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Effect of quenching-partitioning treatment on the microstructure, mechanical and abrasive properties of high carbon steel

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  • Corresponding author:

    Jian-ping Lai    E-mail: ljp@swust.edu.cn

  • Received: 8 June 2020Revised: 3 August 2020Accepted: 5 August 2020Available online: 10 August 2020
  • The present work employed the X-ray diffraction, scanning electron microscopy, electron backscattered diffraction, and electron probe microanalysis techniques to identify the microstructural evolution and mechanical and abrasive behavior of high carbon steel during quenching-partitioning treatment with an aim to enhance the toughness and wear resistance of high carbon steel. Results showed that, with the increase in partitioning temperature from 250 to 400°C, the amount of retained austenite (RA) decreased resulting from the carbide precipitation effect after longer partitioning times. Moreover, the stability of RA generally increased because of the enhanced degree of carbon enrichment in RA. Given the factors affecting the toughness of high carbon steel, the stability of RA associated with size, carbon content, and morphology plays a significant role in determining the toughness of high carbon steel. The analysis of the wear resistance of samples with different mechanical properties shows that hardness is the primary factor affecting the wear resistance of high carbon steel, and the toughness is the secondary one.
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Effect of quenching-partitioning treatment on the microstructure, mechanical and abrasive properties of high carbon steel

  • Corresponding author:

    Jian-ping Lai    E-mail: ljp@swust.edu.cn

  • 1. Key Laboratory of Testing Technology for Manufacturing Process in Ministry of Education, Mianyang 621010, China
  • 2. School of Manufacturing Science and Engineering, Southwest University of Science and Technology, Mianyang 621010, China
  • 3. Central Sichuan Oil-Gas District, Petro-China Southwest Oil and Gas Field Company, Suining 629000, China

Abstract: The present work employed the X-ray diffraction, scanning electron microscopy, electron backscattered diffraction, and electron probe microanalysis techniques to identify the microstructural evolution and mechanical and abrasive behavior of high carbon steel during quenching-partitioning treatment with an aim to enhance the toughness and wear resistance of high carbon steel. Results showed that, with the increase in partitioning temperature from 250 to 400°C, the amount of retained austenite (RA) decreased resulting from the carbide precipitation effect after longer partitioning times. Moreover, the stability of RA generally increased because of the enhanced degree of carbon enrichment in RA. Given the factors affecting the toughness of high carbon steel, the stability of RA associated with size, carbon content, and morphology plays a significant role in determining the toughness of high carbon steel. The analysis of the wear resistance of samples with different mechanical properties shows that hardness is the primary factor affecting the wear resistance of high carbon steel, and the toughness is the secondary one.

    • To date, less attention has been paid to wear-resistant high carbon steel because it is difficult to achieve a good combination of wear resistance and toughness in high carbon steel, which limits its application to conditions requiring considerable toughness [12]. For example, under high-impact conditions, high carbon steel often fails through brittle fracture with no appreciable toughness during abrasion. Thus, the improvement of toughness is of significance to further extend the application of high carbon steel.

      The innovative approach of incorporating a large amount of retained austenite (RA) in a multiphase microstructure through a series of heat treatment processes can help to enhance the toughness of traditional brittle high carbon steels [34]. The presence of soft RA can effectively accommodate strain and inhibit crack propagation at the interface between the surrounding hard phase and the RA via the transformation-induced plasticity (TRIP) effect during deformation, resulting in the significant improvement of toughness [56]. The wear resistance is also enhanced by surface work hardening through the abrasion-induced martensitic transformation of austenite during wear [79]. Given the beneficial role of RA in improving toughness and wear resistance, more attempts should be made to examine the underlying mechanical and abrasive properties of high carbon steel.

      The amount of RA can be remarkably increased by applying quenching–partitioning (Q&P) treatments, which has been validated to improve toughness significantly without compromising the ultimate strength of low carbon and medium carbon steels [1011]. The key point of Q&P treatment is that the untransformed austenite under the quenched condition is enriched with carbon by partitioning carbon from supersaturated quenched martensite to nearby austenite during the partitioning process [12]. The stability of RA considerably increases with the increase in carbon content and subsequently remains unchanged after the final cooling process. However, most studies of Q&P treatment have only focused on low carbon and medium carbon steels [1317]. Thus, it is reasonable to deduce that the novel Q&P treatment, which has been successfully applied to low carbon and medium carbon steels, can be equally applicable to the case of high carbon steel to induce changes in its microstructure and properties.

