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1.0C-1.5Cr钢双相组织的细化和性能双淬火和激光表面淬火相结合的处理外文文献翻译、中英文翻译、外文翻译.doc
1.0C-1.5Cr钢双相组织的细化和性能双淬火和激光表面淬火相结合的处理外文文献翻译、中英文翻译、外文翻译
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1.0C-1.5Cr钢双相组织的细化和性能双淬火和激光表面淬火相结合的处理外文文献翻译、中英文翻译、外文翻译,1.0,1.5,Cr,钢双相,组织,细化,性能,淬火,激光,表面,相结合,处理,外文,文献,翻译,中英文
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Microstructure refinement and properties of 1.0C-1.5Cr steel in a duplex treatment combining double quenching and laser surface quenchingABSTRACTThe 1.0C-1.5Cr steel was subjected to conventional quenching and laser surface quenching treatment. A process combing double quenching and laser surface quenching was proposed for enhancing surface hardness and obtaining finer microstructure. The cementite dissolution and grain growth behavior in the austenitizing process of single quenching, double quenching, and laser surface quenching were studied. The results indicated that compared with single quenching, mean diameter of undissolved cementite particles (UCP) was much finer in double quenching, and the final prior austenite grain size (PAGS) could be decreased by nearly 40% to about 4.5 m. Both grain and cementite particles near the surface will coarsen after laser surface treatment. Compared with single quenching, the PAGS within hardened layer can be decreased by at least 11% through double quenching, and the mean diameter of UCP at the bottom of hardened layer can be decreased by about 20%. Compared with conventional quenching, surface hardness was enhanced by about 20% through laser surface quenching, contributing to the wear resistance. However, the hard and brittle surface layer tends to be crack source during the impact process, leading to the deterioration of final impact toughness. Under the identical laser parameters, the impact absorbed energy is similar in both single and double quenching, which is about 25% of that before laser surface quenching. The impact absorbed energy can be increased from 22J to 28J by preheating at 160 during the laser surface quenching.Keywords：Double quenching；Laser surface quenching；Cementite；Grain size；Impact toughness ；Wear1. Introduction As a typical hypereutectoid steel, 1.0C-1.5Cr (SAE 52100) steel is extensively used in bearings, guide rails, and molds. In order to satisfy the complex working conditions, it is usual for this steel to be subjected to spheroidization annealing, quenching, and low temperature tempering treatment in sequence, which will finally produce a microstructure composed of martensite matrix, undissolved cementite particles, retained austenite, and nanoscale carbides 1,2. Its final mechanical properties can be controlled by adjusting these microstructure parameters. It is commonly accepted that the microstructure refinement is an effective method for improving the comprehensive mechanical properties of 1.0C-1.5Cr steel, especially its fatigue property. At present, there have been several research works surrounding the microstructure refinement of 1.0C-1.5Cr steel. For example, Beswick 3 reported that the prior austenite grain size (PAGS) could be decreased by imposing pre-cold deformation before quenching. Because lots of low angle dislocation cells form during cold deformation, which will increase the nucleation rate of austenite in the subsequent heating process, and thus lead to the refining of final prior austenite grain (PAG). Santos et al. 4 proposed that induction heating and repeated quenching were also feasible for decreasing the PAGS, and their work demonstrated that fatigue property could be enhanced obviously by refining PAG. Mizobe et al. 5 reported that repeated quenching from 850 could decrease the PAGS from 15 m to about 7 m. Li et al. 6,7 reported that thermo-mechanical control process could be used to refine the hot rolled and spheroidized microstructure, which would result in the slight decrease of size of both PAG and undissolved cementite particles (UCP) in the subsequent quenched microstructure. Recently, Salloom et al. 8 proposed that concurrent refinement of PAG and UCP could be achieved by a double quenching process, which consisted of austenitizing at 1050 followed by oil quenching and low-temperature tempering, and reheating to second austenitizing temperature of 850 followed by quenching and tempering. However, it may cause serious decarburization during austenitizing at 1050 , and microcracks may also form during the subsequent oil quenching 9. Lee et al. 10 reported that finer PAG could be obtained by lowering the second austenitizing tem-perature in the double quenching, whereas lower austenitizing tem-perature failed to enhance the final hardness. Therefore, they proposed that nitrocarburization could be conducted during the first austenitizing stage. However, it is time consuming and costly. In the above methods, repeated quenching, especially double quenching, is extremely prom-ising in refining the microstructure of 1.0C-1.5Cr steel. In order to further optimize microstructure and enhance mechanical properties, it is necessary to carry out the research works surrounding the cementite dissolution and grain growth in the double quenching process.It should be noted that both single quenching and double quenching belong to conventional hardening treatment, and usually result in the through-hardening of 1.0C-1.5Cr steel. However, in many working conditions, it is the surface that will bear complex stress and wear. Therefore, the optimization of surface properties of 1.0C-1.5Cr steel has received particular research attentions. There have been several methods for enhancing surface properties, such as surface thermo-chemical treatments (carburizing, carbonitriding, boronizing), shot peening, physical/chemical vapor deposition, laser surface treatment (quenching, cladding, shock peening) 1113. Due to the low cost, high efficiency, and excellent controllability, laser surface quenching is one of the promising methods for enhancing the surface properties of 1.01.5Cr steel. Basu et al. 14 developed a duplex treatment process for 1.01.5Cr steel, which consisted of austempering and laser surface quenching. It was found that laser surface quenching could enhance the surface hardness without affecting its bainitic core. Lin et al. 15 pro-posed that water cooling could be imposed during the laser surface remelting of 1.0C-1.5Cr steel, which would result in higher surface hardness and the formation of an ultrafine microstructure. Due to the beneficial effects on surface hardness, laser surface quenching could be employed to replace the carbonitriding process used in the double quenching treatment. However, until now, few works combined double quenching and laser surface quenching. In addition, the impact tough-ness of this steel should be also paid attention, especially when it is used in the bearings for high-speed rail or maglev train 16,17. However, the impact properties after laser surface quenching was seldom investigated.In this work, the 1.0C-1.5Cr steel was subjected to conventional single quenching, double quenching, and laser surface quenching. The dissolution and size variation of cementite particles, and grain growth behavior in the austenitizing process of conventional quenching and laser surface quenching were studied. The hardness, wear resistance, and impact toughness of 1.0C-1.5Cr steel were analyzed. In addition, the effect of preheating was also studied.2. Experimental procedure2.1. Materialswhich had a microstructure composed of spherical cementite particles and ferrite matrix. The corresponding chemical composition (wt.