Regular Article
Influence of Tempering Temperature on Microstructures and Tensile Properties of a Cr–N Alloyed Medium-Mn Martensitic Steel
2024 Volume 64 Issue 2 Pages 277-285
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2024 Volume 64 Issue 2 Pages 277-285
In this paper, the influence of tempering temperatures on microstructures and tensile properties of a Cr–N alloyed medium Mn martensitic steel was studied. The microstructures formed after the tempering below 400°C are composed of recovered martensite as the matrix, ultrafine retained austenite (RA) and carbonitrides. The tempering at 100°C led to the best combination of 2080 MPa ultrahigh ultimate tensile strength (UTS) and 15% total elongation (TE), which is attributed to the prominent strain hardening capacity caused by both the gradually release of internal stress and the pronounced austenite-to-martensite transformation. The tempering at 400°C resulted in the rapid increase of yield strength (YS) by ~500 MPa due to the relief of internal tensile stress and annihilation of dislocations and the best ductility because it produced the most stable RA grains with the highest C concentration for a sustainable austenite-to-martensite transformation over the large plastic straining. The further increase of temperature to 650°C caused ferrite formed, which decrease both YS and strain hardening rate, leading to the lowest UTS. Moreover, it was found that higher N content increased YS but had little influence on both UTS and TE because it mainly contributed to enhanced precipitation of carbonitrides. It is then concluded that the strength and ductility of medium Mn martensitic steel could be increased by increasing the strain hardening capacity through tailoring both the internal stress in martensite and the mechanical stability of RA via a proper tempering treatment.
Medium-Mn steels (MMS) that usually contain Mn contents from 4 to 10 wt.% have attracted extensive interest due to their outstanding mechanical properties achieved at relatively low alloy cost, and they are considered as the promising 3rd generation advanced high strength steels (AHSS).1) The frontier research on them mainly focus on evading the trade-off of strength and ductility. For example, large amount of forest and mobile dislocations were introduced into one medium Mn steel by a Deformation & Partition process, leading to the ultrahigh yield strength (YS) of 2.2 GPa and total elongation (TE) of 16%.2) Moreover, the ultrafast heating was used to regulate the chemical boundaries of low-carbon medium-Mn steel, which improved the ultimate tensile strength (UTS) and TE to 2.3 GPa and 17%, respectively.3) Recently, a hierarchical microstructure was introduced in,4) which consists of laminated and twofold topologically aligned martensite with finely dispersed retained austenite (RA) that simultaneously activates multiple micro-mechanisms to strengthen and plasticize the medium Mn steel, resulting in the UTS of 2.2 GPa and TE of 18.0%.
Besides mechanical performance, corrosion resistance is also important for high strength steel in automobile, bridge and other structural applications.5) It was reported that increasing the Cr content in medium Mn steels could enhance their protectiveness on NaCl solution spray erosion.6) Furthermore, the alloying of 0.25 wt.%Cr and 0.12 wt.%Mo in a 0.2C–7Mn–3Al (weight percentage unless mentioned elsewhere) steel increased the YS and UTS to 600 MPa and 1600 MPa by precipitation hardening and accelerating the austenite-to-martensite transformation.7) Therefore, the Cr alloying in medium Mn steel should be beneficial in improving both mechanical and corrosion resistance performance. Recently, our research group developed a novel Cr alloyed medium Mn steel with the composition of 0.25C–6.94Mn–0.38Si–3.29Cr–0.23Al–0.12(Nb+V) and a processing route that leads to UTS of 1.8 GPa and TE of 9% and exhibited an enormously reduced oxidation attack during press hardening comparing with the typical press hardening steel (PHS) grade of 22MnB5.8) However, the mechanical performance of these reported corrosion- or oxidation-resistant Cr alloyed medium Mn steel are far below the best mechanical properties achieved in the usual medium Mn steels.
Nitrogen is another typical alloy that was widely used in stainless and wear-resistant steels to enhance both, strength and corrosion resistance.9,10) Moreover, the N alloying in high Mn steel could suppress the occurrence of PLC phenomenon (dynamic strain ageing)11) that often appears in medium Mn steels and leads to surface quality problem during their forming.12) However, no N alloyed medium Mn steel has been reported up to now. Therefore, a Cr+N alloyed medium Mn martensitic steel has been designed in this paper to further improve the mechanical performance of corrosion resistant medium Mn steel. Moreover, the influence of tempering temperatures and N alloying on YS and the strain hardening behavior of this newly developed martensitic steel will be studied and discussed.
