the 200 °C tempered specimen shows quasi-cleavage characteristics featuring small steps, short river patterns, and isolated secondary cracks. The 240 °C tempered condition demonstrates a transition from fine cleavage facets in the pre-crack zone to numerous microvoid coalescence dimples in the fibrous zone, indicating enhanced plasticity during crack propagation. The 280 °C tempered specimen exhibits the most pronounced ductile features, with large tear ridges and cleavage steps in the pre-crack zone, and a bimodal dimple structure in the fibrous zone comprising large dimples surrounded by smaller ones, reflecting superior plastic strain capacity.

4. Analysis and Discussion

4.1. Influence of Quenching Temperature on Microstructure and Properties

The prior-austenite grain structures in the steel, reconstructed via EBSD after quenching between 850 and 910 °C, are presented in Figure 15. The results demonstrate progressive grain coarsening with increasing quenching temperature, attributed to enhanced grain boundary migration driven by elevated thermal energy. The measured prior-austenite grain sizes are approximately 21.23 μm, 26.63 μm, and 29.87 μm for the steel quenched at 850 °C, 880 °C, and 910 °C, respectively, as determined by quantitative analysis. The martensitic transformation, being diffusionless and shear-dominated, initiates and propagates within the parent austenite matrix [26]. Consequently, the prior-austenite grain dimensions establish the spatial constraints for martensite variant development. The observed grain coarsening at higher quenching temperatures consequently leads to enlarged martensite packet and block dimensions, as the expanded austenite grains provide greater accommodation space for the evolving martensitic substructure.

Based on the comprehensive microstructural characterization, the theoretical YS of the quenched specimens was quantitatively evaluated. The YS of the as-quenched martensitic steel is typically described by the following constitutive relationship:

$${\sigma }_{YS}={\sigma }_s+{\sigma }_{GB}+{\sigma }_d$$

(3)

In the formula, ${\sigma }_s$, ${\sigma }_{GB}$, and ${\sigma }_d$ represent the strengthening contributions from the solid solution, grain boundaries, and dislocations, respectively.

Extensive research confirms that the population of carbon atoms retained in the interstitial solid solution within steel is exceptionally low, with the overwhelming majority (~90%) segregating to dislocation cores and lath boundaries as Cottrell atmospheres [27,28]. Therefore, contributions to strengthening from the interstitial solid solution of carbon are not considered in the subsequent analysis. In the experimental steel, alloying elements including silicon, manganese, molybdenum, nickel, and chromium primarily contribute to substitutional solid-solution strengthening. The corresponding YS increment induced by these substitutional solutes can be calculated using Equation (4):

$${\sigma }_s=77+80x_{Mn}+60x_{Si}+45x_{Ni}+60x_{Cr}+11x_{Mo}$$

(4)

where $x_{Mn}$, $x_{Si}$, $x_{Ni}$, $x_{Cr}$, and $x_{Mo}$ correspond to the mass fractions (wt.%) of manganese, silicon, nickel, chromium, and molybdenum in the matrix, respectively. The cumulative strengthening contribution arising from these substitutional solutes amounts to 392.1 MPa.

The Hall–Petch equation describes the grain-refinement strengthening contribution to yield strength (σ_GB), which is expressed in Equation (5) [29,30]:

$${\sigma }_{GB}={\sigma }_0+k_y{d}^{-1/2}$$

(5)

where ${\sigma }_0$ denotes the lattice friction stress of the body-centered cubic (BCC) Fe matrix, taken as 50 MPa based on the established literature [31]. The Hall–Petch coefficient $k_y$ is set to 300 MPa·μm¹/² [32,33,34], consistent with the reported values for martensitic steels. The effective grain size d corresponds to the martensite block dimensions, which govern the Hall–Petch strengthening response. According to Equation (5), the grain boundary strengthening contributions at different quenching temperatures are calculated as 310.2 MPa, 265.1 MPa, and 234.1 MPa, respectively, demonstrating a clear inverse relationship with increasing prior-austenite grain size.

Composed mainly of lath martensite, the as-quenched material contains a high density of dislocations; thus, the corresponding strengthening contribution can be described by the Bailey–Hirsch relation [35,36]:

$${\sigma }_d=\alpha {M_{m}G}_{m}b{\rho }_m^{1/2}$$

(6)

The dislocation density of the quenched specimens was calculated using the modified Williamson–Hall model [37,38]:

$$∆K={\displaystyle \frac{0.9}{D}}+bMK\sqrt{{\displaystyle \frac{\pi \rho \overline{c}}{2}}}$$

(7)

$$\overline{C}={\overline{C}}_{h00}\left(1-q{H}^2\right)$$

(8)

$$F^2=\left(h^2{k}^2+h^2{l}^2+k^2{l}^2\right)/(h^2+k^2+l^2)$$

(9)

where $\beta$ denotes the FWHM of diffraction peaks, $D$ represents the average grain size, and $\theta$ is the Bragg angle. The X-ray wavelength $\lambda$ for Co Kα radiation is taken as 0.17902 nm. $∆K=\beta \mathrm{c}\mathrm{o}\mathrm{s}\theta /\lambda$, $∆K=2\mathrm{s}\mathrm{i}\mathrm{n}\theta /\lambda$. Additional parameters include: the Burgers vector $b$ = 0.258 nm, constant M = 1.4, average dislocation contrast factor $\overline{C}$ for specific (hkl) reflections, ${\overline{C}}_{h00}$ = 0.285 for (h00) reflections, constant $q$ = 2.50 related to dislocation character, and $F^2$ representing the planar index parameter for (hkl) planes.