%) was: 1.0C, 0.28Si, 0.34Mn, 1.57Cr, 0.011P, 0.003S. The mean diameter of spheroidized cementite particles was about 0.34 m.2.2. Conventional quenching and laser surface quenchingIn order to study the cementite dissolution and grain growth behavior in single and double quenching, specimens with dimensions of 3 10 11 mm were prepared for conducting interrupted quenching, as shown in Fig. 1. For single quenching, some specimens were put into a tube furnace with the temperature of 800 and 850 for austeni-tizing, respectively, and held for 5 min, 20 min, 30 min, 60 min, and then they were oil quenched to room temperature. For double quench-ing, some specimens were first put into the furnace with the temperature of 850 (First austenitizing temperature), and held for 30 min, and oil quenched to room temperature. Subsequently, they were put into the furnace with the temperature of 800 (Second austenitizing temper-ature) again. They were oil quenched to room temperature after holding for 5 min, 20 min, 30 min, 60 min, respectively. It should be noted that the first austenitizing temperature and time in double quenching were fixed in this work, and thus the austenitizing temperature and time mentioned in the following text denote the second austenitizing tem-perature and time.Specimens with dimensions of 11 24 100 mm were also pre-pared. Similarly, they were put into the furnace with the temperature of 850 , and held for 30 min, and then oil quenched to room temperature. Subsequently, they were divided into two groups, and one group was directly tempered at 170 for 120 min, and the other group was put into the furnace with the temperature of 800 again. After holding at 800 for 20 min, they were also oil quenched to room temperature, and tempered at 170 for 120 min. These tempered specimens were further machined to 10 24 100 mm, which would remove the sur-face decarburized layers. Subsequently, they were subjected to laser surface quenching. Fig. 2 shows the schematic illustration of laser sur-face quenching. A fiber-coupled diode laser (Laserline) with spot size of 3 8 mm was used in this work. The scanning speed of laser is set to 20 mm/s, and the power was set to 1500 W and 2000 W, respectively. In addition, during the laser surface quenching, some specimens were preheated at 160 . Table 1 further shows the laser surface quenching parameters.2.3. Microstructure characterizationAll the specimens were cut, ground and polished. Subsequently, these specimens were etched with 4% nital. The microstructure was examined by optical microscope (Olympus BX53 M and Zeiss Axio Scope A1), scanning electron microscope (Zeiss PIGMA HV-01-043), electron probe (JEOL JXA8530F), and transmission electron microscope (FEI Tecnai G2 F20). TEM foils were prepared by electro-polishing using a solution composed of 10% perchloric acid and 90% ethanol. In addition, in order to observe the prior austenite grain boundaries, specimens were also etched at 65 for about 70 s with a solution composed of saturated picric acid and a small amount of Teepol wetting etchant. In this work, linear intercept method was used to measure the prior austenite grain size. In order to quantify the UCP, the cementite particles were colored using photoshop in advance, and then they were measured using Image-pro Plus. The volume fraction of UCP was assumed to be equal to its corresponding area fraction 18. Mean diameter of each UCP was calculated from its area by assuming a spherical shape 19. The number of cementite particles measured in each condition ranged from 600 to 2500, which depended on the austenitizing parameters.2.4. Hardness, wear resistance, and impact toughnessHardness was measured using a microhardness tester (HMV-2TADW) with a load of 0.3 kgf. Wear resistance was evaluated using a pin-on-desk wear test rig. The pin was a Si3N4 ceramic ball with 5 mm in diameter, as shown in Fig. 3a. The wear test was conducted at a rotational speed of 500 rpm, and rotational diameter of 5 mm. A normal load of 1.5 kg was imposed on the specimen during the wear test. The impact toughness was measured using a pendulum impact testing machine (SANS ZBC2452-B). The unnotched impact specimens with dimensions of 10 10 55 mm were used in this work 16. It should be noted that, for the specimens subjected to laser surface quenching, the hardened layer was in the opposite side with pendulum during the impact test, as shown in Fig. 3b.3. Results3.1. Microstructure in single and double quenching3.1.1. Cementite dissolution and sizeFig. 4 shows the microstructure of specimens in the single and double quenching treatment. It can be observed that there exist lots of spherical UCP in the martensite matrix. The volume fraction and size distribution of UCP are of great importance for the final mechanical properties. Usually, the increase of austenitizing temperature and extension of austenitizing time will result in the further dissolution of cementite. It is believed that the dissolution behavior of cementite will influence the size variation of UCP. In this work, the volume fraction and mean diameter of UCP under different austenitizing conditions were measured, as shown in Fig. 5.It can be seen from Fig. 5a, with the extension of austenitizing time from 5 min to 60 min, the volume fraction of UCP of the specimens austenitized at 800 C and 850 in single quenching decreased from 10.4% to 7.3%, and 8.0%2.4%, respectively. However, it decreased from 6.8% to 3.4% for the specimens subjected to second austenitizing at 800 in double quenching. Under the identical austenitizing tem-perature, the volume fraction of UCP is smaller in double quenching compared with single quenching. In addition, considering the volume fraction of cementite in the initial microstructure, it can be found that the volume fraction of UCP decreases continuously in the whole auste-nitizing process of single quenching, whereas