The nominal composition of studied steel is Fe–0.2C–7Mn–0.88Si–3Cr–0.2V–0.09N. Such a compositional design is based on the targeted microstructure, which should comprise 10–20% fraction of RA in the martensitic matrix after the solution and quenching treatments. This duplex microstructure led to the superior mechanical properties based on our previous research on ultrastrong and ductile steels.2) Therefore, both C and Mn contents together with Si contents were all carefully designed for ensuring the martensitic matrix and desired RA fraction by quantifying their influences on Ms temperature. High contents of Cr and N is to increase the strength and corrosion resistance, for which, the maximum solute concentration of N in steel was calculated by using the equation proposed in Ref. [13]. The alloying of V is for precipitation strengthening. The equilibrium phase fractions, N/Mn contents in austenite and the compositions of (VCr)(CN) at different temperatures were calculated using ThermoCalc 2019 and TCFe9 database, and the calculated results are shown in Figs. 1(a) through 1(c). It is seen that the possible precipitation formed in the temperature range of 800°C to 1200°C is (VCr)(CN), while the Cr2N, M7C3 and M23C6 could precipitate probably below 800°C (Fig. 1(a)). The equilibrium N content of austenite increases but the Mn content decreases with the increase of temperature (Fig. 1(b)); The N content in (VCr)(CN) is much higher than the C content, the latter gradually decreases with the decreasing temperature particularly below 700°C (Fig. 1(c)).
The designed steel was melted in a 100 kg vacuum induction furnace and cast into ingots, which were then hot-forged into 60 mm thick billets. The billets were heated to 1200°C for 2 h for the solution treatment, and then hot rolled to a thickness of 4.5 mm by eight passes in a pilot two-high hot rolling mill with a finish temperature of 900°C, followed by air cooling to room temperature. Finally, the hot-rolled strips were tempered at 100°C, 400°C and 650°C for 1 h and cooled down to room temperature in air. The achieved hot rolled strips before and after the tempering at these temperatures are termed HR, HR100, HR400 and HR650 respectively.
The standard tensile samples were machined from steel strips along rolling directions with a gauge length of 25 mm and a width of 6 mm. Uniaxial tensile tests were carried out at a displacement rate of 1 mm/min on a WDW-200D tensile testing machine. All the reported tensile properties were taken as the averages of two repeated measurements. The microstructures on the samples were examined by FE-SEM (field emission scanning electron microscope), JSM-6710F, at 20 kV after polishing and etching in 4% nitric acid for 20 s. Electron backscattered diffraction (EBSD) examination was performed on an Auger Nano probe combined with Electron Backscatter Diffraction (Nano-Auger/EBSD), PHI 710, at 20 kV. The data were processed by EDAX OIM software. The X-ray diffraction (XRD) measurements were performed in the Bruker D8 Discovery with Co K-alpha radiation. The volume fraction of RA was calculated from the integrated intensities of (200)bcc, (211)bcc, (200)γ, (220)γ and (311)γ diffraction peaks. The C contents of RA are calculated from the XRD profiles using the method reported in Ref. [14]. During the XRD examination, a two-dimensional detector was selected with the exposure time of 1 s. The raw two-dimensional data was then converted into one-dimensional diffraction pattern by virtue of the commercial software EVA (Bruker). The instrumental broadening was determined by NIST Al2O3 sample. Samples for EBSD and XRD examination were electro-polished in a solution of 20% perchloric acid and 80% ethanol at room temperature. Transmission electron microscopy (TEM) observations were performed in a FEI Tecnai F30 at 300 kV. The foils for TEM were made by first mechanical grinding to a thickness of 40 μm and then twin-jet electro-polishing in a solution of 5% perchloric acid and 95% ethanol (vol.%) at around −20°C.
The engineering stress-train curves of studied HR, HR100, HR400 and HR650 are shown in Fig. 2(a) and their tensile properties including YS, UTS and TE are compared in Fig. 2(b). When the tempering temperature increases to 100°C, YS and UTS change little, while TE increases by 6%; when the tempering temperature increases to 400°C, TE is continuously improved to 21% and yield strength increases substantially by 498 MPa but UTS decreases by 410 MPa, i.e. the strain hardening increment is minimized dramatically; when it further increases to 650°C, YS, UTS and TE all start to decrease significantly.