      The scientific basis for introducing a large amount of RA to high carbon steel by Q&P treatment can be established. In comparison with low carbon steel, high carbon steel is presumed to have a significantly stronger driving force for carbon partitioning because the carbon resource that is available to participate in the diffusion process can be considerably increased by its higher carbon content. The starting temperature of martensitic transformation (Ms) is closely related to the carbon content of steel, and the high carbon content of high carbon steel corresponds to a low Ms temperature [18]. For the Q&P treatment of high carbon steel, the quenching temperature of high carbon steel needs to be low, i.e., ranging between Ms and the finishing temperature of martensitic transformation (Mf) [14]. With the resulting low partitioning temperature, carbide precipitation is suppressed during partitioning. Given these two factors, i.e., increased carbon resource and suppressed carbide precipitation, more carbon is partitioned into the RA, resulting in the increase in RA and the enhancement of the properties of high carbon steel.

      In this work, the microstructural evolution and mechanical properties of high carbon steel subjected to Q&P treatment are investigated. Two partitioning temperatures (i.e., 250 and 400°C) are employed to examine the competition between RA formation and carbide precipitation. The analysis of the two partitioning temperatures with different isothermal holding times shows that the mechanical property and wear resistance of high carbon steel can be adjusted as needed by controlling the fraction, stability, and morphology of RA in the microstructure.

    2.   Experimental
    • The high carbon steel investigated in this work was composed of 1.2C–1.48Si–0.8Mn–1.0Cr–0.2Mo (in wt%). The addition of up to approximately 1.5wt% Si suppressed carbide precipitation and facilitated carbon partitioning [19]. The high carbon steel was subjected to medium-frequency induction melting in a furnace and cast into ingots after deoxidation. The chemical composition was analyzed using the inductively coupled plasma optical emission spectrometer. Before Q&P treatment, the ingots were homogenized at 1100°C for 2 d in an argon atmosphere furnace, followed by furnace cooling to room temperature. The samples for Q&P treatments were cut from the ingots after removing the decarburized layer of the ingots resulting from homogenization heat treatment.

      For Q&P treatment, the quenching temperature was set to be between Ms and Mf [14], which were determined to be 35°C and −130°C, respectively, using the JMatPro software. The schematic diagram of the Q&P treatment process is shown in Fig. 1. First, the samples were austenized at 1000°C for 30 min in an air resistance furnace. Then, the samples were quenched in an oil bath at 20°C. Subsequently, the samples were partitioned in a salt bath at 250 and 400°C for different holding times. The selected holding times were 20, 90, 300, 600, 1800, and 5400 s for the two partitioning temperatures (i.e., 250 and 400°C). Finally, the samples were quenched in water.

      Figure 1.  Schematic of the quenching–partitioning treatment applied in this work.

    • Microstructural characterization of the samples was performed using X-ray diffraction (XRD, Rigaku D/max 2500, Japan), scanning electron microscopy (SEM, FEI, Sirion 200, USA), electron probe microanalyzer (EPMA, JXA-8230, Japan), and electron backscattered diffraction (EBSD, ZEISS EVO MA10, Germany). The preparation of the samples for XRD measurements involved conventional metallographic polishing steps and etching procedures in a 2vol% nital solution to relieve surface stress by removing the deformed layer [4]. XRD analysis was performed with a scanning speed of 2°/min from 35° to 85°. The samples for SEM and EPMA were polished according to standard metallographic procedures and etched with 4vol% nital solution. For EBSD characterization, sample disks with a diameter of 3 mm were mechanically punched from thin foils with a thickness of approximately 80 µm and electropolished using a twin-jet electrochemical polisher at 30 V in 10vol% perchloric acid and 90vol% glacial acetic acid at −20°C. EBSD analysis was performed with a step size of 100 nm, and the raw data obtained were post-processed using the HKL Technology Channel 5 software provided by Oxford Instruments.