it increases first, and then decreases in the later stage of second austenitizing process of double quenching. In other words, some new cementite forms at the early austenitizing stage, which can be observed in Fig. 4c. The size of these newly formed cementite particles is much smaller, which is about 0.1 m.Fig. 5b shows the variation of mean diameter of UCP in the auste-nitizing process. With the increase of austenitizing time, mean diameter of UCP decreases first, and then increases. In early austenitizing process of single quenching, the mean diameter of UCP of the specimen The difference in volume fraction and mean diameter of UCP in single and double quenching should be attributed to the different initial microstructure. For single quenching, the initial microstructure is composed of spherical cementite and ferrite matrix (Fig. 7a). For double quenching, the initial microstructure is mainly composed of UCP and martensite matrix, which is obtained through the previous first quenching. Compared with ferrite matrix, the martensite matrix is metastable, and it contains lots of dislocation and twin boundaries. In addition, carbon is supersaturated in martensite matrix, which induces the serious distortion of lattice. As a result, even at a quite low tem-perature (160 ), nano-carbides can form at the dislocation or twin boundaries within several minutes (Fig. 7b), and all these nano-carbides will transform into -carbides (i.e. cementite) 20,21. In this work, the measured values of volume fraction of cementite in the initial micro-structures of single and double quenching are 19.4% and 4.1%, respectively, which further indicates the initial microstructure of double quenching is supersaturated with carbon.3.1.2. Grain size in single and double quenchingFig. 8 shows the prior austenite grain boundaries of the specimens in single and double quenching. It is recognized that the grain growth is mainly driven by the decrease of interfacial energy. The growth process is a result of grain boundary migration, which is closely related to the diffusion of atoms 22,23. At higher austenitizing temperature and longer time, the PAGS tends to be larger due to the fact that effective diffusion distance of atoms is larger. The PAGS was further measured in this work. In single quenching, with the increase of austenitizing time from 5 min to 60 min, the PAGS of the specimens austenitized at 800 and 850 increases from 7.6 m to 8.1 m, and from 7.8 m to 8.6 m, respectively. There is no obvious difference in the PAGS under these two austenitizing temperatures. In double quenching, the PAGS of the specimens austenitized at 800 increases from 4.1 m to 4.7 m. Compared with single quenching, the PAGS is decreased by almost 40% in double quenching.3.2. Microstructure after laser surface quenching3.2.1. Morphology of the hardened layerAfter single quenching and double quenching, specimens were sub-jected to tempering, and laser surface quenching. Fig. 9 shows the microstructure after laser surface quenching. The hardened layer has a flat bottom except its two sides, which is because the heat dissipation is faster at two sides. Under identical laser parameters, the size of hard-ened layer is almost same, regardless of single and double quenching. When the laser power is 1.5 kW, the width of hardened layer is about 10am.It increases to about 10.6 mm at 2 kW, and preheating further in-creases its width to 11.1 mm. The microstructure of hardened layer is further examined, as shown in Fig. 10. There are few UCP at the surface, which is mainly due to the significant dissolution caused by higher surface temperature. In addition, the subsequent rapid cooling rate prevents the formation of cementite, which further leads to few UCP at the surface. The dissolution of cementite will leave the matrix super-saturated in carbon. Lots of UCP can be observed at the depth of 250 m and 625 m. In other words, a gradient microstructure was obtained after laser surface quenching.3.2.2. Dissolution and size of cementite within hardened layerThe volume fraction and mean diameter of UCP at different depths of hardened layer were measured, as shown in Fig. 11. It is quite obvious that with the increase of depth, the volume fraction of UCP increases, and mean diameter of UCP decreases. The increase of laser power, or preheating will accelerate the cementite dissolution, which is accom-panied with the coarsening of cementite. In addition, although the initial specimens were obtained through single and double quenching, respectively, there is no significant difference in the volume fraction of UCP at the same depth position of these two kinds specimens under identical laser parameters. However, compared with the specimens ob-tained through single quenching, the mean diameter of UCP at the bottom of hardened layer is obviously smaller in the specimens obtained through double quenching, which is decreased by about 20%. In other words, when the mean diameter of UCP is smaller in the initial micro-structure, the mean diameter of UCP is also smaller after laser surface quenching. Before laser surface quenching, the volume fraction of UCP in the specimens of single and double quenching was about 4.1% and 4.2%, respectively, and the corresponding mean diameter of UCP was about 0.24 m and 0.17 m, respectively. After laser surface quenching, at the deeper position, the volume fraction of UCP in the specimens of single and double quenching was at least 5.0% (Fig. 11a), which is slightly larger than that before laser surface quenching. In other words, some new cementite will form at the deeper position during the laser surface quenching. However, the mean diameter of UCP at the deeper position is quite close to that before laser surface quenching. On the whole, finer UCP can be obtained through the process combining double quenching and laser surface quenching.The cementite dissolution behavior is closely related to the temper-ature distribution during the laser surface quenching. However, it is very difficult to measure the temperature distribution along the depth. In this work, the temperature filed was calculated using ABAQUS based on the Fouriers heat conduction equation, as shown in Eq. (1).where , c, and is density, specific heat, and thermal conductivity, respectively. T is temperature, and t is time. Q is the latent heat of phase transformation. The values of , c, and is obtained through the JMatPro (version 7.0). The light intensity of laser heat source is thought to be distributed uniformly in this work. In addition, a 3-D model having size of 10 24 100 mm was used.Fig. 12 shows the calculated results of temperature field. With the increase of depth, the temperature decreases continuously. At the laser power