From the obtained XRD spectra in Fig. 3(a), both the fractions and C concentrations of RA in the four specimens were calculated and the results are shown in Figs. 3(b) and 3(c) respectively. The RA fractions of HR, HR100, HR400 and HR650 are 16.5±1.2%, 16.4±1%, 18.0±0.7% and 22.9±0.8% respectively. This indicates that the tempering below 400°C has little influence on the RA fraction, which increases remarkably only when the tempering temperature increases to 650°C (Fig. 3(b)). In contrast, the C concentration of RA firstly increases with the increase of tempering temperature from 100°C to 400°C, and then decreases when the further increase to 650°C. As a result, the C concentration of RA in HR400 is highest among all the studied samples (Fig. 3(c)).
EBSD results on the microstructures of HR100, HR400 and HR650 are shown in Figs. 4(a) through 4(c). It is seen that the prior austenite grains in HR100 and HR400 were elongated along the rolling direction, indicating the prior austenite were unrecrystallized and work-hardened after the finish rolling (Figs. 4(a), 4(b)). The microstructures of HR100 and HR400 are both composed of martensite as the matrix and RA grains, the latter have a mean diameter of 320±100 nm (Fig. 4(a)) in HR100 and 370±190 nm in HR400 (Fig. 4(b)). This is because the austenite grains can grow at the high tempering temperature of 400°C, and grow even coarser to 433±347 nm at 650°C. The coarsening of austenite grains during the tempering diluted their solute C/Mn concentration and reduced their chemical stability in general; consequently, some RA grains transformed to martensite and others remained during quenching. Moreover, the significant solute partitioning between martensite and austenite and recrystallization of some martensite grains both occurred during the tempering at 650°C, leading to strongly recovered martensite, RA and ferrite formed in HR650, see Fig. 4(c). In Figs. 4(d), 4(e), 4(f), 4(g), precipitates were all found in the martensitic matrix of HR, HR100, HR400 and HR650 specimens, indicating that they should be formed during the hot rolling rather than the tempering at different temperatures. Moreover, the largest quantity of precipitates were observed in HR650 (Figs. 4(e), 4(f), 4(g)), probably because the tempering at 650°C promotes the precipitation. This shall be studied more by TEM as discussed later.
Figures 5(a)–5(d), Figs. 6(a)–6(d) and Figs. 7(a)–7(d) are the TEM bright/dark field images and EDS mappings on the microstructures in HR100, HR400 and HR650 respectively. Both dislocations and precipitates were observed within martensite laths of HR100 having the 100–200 nm width (Figs. 5(a), 5(b) and 5(c)). The EDS mappings in Fig. 5(d) show that the precipitates are V–Cr enriched carbonitride, indicating the precipitates in HR100 are (V,Cr)(NC), which should be formed in the temperature range of 800–1200°C during hot rolling and almost kept unchanged after the tempering at 100°C (Fig. 1(a)). In contrast, dislocation cells were seen in HR400, indicating that a stronger recovery should have occurred during the tempering at 400°C (Figs. 6(a), 6(b) and 6(c)). The EDS mapping in Fig. 6(d) shows the poor presence of C in the precipitate, indicating that it is a nitride of (V,Cr)N more than the carbonitride in HR400 because V/Cr contained nitride is more thermodynamically stable than the carbide, as seen in the decreasing equilibrium C content of (V,Cr)(C,N) with the increase of temperature in Fig. 1(c). Few carbides were found after the tempering at 400°C because most of carbides were dissolved and the released C atoms partitioned into austenite, leading to the highest C concentration in the austenite of HR400. More precipitates having a wide size range of 46–113 nm were observed in HR650 (Figs. 7(a), 7(b) and 7(c)). Among them, the coarser ones are more enriched in C, Mn and Cr (Fig. 7(d)), i.e. close to carbide as indicated by the calculated equilibrium composition at 650°C in Fig. 1(a); while the finer ones are more enriched in N, Cr and V, i.e. (V,Cr)N, which is same to those in HR400. Moreover, the statistic average diameter and the number density of precipitates were as 40±10 nm and 1.7×106 μm−2 in HR100, 65±15 nm and 4.4×106 μm−2 in HR400, and 83±28 nm and 7.7×106 μm−2 in HR650, which is expected to make different contribution to strengthening.
YS (σ0.2) changed little after the tempering at 100°C but increased by 498 MPa at 400°C, which may result from the difference in precipitation hardening. Precipitation hardening increments (σP) in HR100 and HR400 are calculated by using the following Ashby-Orowan equation based on the mean diameter and number density of precipitates.15)
| (1) |
Where G is the shear modus, 83 GPa; b is the Burgers vector, 0.248 nm; X is the mean diameter of precipitates; Vf is the volume fraction of precipitates. The precipitation hardening increments in HR100 and HR400 are calculated as 55 MPa and 72 MPa, the difference between them is much smaller than the one between their YSs. Thus, the precipitation hardening is not the main reason for strengthening.