      The volume fraction of RA was further calculated from the XRD patterns. Errors in determining the volume fraction of RA were estimated by performing three measurements for each sample. From the comparison of the results obtained using different methods to calculate the volume fraction of RA with the experimental results of SEM and EBSD, the following equation was employed to determine the volume fraction of RA (γ), which was assumed to have a similar value to that of the experimental observation [14,20]:

      where Vγ is the volume fraction of RA, Iγ is the integrated intensity of the (111)γ and (200)γ austenite (γ) peaks, and Iα is the integrated intensity of the (110)α, (200)α, and (211)α martensite (α) peaks.

      The carbon content in RA is estimated using the following equation [2122]:

      where xγ is the carbon content in RA (wt%) and αγ is the lattice parameter of RA, which is derived from the (200)γ austenite peak using the following equation:

      where λ is the wavelength of radiation, (hkl) are the Miller indices of a plane, and θ is the Bragg angle.

    • The microhardness experiment was performed on the samples using a Vickers hardness tester (BUEHLER 5104) under the test condition of 2 N load for 15 s. The data of at least 10 measurements were averaged. Impact toughness test was performed on the samples (unnotched, a standard dimension of 10 mm × 10 mm × 55 mm) using an impact testing machine (JB-300B). The data represented the average value of five samples. The abrasion test was performed using a dry reciprocating sliding machine (Bruker, UMT-3, USA). The sliding test was performed on a ball-on-flat configuration with upper counter body of ϕ9.5 mm zirconia and bottom body of the test sample. The following test conditions were applied: normal load of 10 N, single-track distance of 5 mm, sliding frequency of 2 Hz, and testing time of 30 min. The wear scars of the samples subjected to the reciprocating sliding test were characterized using a 3D microscope (Keyence, VHX-5000, Japan). The wear resistance is calculated using the reciprocal value of the wear rate, and the relative wear resistance is determined by unitizing the corresponding values of wear resistance.

    3.   Results
    • The microstructural evolution of the samples subjected to the two partitioning temperatures (i.e., 250 and 400°C) for different times was characterized by the XRD, SEM, EPMA, and EBSD techniques. Fig. 2 shows the XRD patterns and the calculated volume fraction and carbon content of RA as a function of partitioning time. As shown in Fig. 2(c), both curves exhibited a similar trend that the volume fraction of RA increased in the initial partitioning stage, followed by a progressive decrease. Specifically, compared with the sample partitioned at 250°C, a lower peak volume fraction of RA and a shorter time at which the peak value was ultimately reached were observed for the sample partitioned at 400°C. Compared with the sample partitioned at 250°C, the precipitation of cementite (M3C) was more pronounced for the sample partitioned at 400°C, which was manifested by the stronger peaks of cementite in the XRD patterns and the shorter time at which the emergence of cementite peaks was observed. Compared with the sample partitioned at 250°C, the carbon content of RA was slightly higher for the sample partitioned at 400°C (Fig. 2(d)), which might be associated with sufficient carbon enrichment in RA facilitated by the higher partitioning temperature.

      Figure 2.  X-ray diffraction patterns of the samples treated at (a) 250°C and (b) 400°C for different times; corresponding (c) volume fraction and (d) carbon content of RA plotted against partitioning time.

      Fig. 3 shows the SEM micrographs of the quenched sample and samples partitioned at 250°C for different times. Compared with the quenched sample (Fig. 3(a)), even though the partitioning time was prolonged only for 20 s at 250°C, considerable changes in the microstructure of the samples were distinguished, with the multiphase microstructure of the quenched state decomposing into tempered martensite with parallel arrays within individual martensite colonies (Fig. 3(b)). When the partitioning time was prolonged to 90 s, a small amount of austenite with an average size of (1.5 ± 0.5) µm was observed in the microstructure. With further increase in the partitioning time to 600 s, the austenite grains became larger, and the average size of the austenite grains increased to approximately be (3.5 ± 1) µm.

      Figure 3.  Microstructures of (a) the quenched sample and samples partitioned at 250°C for (b) 20 s, (c) 90 s, and (d) 600 s.