of 2 kW, the surface temperature increases from 100 to the maximum temperature (1479 ) within 0.155 s, and decreases to 100 within 0.555 s, which indicates the average heating and cooling rate is about 8896 /s and 2484 /s, respectively. Both the heating and cooling rate are significantly larger than that in the conventional quenching. The larger heating rate leads to a gradient distribution of temperature along the depth, and the larger cooling rate guarantees the formation of hardened layer. In addition, with the increase of laser power from 1.5 kW to 2 kW, the maximum temperature nearing the surface increases obviously, whereas it increases slightly at deeper po-sition (Fig. 12d). At the laser power of 2 kW, the surface temperature exceeds the liquidus line (1460 ) of 1.0C-1.5Cr steel, which indicates that localized micro melting should occur. Due to the surface micro melting, the maximum temperature at the vicinity of surface is not increased obviously through pre-heating, whereas the maximum tem-perature at deeper position is obviously larger than the specimen without preheating. The gradient distribution of temperature along depth leads to the variation of both volume fraction and mean diameter of UCP.3.2.3. Grain size after laser surface quenchingFig. 13 further shows the PAGS at different depths after laser surface quenching. For the specimens obtained through single quenching (SL1, SL2, and SL3), with the increase of depth, the PAGS decreases first, and then increases to about 8.5 m, which is also the PAGS of initial microstructure before laser surface quenching. The PAGS tends to be smaller at the depth of 500 m. For the specimens obtained through double quenching (DL1, DL2, and DL3), with the increase of depth, the PAGS decreases first to 4.5 m, and then changes slightly when the depth exceeds 500 m. The grain size at different depths should be related to the temperature. According to Figs. 12d and 13a, the PAGS at different depths as a function of the maximum temperature at corre-sponding depth was further given in Fig. 13b. It is observed that when the maximum temperature exceeds 800 , the PAGS begins to increase significantly. Under the identical laser parameters, compared with the specimens obtained through single quenching, the PAGS can be decreased by at least 11% at the depth of 125 m in the specimen ob-tained through previous double quenching. It can be decreased by about 40% at the depth of 500 m. In other words, finer PAGS can be obtained through the duplex process combining double quenching and laser surface quenching.3.3. HardnessFig. 14 further shows the hardness after laser surface quenching. It is quite obvious that the surface hardness was enhanced after laser surface quenching. According to the hardness distribution and microstructure, there are mainly three regions in the specimens after laser surface quenching, which includes hardened layer, heat affected zone, and substrate (Fig. 9). The hardened layers represent the regions which have completed the austenite transformation during the laser surface quenching. The heat affected zone is next to the hardened layer, in which region the austenite transformation is insufficient or no austenite transformation occurs. The substrate region is at the innermost side, which is not affected by the laser. It can be observed that, as the laser power increases from 1.5 kW (SL1 and DL1) to 2.0 kW (SL2 and DL2), the thickness of hardened layer increases from about 0.48 mm to about 0.66 mm, and the preheating (SL3 and DL3) further increases its thickness to about 0.72 mm.For the initial specimens obtained through single and double quenching, the hardness before laser surface quenching is about 795 HV and 750 HV, respectively, which is also the corresponding hardness of substrate regions after laser surface quenching. At the laser power of 1.5 kW and 2.0 kW, the maximum hardness within the hardened layer is about 1000 HV. Compared with the hardness of initial specimens, the hardness is increased by about 20% after laser surface quenching. For the specimens with preheating, the maximum hardness within the hardened layer is about 870 HV, which is smaller than that without preheating, whereas it is still significantly larger than the corresponding hardness before laser surface quenching. It should be noted that in all the specimens, the hardness of surface is lower than that of subsurface region, which is mainly due to the coarsening and decarburization of surface microstructure during laser surface quenching 24. Compared with hardened layer and substrate region, the hardness is lower in the heat affected zone due to the tempering effect. In addition, under the identical laser parameters, the hardness of hardened layer is similar for the specimens obtained through single and double quenching. However, compared with the specimens obtained through single quenching, the hardness of both heat affected zone and substrate region is slightly lower in the specimens obtained through double quenching, which corre-sponds with the hardness before laser surface quenching.3.4. Wear resistanceIt is usual for this steel to undergo wear in the actual working con-dition. In this work, the wear resistance was evaluated by the weight loss, as shown in Fig. 15. Note that the wear test was conducted completely in the laser surface quenched region. For the initial speci-mens without laser surface quenching, the wear loss under identical friction time is larger in double quenching, indicating a poor wear resistance. It is mainly because the hardness is lower for the specimens in double quenching compared with that in single quenching. After laser surface quenching, the wear loss under identical friction time is almost same for the two kinds of specimens. However, it is smaller than that of the specimens without laser surface quenching. In other words, the wear resistance can be enhanced by laser surface quenching, which should be attributed to the higher surface hardness.3.5. Impact toughnessAs a hypereutectoid steel, the impact toughness of 1.0C1.5C steel is relatively lower after quenching. Its impact absorbed energy is around 4 J when V-notched impact specimens was used 10. In order to compare the impact toughness, unnotched impact specimens were used in this work. Table 2 shows the impact absorbed energy of the specimens with and without laser surface quenching. For the initial specimens without laser surface quenching, the impact absorbed energy in both single and double quenching is about 84 J, and under the identical laser parameter, the impact absorbed energy of these two kinds of specimens is also similar. However, compared with the initial specimens, the impact absorbed energy decreased significantly after