It has been demonstrated that the yielding of martensitic steel that is composed of martensite and a small fraction of ultrafine austenite grains starts from the plastic deformation in martensitic matrix,16) because the ultrafine austenite grains, embedded in the martensitic matrix like islands, cannot independently bear the applied strain but coordinate the micro-strain with the neighboring martensitic matrix during yielding.17,18) Hutchinson19,20) stated that the yielding of martensitic steel was significantly influenced by the gradual release of internal stress, which is confirmed by the decrease of XRD line width during the yielding. Since higher internal tensile stress leads to yielding at lower tensile strength applied externally, the tempering at 400°C should relieve the internal stress more than that at 100°C, leading to YS raised with the increase of temperature from 100°C to 400°C. However, the release of internal stress should not be the only factor that influence the YS of studied steel, because the tempering at 100°C also released the internal stress in some extent21) but only increased the YS by 20 MPa. A. J. McEvily et al. attributed the yielding of martensitic steel to the slip of mobile dislocations that had been generated due to the austenite-to-martensite transformation during quenching.22) Although the recent results by atom probe tomography23) showed an extensive segregation of carbon atoms into dislocations, i.e. the Cottrell atmospheres should be formed already during the quench to immobilize dislocations, it cannot be ruled out that some dislocations may not be locked and still have mobility. It is noted that the tempering at 400°C resulted in some mobile dislocations in the quenched martensite annihilated due to more significant recovery, as confirmed by the formation of dislocation cells within martensite laths in HR400 (Fig. 6(a)). In this case, few mobile dislocations are available to slip and then be multiplicated at low stress level, resulting in the increase of elastic limit. In contrast, limited recovery occurred during the tempering at 100°C (Fig. 5(a)). In summary, both the pronounced release of internal stress and the significant annihilation of mobile dislocations in martensite during the tempering at 400°C led to much larger increase of YS of HR400 than at 100°C.
Different from HR100 and HR400, the microstructures of HR650 consist of ferrite, RA, and fresh martensite (Fig. 4(c)). The formation of quenched martensite led to many mobile geometrically necessary dislocations (GNDs) generated in ferrite near the martensite/ferrite interface due to the volume expansion associated with γ-to-α’ transformation during cooling.24) These mobile dislocations within ferrite can easily slip at lower stress due to much lower C/N concentrations than that in martensite. As a consequence, the increase of tempering temperature from 400°C to 650°C led to the rapid decrease of YS by 640 MPa.
4.2. Strain Hardening MechanismWork hardening rates varied during the tensile deformation, as seen in Fig. 8(a). All the strain-hardening rates first decreased rapidly, and then remained almost constant until fracture. Among them, the HR specimen exhibited the initially highest strain hardening rate but decreased most rapidly. HR400 exhibited the lowest average strain hardening rate but the most sustainable strain hardening capacity. The strain hardening rate of HR650 is lowest at the beginning, but then increased to a level between HR400 and HR100 until fracture.
It is well known that the austenite-to-martensite transformation can significantly influence the strain hardening capacity of medium Mn steel. The RA fractions of studied steel interrupted at the 2% tensile strain and after fracture were further examined by using XRD method, and they are shown in Fig. 8(b) together with the RA fractions before deformation. The smallest and largest RA fractions were transformed in HR400 and HR650 during the tensile straining to 2%, suggesting the highest and lowest mechanical stability of RA in HR400 and HR650. The mechanical stability of austenite should be influenced by various aspects, including C, Mn, N concentrations and grain size et al. The equilibrium N concentrations of austenite at 400°C and 650°C are calculated as 0.00112% and 0.00271% (Figs. 1(b) and 1(c)). This indicates that the N content of austenite in HR650 should be higher than that in HR400, but the austenite in HR650 exhibited much lower stability than that in HR400. Therefore, the N enrichment in austenite should not dominant the difference in austenite stability of studied samples. The Mn concentration of austenite in HR400 and HR650 were further examined by EDS, which are 7.46 ±0.8 wt.% and 6.85±0.6 wt.%, indicating the lower Mn content of austenite in HR650 (Fig. 9). That is because the coarsening of austenite diluted the C/Mn concentration and the formation of large amount of carbides led to the C/Mn atoms were enriched in precipitates after tempering at 650°C (Fig. 7). Moreover, the EBSD and XRD examinations have shown that the austenite grains in HR650 are coarsest, while the C concentrations of austenite in HR400 is highest (Figs. 3 and 4). Therefore, the highest mechanical stability of RA in HR400 due to the most significant C partitioning between martensite and austenite during tempering at 400°C (Fig. 3(c)) contributed to more sustainable strain hardening (Fig. 8(a)), resulting in the largest TE. In contrast, the RA grains in HR and HR100 transformed more rapidly due to their lower C concentrations. In HR650, the coarsest austenite grain size and the lower C/Mn concentration led to the lowest austenite stability, resulting in the most RA grains transformed to martensite after the tensile straining to 2%.