      The microstructures of the samples partitioned at 400°C are further characterized to analyze the effect of partitioning temperature on the microstructure of high carbon steel. Fig. 4 shows the microstructural evolution of the samples treated at 400°C. The microstructure of the sample partitioned for 20 s was similar to the sample partitioned at 250°C for 90 s, indicating that the formation of RA controlled by carbon diffusion was accelerated with the increase in partitioning temperature. From the comparison of the samples with the maximum volume fraction of RA achieved at 250 and 400°C (Figs. 3(d) and 4(b)), finer blocky austenite grains accompanied by a larger amount of film-like austenite were observed in the sample treated at 400°C than those in the sample treated at 250°C. With the partitioning time further increased to 1800 s, several coarse fresh martensites with feather-like morphology were distinguished as a result of the decomposition of block austenite grains.

      Figure 4.  Microstructures of the samples partitioned at 400°C for different times: (a) 20 s; (b) 300 s; (c) 1800 s; (d) 5400 s.

      The microstructures of the samples with the peak volume fraction of RA for both partitioning temperatures are selected for EPMA to obtain more details on the carbon distribution of individual phase in the microstructure. As shown in Fig. 5(a), two types of martensite were distinguished in the microstructure because of different etching contrasts. The fresh martensite (pointed by the yellow arrows in Fig. 5(a)) in the vicinity of blocky austenite grains (pointed by the green arrows in Fig. 5(a)) exhibited a lower carbon content than the quenched martensite (pointed by the red arrows in Fig. 5(a)). Further analysis showed that carbon was constrained in martensite, which had a higher carbon content than RA. Compared with the sample treated at 400°C, the carbon content in the martensite regions was higher in the sample treated at 250°C, which was evidenced by the redder martensite regions (pointed by the red arrows in Fig. 5(b)).

      Figure 5.  Microstructures of the samples and the corresponding elemental maps of carbon: (a, b) sample treated at 250°C for 600 s; (c, d) sample treated at 400°C for 300 s (C, Ave, and Conc in color scale denote carbon, average, and concentration, respectively).

      EBSD characterization was performed to investigate the phase evolution during partitioning, as shown in Fig. 6. The volume fraction of RA in the sample treated at 250°C showed a gradual increase from 1.5% to 26% with the increase in partitioning time from 20 to 600 s. In the initial partitioning stage (20 s), the nucleation of RA was distinguished at the intersections of martensite colonies with large angles of orientation. The corresponding inverse pole figures showed that the small austenite grains at the boundaries had the same crystallographic orientation, indicating that the nucleation of austenite occurred at the same prior austenite grain boundaries. As the partitioning process continued, the size of RA became larger with the major axis growing along the boundaries and the minor axis developing toward the interior of martensite. This finding can be explained by the higher diffusion coefficient at the boundaries, which results in the migration rate of austenite–martensite interface along boundaries being faster than that of other regions.

      Figure 6.  (a)–(c) Color-coded phase maps and (d)–(f) corresponding inverse pole figures of samples partitioned at 250°C for different times: (a, d) 20 s; (b, e) 300 s; (c, f) 600 s (bcc—Body-centered cubic; fcc—Face-centered cubic; ZO—Z orientation).

    • Fig. 7 shows the impact toughness and hardness of the samples as a function of partitioning time. As shown in Fig. 7(a), the impact toughness of the samples treated at 400°C was almost higher than those of samples treated at 250°C. The peak impact toughness were determined to be 22 J·cm−2 for the sample partitioned at 250°C for 600 s and 26 J·cm−2 for the sample partitioned at 400°C for 300 s. In the hardness versus partitioning time curves, a roughly progressive decrease was observed with the increase in partitioning time for both partitioning temperatures, which was associated with the recovery of martensite during partitioning [17]. The difference was that the decrease was more pronounced for the samples treated at 400°C than the samples treated at 250°C. Hence, the samples partitioned at 250 and 400°C that reached the peak values of impact toughness were designated as PT-250 and PT-400, respectively.

      Figure 7.  (a) Impact toughness and (b) hardness versus partitioning time of the samples.

      The SEM images of the fracture features of PT-250 and PT-400 samples after impact toughness tests are shown in Fig. 8. A mixed fracture mode of ductile and quasi-cleavage features was observed for both samples. For the PT-400 sample, the ductile fracture was more prominent, which was characterized by varying sizes of dimples and lacerated ridges on the fracture surface. For the PT-250 sample, the quasi-cleavage mode was more prevalent, which was characterized by smooth cleavage facets throughout the fracture surface (pointed by the white arrows in Fig. 8(a)).