laser surface quenching, which is about 25% of that before laser surface quenching. At the laser power of 1.5 kW and 2.0 kW, the impact absorbed energy is about 24 J and 22 J, respectively. It can be slightly increased to about 28 J through preheating at 160 , which should be related to the softening of hardened layer and the decrease of residual stress 13.The impact fracture morphology was further observed, as shown in Fig. 16. For the specimens without laser surface quenching (Fig. 16a and b), the fracture surface appears to quasi-cleavage fracture. Lots of UCP can be observed in the fracture surface 25. The appearance of UCP indicates that some microcracks form at the interface between UCP and martensite matrix. For the specimens with laser surface quenching (Fig. 16c and d), microcracks initiate at the surface of hardened layer (Arrow A), and then propagate towards the substrate. At the vicinity of crack origin, the fracture surface appears intergranular fracture, indi-cating the hardened layer is quite brittle. At the middle of hardened layer (Arrow B) and heat affected zone (Arrow C), the fracture surface appears quasi-cleavage fracture, which is similar with the fracture sur-face of specimens without laser surface quenching. In addition, it is observed that the fracture surface is smoother in the hardened layer compared with that in heat affected zone, which further indicates the brittle nature of hardened layer. According to the fracture surface, the fracture can be divided into three stages in the specimens with laser surface quenching. First, microcracks form at the surface of hardened layer when the stress exceeds the strength limit. And then, cracks propagate quickly within the hardened layer. At last, it further propa-gates at a slower rate until the thorough fracture of specimens. The interface between martensite and cementite particles are potential sites for crack formation and propagation.4. Discussion4.1. Cementite dissolution and size variation4.1.1. Cementite in single and double quenchingDue to the different initial microstructure, cementite dissolution behavior is different in the austenitizing process of single and double quenching. Note that in this work, the austenitizing process of double quenching specifically denotes the second austenitizing process (Fig. 1b). For single quenching, the initial microstructure is composed of spherical cementite and ferrite matrix (Fig. 7a). However, for double quenching, the initial microstructure before second austenitizing pro-cess is mainly composed of UCP and martensite matrix, which is ob-tained through the previous first oil quenching. In the austenitizing process of single quenching, ferrite will transform into austenite, and cementite will dissolve continuously due to the temperature rise and significantly different carbon solubility between ferrite and austenite. After ferrite-to-austenite transformation, the cementite continues dis-solving into austenite to obtain a homogeneously distributed austenite. During the heating stage, the dissolution rate of cementite increases with time, whereas it decreases gradually with time during the holding stage 26. In the austenitizing process of double quenching, lots of nano-carbides will form before the completion of austenite trans-formation. It was reported that when this steel was tempered at 160 for 4 h after quenching, the volume fraction of cementite could increase from 4.0% to 14.1% due to the precipitation of nano-carbides 21. After the completion of austenite transformation, these newly formed car-bides start dissolve into austenite. Therefore, the volume fraction of UCP increases first, and then decreases in the austenitizing process of double quenching.The size variation of cementite particles depends on many factors including dissolution behavior, Gibbs-Thomson effect, and precipitation of nano-carbides. In the early austenitizing process of single quenching, the cementite dissolution is dominated, leading to the decrease of mean diameter of UCP. However, in the later stage, the cementite dissolution is slow, and the cementite will coarsen at the expense of smaller ones due to Gibbs-Thomson effect, which is similar with the Ostwald ripening 27,28. In the early austenitizing process of double quenching, espe-cially the heating stage, lots of nano-carbides will form, and finally evolve into tiny cementite particle. Therefore, the size of UCP is smaller in the early austenitizing stage. With the increase of austenitizing time,these cementite particles will also coarsen due to the Gibbs-Thomson effect. In single quenching, the UCP appears to be bimodal distribu tion at the austenitizing time of 5 min, which is mainly related to the size distribution of cementite in the initial spheroidized microstructure. It was reported that the size distribution of cementite in spheroidized microstructure is also bimodal due to the different growth rates of cementite at grain boundaries and within grains during spheroidization annealing 29,30. In the early austenitizing process of single quench ing, accompanying with the cementite dissolution, the mean diameter of both small-sized (00.15 m) and large-sized (0.150.6 m) cementite decreases. As a result, the size distribution of UCP is still bimodal. However, at the later austenitizing stage, some small-sized cementite particles will dissolve completely, and some large-sized cementite par ticles will coarsen instead of dissolution due to the Gibbs-Thomson ef fect. Therefore, the number fraction of small-sized UCP will decrease, and number fraction of large-sized UCP will increase. The boundary between number fraction of small-sized cementite and that of large-sized cementite becomes ambiguous, which results in the less obvious bimodal distribution. In double quenching, lots of tiny cementite particles will form at the early stage of second austenitizing process. As a result, the total number of cementite per unit area will be quite large, and the number fraction of small-sized cementite is far larger than that of large-sized cementite. The peak caused by large-sized cementite is quite small, and thus the bimodal character is not obvious in double quenching.4.1.2. Cementite in laser hardened layerThe initial microstructure before laser surface quenching is mainly composed of cementite and low temperature tempered martensite ma trix. During the laser heating, the martensite will transform into austenite, and cementite will dissolve into austenite. Simulated results of temperature field indicate the surface temperature is obviously larger than that at deeper position. At the laser power of 2 kW, the surface temperature even exceeds the liquidus line. Therefore, significant dissolution of cementite occurs at