The calculated average dislocation density in the as-quenched specimens ranges from (17.68 to 20.91) × 10¹⁴ m⁻². According to Equation (6), the corresponding dislocation strengthening contribution is determined to be 480–500 MPa. Overall, the dislocation strengthening contribution exhibits minimal temperature dependence within the investigated quenching temperature range.

The relative contributions of several strengthening mechanisms to the theoretical YS in the as-quenched condition are quantified in Figure 16. The anticipated YSs for quenching temperatures of 850 °C, 880 °C, and 910 °C are 1202 MPa, 1136 MPa, and 1081 MPa, respectively, based on the current computations, demonstrating agreement within 5% of the experimental results. The findings demonstrate that grain-refinement strengthening, which is closely related to the size of martensite blocks, is the main factor influencing YS.

Experimental results demonstrate that YS and UTS are governed by distinct microstructural mechanisms. The sample quenched at 850 °C exhibits the highest YS, which is primarily attributed to the finest martensite block size that enhances the Hall–Petch strengthening effect. In contrast, the sample quenched at 880 °C achieves the maximum UTS. This condition corresponds to the most refined martensitic laths and the highest dislocation density. During plastic deformation, these two factors act synergistically, leading to the formation of a network of high-density dislocation tangles and LAGBs that effectively impede dislocation glide, thereby resulting in a pronounced work-hardening capacity [39,40,41]. Furthermore, the higher fraction of LAGBs in the 850 °C quenched sample serves as “plastic accommodation zones” at the micro-scale. Refined lath structures can increase the tortuosity of crack propagation paths, promoting crack deflection and branching, which improves impact toughness while maintaining high strength. Refined lath structures create a high density of internal boundaries. During fracture, these boundaries act as obstacles, forcing the propagating crack to follow a more tortuous path involving deflection and branching. This process consumes significantly more energy per unit of crack advance, thereby enhancing impact toughness (crack propagation resistance) while the high strength is maintained by the fine microstructure itself. The 880 °C quenched condition, characterized by a combination of refined martensitic laths, high dislocation density, and a favorable grain boundary distribution, achieves an optimal balance between strength and toughness [42,43,44].

4.2. Influence of Tempering Temperature on Microstructure and Properties

Tempering treatment optimizes the strength–toughness balance in the experimental steel through microstructural regulation involving dislocation reorganization, carbide precipitation, and RA stabilization. Compared to the as-quenched condition, tempered specimens demonstrate significantly enhanced YS, where precipitation strengthening from ε-carbides effectively compensates for softening effects associated with reduced dislocation density and martensite matrix recovery [45]. With the tempering temperature rising from 160 °C to 280 °C, a steady decline occurs in YS, UTS, and hardness; conversely, the material’s fracture toughness is progressively enhanced. Notably, both impact energy and A peak following 200 °C tempering, indicating optimal strength–toughness synergy at this condition.

Figure 17 schematically illustrates the microstructural evolution during tempering: (1) At 160 °C, high-density needle-like ε-carbides precipitate within martensite laths, providing substantial precipitation strengthening through dislocation pinning. However, this strong pinning effect restricts dislocation mobility and rearrangement, resulting in quasi-cleavage fracture characteristics and a compromised ductility–toughness combination. (2) At 200 °C, the ε→θ carbide transition moderates dislocation pinning while facilitating coordinated plastic deformation. This condition maintains relatively high dislocation density and stable film-like RA, enabling an optimal toughness–ductility combination with minimal strength sacrifice [46]. The corresponding impact fractures exhibit maximum shear lip width (514 μm) and characteristic microvoid coalescence dimples [47]. (3) At elevated temperatures (240–280 °C), θ-carbide coarsening with loss of matrix coherence reduces precipitation strengthening effectiveness. Concurrent martensite lath widening, significant dislocation density reduction, and partial RA decomposition collectively diminish crack propagation resistance. These microstructural changes manifest as increased cleavage facets on impact fracture surfaces and a progressive deterioration of toughness and ductility [48].

5. Conclusions

This study systematically designed heat-treatment parameters with varying quenching and tempering temperatures for a low-alloy UHSMS. The individual and combined effects of these thermal processing parameters on microstructural evolution and mechanical properties were systematically investigated, followed by comprehensive characterization of the quenched and tempered states. The main findings are summarized below:

(1)

The new low-alloy steel achieves the optimal strength–toughness matching when quenched at 880 °C for 60 min followed by tempering at 200 °C for 120 min: the UTS reaches 2017 MPa, the A is 10.4%, and the impact toughness attains 80.26 J/cm². The microstructure in this state consists of fine tempered martensite, ε-carbides, and θ-carbides, realizing a good balance between strength and plasticity, making this steel a promising candidate for critical aerospace structural components such as landing gear, rocket motor casings, and high-stress fasteners, where both high strength and damage tolerance are essential.

(2)

When the quenching temperature is 880 °C, the synergistic effect between the refinement of martensitic laths and the peak dislocation density at this temperature enables the new low-alloy steel to possess high strength (with an UTS of 2308.8 MPa) while maintaining good plasticity (with an A of 6.2%). Solid-solution strengthening, grain boundary strengthening, and dislocation strengthening are the main reasons for the high strength of the as-quenched specimens. The UTS is dominated by the work-hardening capacity, while the YS is controlled by the Hall–Petch effect.

(3)

Specimens tempered at 200 °C exhibit higher YSs and improved plasticity compared to as-quenched specimens. This excellent combination of properties stems from the nano-scale ε→θ carbide transformation, along with the retention of fine martensitic laths, high dislocation density, and RA films, which significantly enhance toughness and plasticity.