      Figure 8.  SEM images of fracture features for samples after impact toughness tests: (a) PT-250; (b) PT-400.

    • Because high carbon steels are utilized as a wear-resistant material in most cases, the wear resistance of the samples needs to be evaluated. The microstructure of completely tempered martensite, achieved through a tempering process by annealing at 250°C for 2 h, was added in wear tests to increase the comparability with PT-250 and PT-400 samples. The tempered sample was designated as QT-250 sample. Tempered martensite dominated in the microstructure of QT-250 sample and only a small amount of RA was distinguished due to long tempering times. The hardness and wear rates were measured to be HV (610 ± 10) and 1 × 10−5 mm3·N−1·m−1 for PT-250, HV (525 ± 10) and 1.9 × 10−5 mm3·N−1·m−1 for PT-400, HV (530 ± 10) and 2.6 × 10−5 mm3·N−1·m−1 for QT-250, respectively. Fig. 9 shows the wear data of the samples after sliding tests. From the analysis of wear features of the samples, the degree of wear severity was ranked as QT-250 > PT-400 > PT-250, which was manifested by the increase in volume loss in wear scars.

      Figure 9.  Wear features of the (a) PT-250, (b) PT-400, and (c) QT-250 samples; (d) hardness and relative wear resistance of the samples.

    4.   Discussion
    • According to the Gibbs free energy diagram of martensite (α) and austenite (γ) phases at a fixed partitioning temperature [23], the Gibbs energy gap between two phases (∆Gα→γ) is the driving force for carbon partitioning from martensite into austenite during Q&P treatment, as shown in Fig. 10(a). In contrast to that in the low carbon and medium carbon steels, the martensite in high carbon steel has a higher carbon content, thereby shifting the carbon content rightward and leading to a larger ∆Gα→γ in the diagram. Consequently, the driving force for carbon partitioning from martensite into austenite increases with the increase in carbon content of steel. The paraequilibrium carbon content of two phases shown in Fig. 10(a) and the schematic diagram of carbon distribution at α/γ interface shown in Fig. 10(b) indicate the existence of a carbon content gradient at the interface, illustrating the carbon partitioning direction from martensite into austenite. Figs. 10(c)10(f) show the microstructures and corresponding line scanning of carbon on the retained austenite of the as-quenched and PT-250 samples, respectively. The carbon distribution of the PT-250 sample has a parabolic-shaped gradient in blocky austenite with the carbon content increasing progressively with line scanning from the center to the edge (Fig. 10(f)). This finding is in sharp contrast to that of austenite in the quenched state, in which the carbon content in austenite grains remains nearly constant with only slight fluctuations (Fig. 10(d)). With further partitioning at 250°C for 600 s, the carbon content of supersaturated martensite decreases from (2.3 ± 0.1)wt% to (1.5 ± 0.1)wt% and the carbon content of RA increases from 0.65wt% to 1.1wt% simultaneously; thereby corroborating the intrinsic mechanism of carbon diffusion and kinetics at the martensite–austenite interface during partitioning [2425].

      Figure 10.  (a) Gibbs free energy diagram showing the effect of carbon content on ∆Gα→γ; (b) carbon distribution diagram illustrating the carbon partitioning mechanism (∆Gα→γ—Gibbs energy gap between α and γ phases; Cα and Cγ—Carbon contents of α and γ phases away from α/γ interface; Cα/γ and Cγ/α—Paraequilibrium carbon contents of α and γ phases at α/γ interface); microstructures and corresponding line scanning of carbon on the retained austenite phases of the (c, d) as-quenched and (e, f) PT-250 samples, respectively.

      On the basis of the carbon partitioning mechanism, the evolution of RA is analyzed as a function of partitioning time. In the initial partitioning stage, because of the large carbon content gradient at the martensite–austenite interface, carbon is transported rapidly from the supersaturated martensite to adjacent unstable austenite, increasing the carbon content of RA and thus the stability of RA remarkably. As the partitioning process continues, the carbon content of martensite diminishes gradually, and the driving force for carbon partitioning decreases progressively. When the carbon that flows into RA through the partitioning process is less than that exhausted by the competing reactions, such as carbide precipitation, the carbon content of RA decreases inevitably. As the carbon content of RA decreases below a critical point, the decomposition of RA occurs, leading to the decrease in the fraction of RA.