the vicinity of surface, and with the increase of depth, the dissolution degree of cementite decreases, i.e. the volume fraction of UCP increases. It should be noted that accompanied with the significant dissolution of cementite, the overall mean diameter of UCP will not decrease. It is mainly because the surface temperature is so high that smaller-sized cementite particles will dissolve into the matrix completely at the vicinity of surface. As a result, the overall mean diameter of UCP is larger at the vicinity of surface. It was also reported that the mean diameter of undissolved carbides was larger at higher austenitizing temperature in the conventional quenching treatment 31. In order to obtain finer UCP, the surface temperature should be not too high. Compared with the surface position, the smaller-sized cementite particles will not dissolve completely at deeper position of hardened layer due to the lower maximum temperature. Before laser surface quenching, there are lots of smaller-sized cementite in the initial microstructure obtained through double quenching. Therefore, when the initial microstructure is obtained through double quenching, the overall mean diameter of UCP is obviously smaller at deeper position of hardened layer.4.2. Grain sizeThe final PAGS depends on the comprehensive effect of nucleation rate and growth kinetics. The PAGS can be decreased by increasing austenite nucleation rate and decreasing growth rate. In the conven-tional heat treatment, the PAGS can be refined significantly through double quenching. It is mainly because the initial microstructure of second austenitizing process of double quenching is metastable martensite. Both the dislocation and twin crystal in the martensite are potential nucleation sites for austenite, which has a beneficial effect on increasing the nucleation rate of austenite, and thus leads to the refinement of PAG.For both initial microstructures, after laser surface quenching, the PAGS at the vicinity of surface is obviously larger than that at deeper position of hardened layer. It is mainly because the maximum temper-ature decreases with the increase of depth. At lower temperature, the diffusivity of atoms is smaller, which is not beneficial to the migration of grain boundary. In addition, there are more UCP at deeper position, which also hinders the migration of grain boundary. Therefore, compared with the surface, the PAGS tends to be smaller at deeper po-sition. However, when the depth extends a certain value, the nucleation of austenite is insufficient, or there is no austenite formation due to lower temperature. As the depth continues increasing, the number of newly formed austenite grains decreases. Because the size of newly formed austenite grains is relatively small. When its number decreases, the final PAGS will increase. It should be noted that when the depth exceeds 500 m, the PAGS changes slightly in the specimens obtained through double quenching. That is mainly because the PAGS of the initial microstructure is too small in double quenching. The decrease of the number of newly formed austenite grains will not influence the final PAGS significantly. After laser surface quenching, the PAGS is much smaller in the specimens obtained through double quenching, which can be attributed to higher austenite nucleation rate. Before laser surface quenching, the PAGS is much smaller in the specimens obtained through double quenching (Fig. 8), and thus the area of corresponding grain boundaries per unit volume is larger. The grain boundaries are potential nucleation site for new austenite grain nucleus. With the increase of area of grain boundary in the initial microstructure, the nucleation rate of austenite will also increase during the laser surface quenching, which contributes to obtaining finer PAG.4.3. Deterioration of impact toughness after laser surface quenchingBoth the surface hardness and wear resistance can be enhanced after laser surface quenching. However, the impact toughness becomes worse. It should be attributed to the brittle nature of laser hardened layer. During the impact process, the hardened layer will bear tensile stress. Once the tensile stress exceeds the fracture limit, microcrack will form at the outer surface of brittle hardened layer. The stress concentration nearing tip of microcrack will further accelerate the propagation of crack, and finally lead to the thorough fracture. In this work, the hard-ened layer was softened by preheating at 160 , which was beneficial to decrease its brittleness, and thus delay the formation of crack. Therefore, the impact toughness can be improved by preheating during the laser surface quenching. However, the preheating will lead to the decrease of hardness of hardened layer due to the fact that cooling rate is decreased. In addition, when the preheating temperature exceeds martensite start transformation, the martensite transformation will also be retarded,weakening the strengthening effect of laser surface treatment. Consid-ering the fact that the 1.0C-1.5Cr steel is usually tempered below 170 after quenching, the preheating temperature should be also lower than 170 , which will not decrease the hardness of matrix significantly.According to the above results, double quenching can be combined with laser surface quenching to obtain finer microstructure. The microstructure refinement will enhance the fatigue properties of 1.0C-1.5Cr steel 4,10. Although compared with the specimens obtained through single quenching, the hardness is smaller for the specimens obtained through double quenching, whereas the hardness within the hardened layer is similar for these two kinds of specimens after laser surface quenching. Furthermore, the wear resistance and impact toughness are also similar for these two kinds of specimens after laser surface quenching. In other words, compared with the duplex treatment composed of single quenching and laser surface quenching, the duplex treatment composed of double quenching and laser surface quenching will not decrease the wear resistance and impact toughness. Therefore, in terms of microstructure refinement, it should be feasible to combine double quenching and laser surface quenching. However, the laser surface quenching will lead to the deterioration of impact toughness. The duplex treatment composed of double quenching and laser surface quenching can only be used for the working conditions of small impact loads. Further work will be carried out to improve the impact toughness of 1.0C-1.5Cr steel in laser surface quenching.5. ConclusionsIn this work, the microstructure evolution of 1.0C-1.5Cr steel in the conventional quenching and laser surface quenching process was stud-ied. Microstructure refinement was achieved through double quenching, and a process combining double quenching and laser surface quenching was proposed. The hardness, wear resistance, and impact toughness were analyzed. The following conclusions are drawn:(1)In the austenitizing process of single quenching, with the increase of austenitizing time, cementite dissolves continuously, whereas in the austenitizing process of double quenching, new small-sized cementite particles will form during the heating and early hold-ing stage, and then cementite begins to dissolve into austenite. In the austenitizing process of both single and double quenching, the mean diameter of UCP decrease first, and then increases, which are due to the comprehensive effect of dissolution, Gibbs-Thomson effect, and formation of new cementite particles.(2)Compared with single quenching, the minimum value of mean diameter of UCP can be decreased by about 40% to 0.14 m through double quenching, and PAGS can also be decreased by about 40% to 4.5 m. The size distribution of UCP appears to be bimodal in single quenching. At the austenitizing temperature of 800 and austenitizing time of 5 min, the number fraction of UCP with the diameter smaller than 0.15 m is 26.4% in single quenching, and it increases to 72.4% in double quenching.(3)In laser surface quenching, both cementite and PAG coarsened obviously at the surface due to the higher surface temperature. At a deeper position of hardened layer, the mean diameter of UCP is smaller. Double quenching can be combined with laser surface quenching to obtain finer microstructure.(4)Compared with conventional quenching, the surface hardness can be increased by about 20% through laser surface quenching, which is beneficial to enhance the wear resistance. However, microcracks are easy to form within the surface hardened layers, which will induce stress concentration, and thus lead to the deterioration of impact toughness. Preheating can increase the impact absorbed energy to about 28 J due to the softening of hardened layer.(5)In conventional quenching, the fracture surface appears to be quasi-cleavage fracture. In laser surface quenching, the intergranular fracture occurs at the vicinity of crack source, and the fracture surface appears to be quasi-cleavage at deeper posi tion. The interface between UCP and martensite matrix is the potential site for crack nucleation and propagation.Declaration of competing interestThe authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.CRediT authorship contribution statementZhen-xing Li: Conceptualization, Investigation, Writing - originaldraft. Bing-qian Tong: Investigation, Validation. Qun-li Zhang: Resources. Jian-hua Yao: Supervision, Conceptualization. VolodymyrKovalenko: Writing - review & editing.AcknowledgmentsThis work was supported by National Key R&D Program of China (2018YFB0407301); and Fundamental Research Funds for the Provin cial Universities of Zhejiang (RF-C2019003).References1H.K.D.H. Bhadeshia, Steels for bearings, pro, Mater. Sci. 57 (2012) 2684352Z.X. Li, C.S. Li, J.Y. Ren, B.Z. Li, J. Zhang, Y.Q. 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Trans. 48 (2017) 232023351.0C-1.5Cr钢双相组织的细化和性能双淬火和激光表面淬火相结合的处理摘要对1.0C-1.5Cr钢进行常规淬火和激光表面淬火处理。为提高表面硬度，获得更精细的显微组织，提出了一种双淬火和激光表面淬火相结合的工艺。研究了单淬火、双淬火和激光表面淬火奥氏体化过程中渗碳体的溶解和晶粒长大行为。结果表明:与单次淬火相比，二次淬火时未溶渗碳体颗粒的平均直径(UCP)更细，最终奥氏体晶粒尺寸(PAGS)减小近40%，约为4.5 m;激光表面处理后，表面附近的晶粒和渗碳体颗粒都会变粗。与单次淬火相比，双次淬火可使硬化层内PAGS减少至少11%，硬化层底部UCP平均直径减少约20%。与常规淬火相比，激光表面淬火使表面硬度提高了20%左右，提高了耐磨性。但在冲击过程中，表面的硬脆层往往是裂纹源，导致最终冲击韧性的恶化。在相同的激光参数下，单次和双次淬火时的冲击吸收能量基本相同，约为激光表面淬火前的25%。在激光表面淬火过程中，通过在160处预热，可以使冲击吸收能从22J提高到28J。关键词:双淬火;激光表面淬火;渗碳体1.介绍1.0C-1.5Cr (SAE 52100)钢是一种典型的过共析钢，广泛用于轴承、导轨和模具。为了满足复杂的工作条件，通常对该钢依次进行球化退火、淬火和低温回火处理，最终形成由马氏体基体、未溶解的渗碳体颗粒、残余奥氏体和纳米级碳化物组成的组织1,2。通过调整这些微观组织参数，可以控制其最终的力学性能。组织细化是提高1.0C-1.5Cr钢综合力学性能，特别是疲劳性能的有效方法。目前，围绕1.0C-1.5Cr钢组织细化的研究工作较多。例如，Beswick3报告说，可以通过在淬火前施加预冷变形来降低先前的奥氏体晶粒尺寸(PAGS)。因为在冷变形过程中会形成大量的低角度位错胞，在后续的加热过程中会提高奥氏体的形核速率，从而导致最终预奥氏体晶粒(PAG)的细化。Santos等人4提出感应加热和重复淬火对降低PAG也是可行的，他们的工作表明细化PAG可以明显提高疲劳性能。Mizobe等人5报道，850重复淬火可以使PAGS从15 m下降到7 m左右。Li等人6,7报道，热机械控制过程可以细化热轧和球化的组织，这将导致后续淬火组织中PAG和未溶解渗碳体颗粒(undissolved cementite particles, UCP)的尺寸略有减小。最近,Salloom et al。8提出并发PAG的细化和跟单信用证可以通过双重淬火过程中,由奥氏体化的1050油淬火和低温回火,紧随其后,再热二次奥氏体化温度850淬火及回火紧随其后。但在1050下进行奥氏体化时可能导致严重脱碳，后续油淬9时也可能形成微裂纹。Lee等人10报道在二次淬火中降低二次奥氏体化温度可以得到更细的PAG，而降低奥氏体化温度并不能提高最终硬度。因此，他们提出在奥氏体化第一阶段可以进行氮碳化。然而，这是耗时和昂贵的。在上述方法中，重复淬火，特别是二次淬火是细化1.0C-1.5Cr钢组织的极有前景的方法。为了进一步优化组织，提高力学性能，有必要围绕二次淬火过程中渗碳体的溶解和晶粒长大进行研究工作。需要注意的是，无论是单淬火还是双淬火都属于常规淬火处理，通常会导致1.0C-1.5Cr钢的贯穿淬火。然而，在许多工作条件下，它是表面将承受复杂的应力和磨损。因此，1.0C-1.5Cr钢的表面性能优化受到了特别的研究关注。提高表面性能的方法有几种，如表面热化学处理(渗碳、碳氮共渗、渗硼)、喷丸、物理/化学气相沉积、激光表面处理(淬火、熔覆、冲击喷丸)11-13。激光表面淬火具有成本低、效率高、可控性好等优点，是提高1.0-1.5Cr钢表面性能的一种很有前途的方法。Basu等人14开发了一种1.0-1.5Cr钢的双重热处理工艺，即等温淬火和激光表面淬火。结果表明，激光表面淬火可以在不影响贝氏体芯的情况下提高合金的表面硬度。Lin等人提出在1.0C-1.5Cr钢的激光表面重熔过程中可以施加水冷却，使其表面硬度提高，并形成超细组织。由于激光表面淬火对表面硬度的有利影响，可以替代碳氮共渗工艺用于双淬火处理。然而，目前将双淬火与激光表面淬火相结合的工作还很少。此外，还应注意这种钢的冲击韧性，特别是用于高速铁路或磁悬浮列车的轴承时16,17。但对激光表面淬火后的冲击性能研究较少。本文对1.0C-1.5Cr钢进行了常规单淬火、双淬火和激光表面淬火。研究了常规淬火和激光表面淬火奥氏体化过程中渗碳体颗粒的溶解、尺寸变化及晶粒长大行为。分析了1.0C-1.5Cr钢的硬度、耐磨性和冲击韧性。此外，还对预热的效果进行了研究。2.实验的程序2.1。材料其微观结构由球形渗碳体颗粒和铁氧体基体组成。对应的化学成分(wt.%)为:1.0C, 0.28Si, 0.34Mn, 1.57Cr, 0.011P, 0.003S。球化渗碳体颗粒的平均直径约为0.34 m。2.2。常规淬火和激光表面淬火为了研究渗碳体在单淬火和双淬火时的溶解和晶粒长大行为，制备了尺寸为3 10 11 mm的试样进行间断淬火，如图1。单次淬火时，将部分试样分别放入800和850的管式炉中进行奥氏体化处理，保温5 min、20 min、30 min、60 min，油淬至室温。二次淬火时，先将部分试样放入850(首次奥氏体化温度)的炉中保温30 min，油淬至室温。随后再次放入800(第二次奥氏体化温度)的炉中。分别保温5 min、20 min、30 min、60 min后油淬至室温。需要注意的是，本工作中二次淬火时的第一次奥氏体化温度和时间是固定的，因此下文提到的奥氏体化温度和时间是指第二次奥氏体化温度和时间。同时还制备了尺寸为11 24 100 mm的试样。同样放入温度为850的炉中保温30 min，然后油淬至室温。随后将其分为两组，一组在170下直接回火120 min，另一组再次放入800的炉中。800保温20 min后，油淬至室温，170回火120 min。回火后再加工至10 24 100 mm，去除表面脱碳层。随后，对其进行激光表面淬火。图2为激光表面淬火示意图。本文采用光斑尺寸为3 8 mm的光纤耦合二极管激光器(Laserline)。设置激光扫描速度为20mm /s，功率分别为1500 W和2000 W。此外，在激光表面淬火过程中，对部分试样进行了160的预热。表1进一步给出了激光表面淬火参数。2.3。微观结构表征所有的标本都经过了切割、研磨和抛光。随后，这些标本用4%的硝酸盐酸盐蚀刻。采用光学显微镜(Olympus BX53 M和Zeiss Axio Scope A1)、扫描电子显微镜(Zeiss PIGMA HV-01-043)、电子探针(JEOL JXA8530F)和透射电子显微镜(FEI Tecnai G2 F20)对其显微组织进行观察。用10%高氯酸和90%乙醇组成的电解抛光溶液制备TEM箔。此外，为了观察早期的奥氏体晶界，试样还在由饱和苦酸和少量Teepol润湿腐蚀剂组成的溶液中进行了65约70 s的腐蚀。本研究采用线性截距法测量奥氏体前期晶粒尺寸。为了量化UCP，预先用photoshop对渗碳体颗粒着色，然后用Image-pro Plus对其进行测量。假设UCP的体积分数等于其相应的面积分数18。每个UCP的平均直径是通过假设其面积为球形19来计算的。各条件下测定的渗碳体颗粒数在600 2500之间，这取决于奥氏体化参数。2.4。硬度、耐磨性和冲击韧性用显微硬度计(HMV-2TADW)测定硬度，载荷为0.3 kgf。耐磨性的评估使用的是一个销桌上磨损试验台。引脚为直径为5mm的氮化硅陶瓷球，如图3a所示。磨损试验在转速为500 rpm，旋转直径为5 mm的条件下进行。在磨损试验中，对试样施加1.5公斤的正常载荷。冲击韧性测试采用摆锤式冲击试验机(SANS ZBC2452-B)。16采用尺寸为10 10 55 mm的无缺口冲击试件。需要注意的是，对于激光表面淬火的试样，冲击试验时硬化层与摆锤相反，如图3b所示。3.结果3.1。单双淬时的显微组织3.1.1。渗碳体溶解和粒度图4为试样单次淬火和二次淬火的显微组织。可以观察到在马氏体基体中存在大量的球状UCP。复合材料的体积分数和尺寸分布对复合材料的最终力学性能有重要影响。通常，奥氏体化温度的升高和奥氏体化时间的延长会导致渗碳体的进一步溶解。认为渗碳体的溶解行为会影响UCP的尺寸变化。本文测量了不同奥氏体化条件下UCP的体积分数和平均直径，如图5所示。由图5a可以看出，随着奥氏体化时间从5 min延长到60 min, 800和850下单次淬火奥氏体化试样的UCP体积分数分别从10.4%降低到7.3%和8.0% 2.4%。