      To summarize the results obtained and the arguments previously reported, the formation of RA during Q&P treatment is closely associated with carbon diffusion and the amount and stability of RA depend on the degree of how much carbon is involved in the diffusion from supersaturated martensite into RA. Because of the stronger driving force for carbon partitioning in high carbon steel, the maximum fraction of RA is higher than that obtained in low carbon and medium carbon steels [13,26]. With the increase in partitioning temperature from 250 to 400°C, the carbide precipitation effect that exhausts the carbon source becomes more pronounced, and the residual amount of carbon available to participate in the formation of RA decreases, thereby leading to the decrease in the fraction and size of RA.

      Another phenomenon is worth noting. The decomposition of quenched untransformed austenite into martensite occurs in the initial partitioning stage, which is not observed in low carbon steel [13]. Because the quenching process occurs from the fully austenitic temperature region, the microstructure is in a high energy state and relatively metastable. The carbon content in quenched austenite is determined to be approximately 0.65wt%, as shown in Fig. 10(d). By substituting the composition of quenched austenite into the following equation [26], Ms (°C) is estimated to be 229°C.

      where WC, WMn, WSi, and WAl are the contents (wt%) of C, Mn, Si, and Al, respectively.

      Upon heating to a partitioning temperature above Ms, the decomposition of quenched austenite is expected to occur. In addition to the carbon content, another factor should be taken into consideration in the decomposition of quenched austenite. Previous studies reported that the stability of austenite was affected by the constraining effect from phases surrounding austenite, i.e., an increase in stress concentration surrounding metastable austenite decreases its stability [2729]. The level of stress concentration at the interface between quenched martensite and austenite is assumed to be high in high carbon steel because the volume dilatation effect resulting from austenite- martensitic transformation during quenching is prominent in high carbon steel due to its high carbon content. Upon heating to the partitioning temperature, the stress concentration tends to be relieved instantaneously through the decomposition of RA into martensite.

    • In the microstructural design of structural steel, the characteristics of RA (i.e., morphology, stability, and volume fraction) are of significance in determining the toughness of high carbon steel. However, determining which characteristic, i.e., volume fraction or stability, has a dominant role in affecting the toughness of high carbon steel is difficult. In this work, different characteristics of RA, obtained through different processing parameters, are analyzed to determine which characteristic of RA affects the toughness of high carbon steel. The results shown in Figs. 2 and 7 illustrate that the partitioning time at which the maximum toughness is achieved is identical to the time at which the peak volume fraction of RA is obtained for both partitioning temperatures, indicating the important role of RA fraction in determining the toughness of high carbon steel.

      Another interesting feature is noted by comparing the toughness and volume fraction of RA between PT-250 and PT-400 samples. Despite the higher fraction of RA transformed during the impact toughness test, a lower toughness is observed for the PT-250 sample than that of the PT-400 sample, which seems to show a paradox based on the positive correlation between the fraction of RA and toughness (Table 1). Previous studies reported that, apart from the fraction of RA, the toughness of high carbon steel was closely correlated with the mechanical stability of RA, which was affected by the carbon content of austenite [3031], the size of austenite grains [32], the morphology [33], and the constraining effect [2728]. Generally, the film-like austenite is more preferable than blocky austenite in microstructural design of high carbon steel because of its higher stability [33]. When subjected to the impact toughness test, the large blocky austenite grains in PT-250 sample trigger strain-induced martensitic transformation at a relatively small strain level because of its low stability and transform into coarse plate martensite [26,3334], which is prone to stress concentration at the interface and promotes crack propagation during deformation [35]. In contrast, finer blocky austenite grains accompanied with a larger amount of film-like austenite in PT-400 sample help in achieving a higher toughness by accommodating a larger strain. Moreover, the surrounding tempered martensites with a lower hardness in PT-400 sample can accommodate the strain more effectively than the harder martensite in PT-250 sample and lead to a lower stress concentration at the austenite–martensite interface, which, in turn, delays the strain-induced martensitic transformation during impact toughness test. Generalizing from the previous analysis, compared with the volume fraction of RA, the stability of RA in the microstructure plays a more significant role in determining impact toughness of high carbon steel.