而经800二次奥氏体化后，在二次淬火时则由6.8%降至3.4%。在相同的奥氏体化温度下，双淬火时UCP的体积分数比单淬火时小。此外,考虑到渗碳体的体积分数在最初的微观结构,它可以发现,跟单信用证的体积分数不断减少在整个auste-nitizing单一淬火的过程,而先增加,然后降低后期二次奥氏体化过程的双重淬火。换句话说，在奥氏体化早期，一些新的渗碳体形式可以在图4c中观察到。新形成的渗碳体颗粒尺寸较小，约为0.1 m。图5b显示了奥氏体化过程中UCP平均直径的变化。随着奥氏体化时间的增加，平均直径先减小后增大。在单次淬火的早期奥氏体化过程中，试样的平均UCP直径与单次淬火的体积分数和平均UCP直径的差异应归因于初始组织的不同。单次淬火时，初始组织为球状渗碳体和铁素体基体(图7a)。二次淬火时，初始组织主要为UCP和马氏体基体，为首次淬火后的组织。与铁氧体基体相比，马氏体基体是亚稳态的，含有大量位错和孪晶界。此外，碳在马氏体基体中过饱和，导致晶格严重变形。因此，即使在相当低的温度(160)下，纳米碳化物也能在几分钟内在位错或孪晶界处形成(图7b)，所有这些纳米碳化物都将转变为-碳化物(即渗碳体)20,21。在本研究中，单次淬火和双次淬火初始组织中渗碳体体积分数的实测值分别为19.4%和4.1%，进一步表明双次淬火初始组织为碳过饱和。3.1.2。一次淬火和二次淬火晶粒尺寸图8为试样单次淬火和二次淬火时的优先奥氏体晶界。认识到晶粒长大主要是由界面能的降低驱动的。生长过程是晶界迁移的结果，晶界迁移与原子的扩散密切相关22,23。当奥氏体化温度和时间越长，原子的有效扩散距离越大，PAGS越大。本研究进一步测量了PAGS。单次淬火时，随着奥氏体化时间从5 min增加到60 min, 800和850下奥氏体化试样的PAGS分别从7.6 m增加到8.1 m和7.8 m增加到8.6 m;两种奥氏体化温度下PAGS无明显差异。在800双淬条件下，奥氏体化试样的PAGS由4.1 m增加到4.7 m。与单次淬火相比，双淬火的PAGS含量降低了近40%。3.2。激光表面淬火后的显微组织3.2.1之上。硬化层的形貌试样经过单淬火和双淬火后进行回火和激光表面淬火。图9为激光表面淬火后的显微组织。硬化层除两侧外底部平整，这是因为两侧散热较快。在相同的激光参数下，无论是单次淬火还是双次淬火，硬化层的尺寸基本相同。当激光功率为1.5 kW时，硬化层宽度约为10点。当功率为2kw时，它会增加到10.6 mm左右，预热后，它的宽度会增加到11.1 mm。进一步观察硬化层的微观组织，如图10所示。表面UCP很少，这主要是由于较高的表面温度导致了明显的溶解。此外，随后的快速冷却速度阻止了渗碳体的形成，从而导致表面UCP较少。渗碳体的溶解会使基体中的碳处于过饱和状态。在250 m和625 m深度可观察到大量UCP。也就是说，激光表面淬火后得到了梯度组织。3.2.2。硬化层内渗碳体的溶解和尺寸测量了不同硬化层深度下UCP的体积分数和平均直径，如图11所示。很明显，随着深度的增加，UCP的体积分数增加，UCP的平均直径减小。增加激光功率或预热会加速渗碳体的溶解，并伴随渗碳体的粗化。此外，虽然初始试样分别通过单淬火和双淬火获得，但在相同激光参数下，两种试样在相同深度位置的UCP体积分数没有显著差异。但与单次淬火试样相比，双次淬火试样淬硬层底部UCP的平均直径明显变小，减少了约20%。也就是说，当初始微观结构中UCP的平均直径较小时，激光表面淬火后UCP的平均直径也较小。激光表面淬火前，单双淬火试样中UCP的体积分数分别为4.1%和4.2%，对应的UCP平均直径分别为0.24 m和0.17 m。激光表面淬火后，在较深的位置，单淬火和双淬火试样中UCP的体积分数至少为5.0%(图11a)，略大于激光表面淬火前。也就是说，在激光表面淬火过程中，较深的位置会形成一些新的渗碳体。但深孔处的平均直径与激光淬火前相当接近。总体而言，采用双淬火和激光表面淬火相结合的工艺可获得较细的单晶面淬火。渗碳体的溶解行为与激光表面淬火过程中的温度分布密切相关。然而，沿深度方向的温度分布测量是非常困难的。本文基于傅里叶热传导方程，利用ABAQUS计算温度场，如式(1)所示。其中，c和分别为密度，比热和热导率。T是温度，T是时间。Q为相变潜热。， c， 的值是通过JMatPro(7.0版)获得的。本文认为激光热源的光强分布是均匀的。此外，还使用了一个尺寸为10 - 24 - 100毫米的三维模型。图12为温度场计算结果。随着深度的增加，温度不断降低。激光功率为2kw时，表面温度在0.155 s内从100上升到最高温度(1479)，在0.555 s内下降到100，平均升温速率约为8896/s，冷却速率约为2484/s。加热速率和冷却速率均明显大于常规淬火。较大的升温速率使温度沿深度呈梯度分布，较大的冷却速率保证了硬化层的形成。此外，随着激光功率从1.5 kW增加到2 kW，接近表面的最高温度明显升高，而在更深的位置则略有升高(图12d)。当激光功率为2 kW时，1.0C-1.5Cr钢的表面温度超过了液相线(1460)，表明发生了局部微熔化。由于表面微熔化，经过预热后，表面附近的最高温度没有明显升高，而较深位置的最高温度明显大于未经预热的试样。温度沿深度的梯度分布导致了UCP体积分数和平均直径的变化。3.2.3。激光表面淬火后的晶粒尺寸图13进一步显示了激光表面淬火后不同深度的PAGS。对于单次淬火(SL1、SL2和SL3)试样，PAGS随深度的增加先减小后增大，达到8.5 m左右，这也是激光表面淬火前初始组织的PAGS。当深度为500 m时，PAGS会变小。对DL1、DL2和DL3双淬试样，PAGS随深度的增加先减小到4.5 m，超过500 m后略有变化;不同深度的晶粒尺寸应与温度有关。根据图12d和图13a，图13b进一步给出了不同深度PAGS与相应深度最高温度的函数关系。观察到，当最高温度超过800时，PAGS开始显著增加。在相同激光参数下，与单次淬火试样相比，双次淬火试样在125 m深度处的PAGS至少降低了11%。在500 m深度可降低约40%。也就是说，通过双猝灭和激光表面猝灭相结合的双相工艺，可以得到更精细的PAGS。3.3。硬度图14进一步显示了激光表面淬火后的硬度。激光表面淬火后，表面硬度明显提高。从硬度分布和显微组织来看，激光表面淬火后试样主要存在三个区域，分别为硬化层、热影响区和基体(图9)。硬化层是激光表面淬火过程中完成奥氏体相变的区域。热影响区靠近硬化层，在该区域奥氏体相变不足或未发生奥氏体相变。衬底区域位于最内侧，不受激光的影响。可以看出，随着激光功率从1.5 kW (SL1和DL1)增加到2.0 kW (SL2和DL2)，硬化层的厚度从约0.48 mm增加到约0.66 mm，而预热(SL3和DL3)使硬化层的厚度进一步增加到约0.72 mm。通过单次和双次淬火获得的初始试样，激光表面淬火前硬度分别为795 HV和750 HV左右，这也是激光表面淬火后基体区域的硬度。在1.5 kW和2.0 kW的激光功率下，硬化层的最大硬度约为1000 HV。激光表面淬火后的硬度比初始试样的硬度提高了20%左右。经过预热处理的试样硬度最大值约为870 HV，小于未预热时的硬度值，但仍明显大于激光表面淬火前的硬度值。值得注意的是，所有试样的表面硬度都低于亚表面硬度，这主要是由于激光淬火24时，表面组织发生粗化和脱碳。由于回火效应，热影响区硬度低于硬化层和基体区。此外，在相同的激光参数下，单次淬火和双次淬火试样的硬化层硬度相近。但与单次淬火试样相比，双次淬火试样的热影响区和基体区硬度均略低，与激光表面淬火前的硬度一致。3.4。耐磨性这种钢在实际工况下发生磨损是很常见的。在这项工作中，耐磨性是通过失重来评估，如图15所示。需要注意的是，磨损试验完全是在激光表面淬灭区域进行的。对于未进行激光表面淬火的初始试样，双淬火时相同摩擦时间下的磨损量较大，表明其耐磨性较差。这主要是由于双淬火试样的硬度低于单淬火试样的硬度。激光表面淬火后，两种试样在相同摩擦时间下的磨损量基本相同。但与未进行激光表面淬火的试样相比，其表面粗糙度要小。换句话说，激光表面淬火可以提高耐磨性，这应该归功于较高的表面硬度。3.5。冲击韧性1.0C-1.5C钢作为过共析钢，淬火后冲击韧性相对较低。采用10的v形缺口冲击试样时，其冲击吸收能约为4 J。为了比较冲击韧性，本文采用了非缺口冲击试样。表2为激光表面淬火前后试样的冲击吸收能。对于未激光表面淬火的初始试样，单、双淬火时的冲击吸收能均为84j左右，在相同的激光参数下，两种试样的冲击吸收能也基本相同。但与初始试样相比，激光表面淬火后的冲击吸收能量明显下降，约为激光表面淬火前的25%。当激光功率为1.5 kW和2.0 kW时，冲击吸收能量分别约为24 J和22 J。通过160的预热，可以略微提高到28j左右，这与硬化层的软化和残余应力13的降低有关。进一步观察冲击断口形貌，如图16所示。未进行激光表面淬火的试样(图16a和b)断口形貌为准解理断裂。在断口25中可以观察到大量的UCP。UCP的出现表明在UCP与马氏体基体的界面处形成了微裂纹。对于激光表面淬火试样(图16c和d)，淬硬层表面(箭头A)萌生微裂纹，并向基体传播。在裂纹萌生处附近，断口出现沿晶断裂，表明硬化层相当脆。在淬火层中间(箭头B)和热影响区(箭头C)，断口出现准解理断裂，与未进行激光表面淬火的试样断口相似。此外，淬硬层的断口较热影响区更光滑，进一步表明了淬硬层的脆性。根据断口形貌，激光表面淬火试样的断口可分为三个阶段。首先，当应力超过强度极限时，硬化层表面形成微裂纹。然后，裂纹在硬化层内迅速蔓延。最后，它进一步以较慢的速率膨胀，直至试样彻底断裂。马氏体和渗碳体颗粒之间的界面是裂纹形成和扩展的潜在位置。4.讨论4.1。渗碳体溶解和尺寸变化以下4.4.1。渗碳体一次淬火和二次淬火由于初始组织不同，渗碳体在单次淬火和二次淬火奥氏体化过程中的溶解行为不同。需要注意的是，在本研究中，二次淬火的奥氏体化过程具体指的是第二次奥氏体化过程(图1b)。单次淬火时，初始组织为球状渗碳体和铁素体基体(图7a)。而在二次淬火过程中，第二次奥氏体化前的初始组织主要由UCP和马氏体基体组成，这是通过第一次油淬火得到的。在单次淬火的奥氏体化过程中，由于温度升高，铁素体与奥氏体的碳溶解度差异显著，导致铁素体向奥氏体转变，渗碳体不断溶解。在铁素体向奥氏体转变后，渗碳体继续溶解为奥氏体，得到均匀分布的奥氏体。在加热阶段，渗碳体的溶解速率随时间增加，而在保温阶段26，渗碳体的溶解速率随时间逐渐降低。在二次淬火的奥氏体化过程中，在奥氏体转变完成之前会形成大量的纳米碳化物。淬火后在160回火4 h后，由于纳米碳化物21的析出，渗碳体体积分数从4.0%提高到14.1%。奥氏体转变完成后，这些新形成的汽车开始溶解为奥氏体。因此，在双淬火奥氏体化过程中，UCP的体积分数先增大后减小。渗碳体颗粒尺寸的变化取决于溶解行为、吉布斯-汤姆逊效应和纳米碳化物的析出等因素。在单次淬火的早期奥氏体化过程中，渗碳体溶解占主导地位，导致UCP平均直径减小。但后期渗碳体溶解缓慢，由于Gibbs-Thomson效应，渗碳体会以较小的渗碳体为代价而粗化，这与Ostwald成熟相似27,28。在双淬奥氏体化早期，特别是加热阶段，会形成大量的纳米碳化物，最终演变为微小的渗碳体颗粒。因此，在奥氏体化早期，UCP的尺寸较小。随着奥氏体化时间的增加，由于吉布斯-汤姆逊效应，这些渗碳体颗粒也会变粗。在单次淬火时，在5 min的奥氏体化时间，UCP呈双峰分布，这主要与初始球化组织中渗碳体的尺寸分布有关。据报道，由于在球化退火过程中，渗碳体在晶界处和晶内的生长速度不同，因此在球化组织中渗碳体的尺寸分布也是双峰的29,30。在单次淬火奥氏体化早期，随着渗碳体的溶解，小尺寸(0-0.15 m)和大尺寸(0.15-0.6 m)渗碳体的平均直径均减小。因此，UCP的规模分布仍然是双峰的。但在奥氏体化后期，由于吉布斯-汤姆逊效应，小颗粒的渗碳体会完全溶解，大颗粒的渗碳体不会溶解，反而会变粗。因此，小型UCP的数量分数会减少，而大型UCP的数量分数会增加。小尺寸渗碳体数分数与大尺寸渗碳体数分数的边界变得模糊，导致其双峰分布不明显。二次淬火时，在第二次奥氏体化初期会形成大量细小的渗碳体颗粒。因此，单位面积的渗碳体总数将相当大，而且小尺寸渗碳体的数量比例远远大于大尺寸渗碳体。大尺寸渗碳体引起的峰值很小，因此在双淬过程中双峰特征不明显。4.1.2。激光硬化层中的渗碳体激光表面淬火前的初始组织主要由渗碳体和低温回火马氏体组成。在激光加热过程中，马氏体转化为奥氏体，渗碳体溶解为奥氏体。温度场模拟结果表明，地表温度明显大于深部温度。当激光功率为2kw时，表面温度甚至超过液相线。因此，渗碳体的明显溶解发生在表面附近，且随着深度的增加，渗碳体的溶解程度减小，即UCP的体积分数增加。需要注意的是，随着渗碳体的显著溶解，UCP的总平均直径并不会减小。这主要是因为表面温度过高，较小尺寸的渗碳体颗粒会在表面附近完全溶解到基体中。因此，在表面附近，UCP的总平均直径较大。在常规淬火处理31中，在较高的奥氏体化温度下，未溶碳化物的平均直径较大。为了获得更细的UCP，表面温度不宜过高。与表面位置相比，较小尺寸的渗碳体颗粒在硬化层较深处由于较低的最高温度而不能完全溶解。在激光表面淬火前，通过二次淬火得到的初始组织中存在大量的小尺寸渗碳体。因此，二次淬火获得初始组织时，淬火层较深处UCP的总平均直径明显变小。4.2。晶粒尺寸最终的PAGS取决于成核速率和生长动力学的综合影响。通过提高奥氏体成核速率和降低生长速率可以降低PAGS。在常规热处理中，通过双淬可显著细化PAGS。这主要是因为二次淬火第二次奥氏体化过程的初始组织为亚稳马氏体。马氏体中的位错和孪晶都是奥氏体的潜在形核位点，有利于提高奥氏体的形核率，从而促进PAG的细化。对于两种初始组织，激光表面淬火后，表面附近的PAGS明显大于硬化层深处的PAGS。这主要是由于最大温度随深度的增加而减小。在较低的温度下，原子的扩散系数较小，不利于晶界的迁移。此外，在更深的位置有更多的UCP，这也阻碍了晶界的迁移。因此，与表面相比，PAGS在较深的位置趋于较小。但当深度扩展到一定值时，奥氏体的形核不足，或者由于温度较低而没有形成奥氏体。随着深度的增加，新形成的奥氏体晶粒数量减少。因为新形成的奥氏体晶粒尺寸相对较小。当其数量减少时，最终的PAGS会增加。值得注意的是，当深度超过500 m时，双淬试样的PAGS变化不大。这主要是因为二次淬火时初始组织的PAGS太小。新形成的奥氏体晶粒数的减少对最终的PAGS没有明显的影响。激光表面淬火后，二次淬火试样的PAGS明显减小，这与较高的奥氏体形核速率有关。激光表面淬火前，双淬试样的PAGS要小得多(图8)，因此单位体积对应晶界面积更大。晶界是新奥氏体晶核的潜在形核位置。激光表面淬火过程中，随着初始组织晶界面积的增加，奥氏体形核速率也会增加，有利于获得更细的PAG。4.3。激光表面淬火后冲击韧性下降激光表面淬火可提高表面硬度和耐磨性。然而，冲击韧性变差。这应归因于激光硬化层的脆性。在冲击过程中，硬化层将承受拉应力。一旦拉应力超过断裂极限，脆性硬化层的外表面就会形成微裂纹。微裂纹尖端附近的应力集中会进一步加速裂纹的扩展，最终导致微裂纹的彻底断裂。对硬化层进行160预热软化，有利于降低其脆性，从而延缓裂纹的形成。因此，在激光表面淬火过程中通过预热可以提高冲击韧性。但预热会导致冷却速度降低，导致硬化层硬度降低。此外，当预热温度超过马氏体开始转变时，马氏体转变也会受阻，削弱激光表面处理的强化效果。考虑到1.0C-1.5Cr钢淬火后一般回火温度在170以下，预热温度也应低于170，这样不会显著降低基体的硬度。根据以上结果，可以将双淬火与激光表面淬火相结合，获得更精细的组织。组织细化可以提高1.0C-1.5Cr钢的疲劳性能4,10。与单次淬火试样相比，双次淬火试样的硬度较小，而激光表面淬火
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