      SampleHeat treatmentAfter the impact toughness testAfter the abrasion test
      PT-25032 ± 1.516 ± 10
      PT-40023 ± 112 ± 10

      Table 1.  Changes in the volume fractions of retained austenite under different conditions vol%

      Further analysis of Fig. 7(a) revealed a second increase in impact toughness at longer partitioning times (i.e., 1800–5400 s) for both temperatures. Given the small amount of RA at this stage (Fig. 2), the contribution of TRIP effect of RA to impact toughness of high carbon steel is less significant, and the residual tempered martensites in the microstructure begin to play an important role. It was reported that the tempered martensite with a low hardness was ductile, whereas tempered martensite with a high hardness was brittle [34]. This finding is supported by the fact that the sample partitioned at 400°C for 5400 s with a hardness of HV 530 has a higher impact toughness than the sample partitioned at 250°C with a hardness of HV 590. Given the progressive decrease in hardness at this stage, the ductility of tempered martensites is expected to increase gradually, leading to the second increase in impact toughness.

      The fracture feature was characterized to some degree by its parent microstructure, in which the transformation of austenite grains to martensite occurred during impact toughness tests [36]. The microstructure of PT-250 sample, which is dominated by large blocky austenite grains, is expected to exhibit a low resistance to crack propagation and leads to martensitic transformation at low strains, thus leading to the resulting quasi-cleavage fracture. Meanwhile, the small austenite grains show a high capability to accommodate strains and contribute to a high degree of ductile fracture features [37], in which the fine dimples may be associated with the coalescence of voids resulting from secondary carbides in tempered martensitic matrix [14,38].

    • The wear resistance of a material is governed by hardness, and a higher hardness usually corresponds to a better wear resistance. This rule holds in this work, in which the PT-250 sample with the highest hardness of HV 610 exhibits the best wear resistance among three samples, as evidenced by the role of hardness in enhancing wear resistance. Moreover, the contribution of RA to the wear resistance is identified by comparing the wear resistance and hardness of the QT-250 and PT-400 samples. Despite the similar hardness, the PT-400 samples exhibits a higher wear resistance than QT-250 sample, suggesting the role of microstructural characteristic in affecting wear performance. It was reported that abrasion-induced martensitic transformation of RA contributed significantly to wear resistance during abrasion [9,3940]. The comparison of the microstructural difference between QT-250 and PT-400 samples shows that the presence of a considerable amount of RA in PT-400 sample is responsible for its higher wear resistance.

      An increase in hardness could enhance the resistance of steel to penetration by abrasives, whereas an increase in toughness could hinder the nucleation and propagation of microcracks during wear, thus increasing the wear resistance of materials [39,4142]. Both hardness and toughness contribute to the wear resistance of high carbon steel; however, it is unclear which factor has a dominant role. Despite its lower toughness, the PT-250 sample with a higher hardness showed a higher wear resistance than PT-400 sample, indicating the more significant role of hardness in determining wear resistance of high carbon steel. In conclusion, under the low-impact abrasive conditions in this work, hardness is the primary factor affecting the wear resistance of high carbon steel, followed by toughness.

      Because of the beneficial role of RA in affecting the mechanical and wear properties of high carbon steel, these properties can be adjusted as needed to best fit the application conditions by purposely modifying the characteristics of RA in multiphase microstructure by controlling the processing parameters.

    5.   Conclusions
    • In this work, the microstructure, mechanical, and abrasive properties of high carbon steel were investigated. The influence of RA on impact toughness and wear resistance of high carbon steel was analyzed in terms of stability, morphology, and volume fraction. The detailed conclusions are listed as follows:

      (1) The driving force for carbon partitioning from martensite into austenite was enhanced by increasing the carbon content of high carbon steel, leading to the increase in the volume fraction of RA in the microstructure. With the increase in the partitioning temperature from 250 to 400°C, the volume fraction and average size of RA after longer partitioning times decreased because of carbide precipitation that exhausted the carbon sources.

      (2) The comparison of the microstructure with properties showed that the stability of RA had a larger contribution to toughness. Meanwhile, hardness is the primary factor affecting the wear resistance of high carbon steel, followed by toughness.

    Acknowledgements
    • This work was supported by the Natural Science Foundation of Southwest University of Science and Technology (No. 19zx7163) and the National Natural Science Foundation of China (No. 51975492).

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