Moreover, the release of residual stress and the dislocation multiplication in a phase cause the narrowing and broadening of full width at half maximum (FWHM) of its crystal plane respectively.20) Therefore, the FWHMs of (200)bcc and (211)bcc peaks of four samples before deformation and at the tensile deformation to 2% strain were measured by XRD to reveal the changes in both internal stress and dislocation density in martensite, and their results are summarized in Table 1. It can be seen that the FWHM of (200)bcc and (211)bcc peaks in HR and HR100 both decreased but changed little in HR400 after the tensile straining to 2%, indicating the internal stress in martensite gradually released in HR and HR100 at the beginning of tensile deformation but not in HR400. This is because more internal stress was released after the tempering at higher temperature of 400°C, as confirmed by the rapid decrease of FWHMs of both (200)bcc and (211)bcc peaks with the increase of tempering temperature from 100°C to 400°C (Table 1).
| Samples | Full width at half maximum (FWHM) | |||
|---|---|---|---|---|
| (200)bcc-0% | (200)bcc-2% | (211)bcc-0% | (211)bcc-2% | |
| HR | 1.837 | 1.524 | 1.536 | 1.353 |
| HR100 | 1.758 | 1.663 | 1.518 | 1.446 |
| HR400 | 1.329 | 1.362 | 1.382 | 1.305 |
| HR650 | 0.755 | 1.022 | 0.855 | 1.116 |
It is then summarized that both the more rapid austenite-to-martensite transformation and the release of internal stress in martensite during the tensile test contributed to the high strain hardening rate in HR and HR100, leading to the ultrahigh UTS over than 2 GPa. However, higher internal stress could lead to more heterogeneous plastic deformation and even the premature fracture, like in HR. In contrast, almost no internal stress remained prior to the tensile test in HR400 due to significant recovery and accordingly, the RA grains transformed most slowly due to the highest mechanical stability. This contributed to the lowest but most sustainable strain hardening capacity, leading to the largest TE but lower UTS. In contrast, the moderate residual internal stress in martensite and RA grains having a proper mechanical stability in HR100 both resulted in the best combination of strength and ductility among all the specimens.
In Table 1, the FWHMs of (200)bcc and (211)bcc peaks in HR650 increased after the straining to 2%, indicating the dislocations in bcc phase should be multiplicated. HR650 is different from three other specimens as it contains ferrite grains. These ferrite grains are always preferentially deformed in the multi-phase steels,25) which reduced the plastic strain allocated into martensite and resulted in much lower strain hardening rate at the beginning of deformation than other specimens that are composed of martensite and austenite.26) Therefore, HR650 has the lowest YS and UTS among all the specimens due to the presence of ferrite.
4.3. Influence of N Alloying on Tensile Properties of Medium Mn Martensitic SteelThe tensile properties of studied Cr/N alloyed medium Mn steel appears not much influenced by the alloying of N as expected, after we compared the tensile properties of HR200 with another steel, which has the similar composition except for no alloying of N (Fe–0.22C–6.94Mn–0.38Si–3.29Cr–0.12V) and was prepared by the same hot rolling process and tempered at 200°C for 1 h (termed HR200-NF), see Fig. 9. It is seen that the N alloying leads to the increase of YS by 130 MPa but the similar UTS and TE. The precipitation hardening increment of HR200 should be 55–72 MPa, since that the calculated values in HR100 and HR400 are 55 MPa and 72 MPa respectively; while the solid solution strengthening increment of HR200 due to N alloying is then estimated as 58–75 MPa by subtracting precipitation hardening from the total YS increment.
HR200 and HR200-NF exhibit the similar strain hardening behavior due to the following reasons except for the different YS. Although the nominal N content of HR200 is up to 0.09 wt.%, the N content of austenite should decrease to approximate 0.025% at the finish rolling temperature of 900°C before quenching (Fig. 1(b)) due to the precipitation of nitride and carbonitride. In other word, most of N atoms are to form precipitates during hot rolling (Fig. 1(c)). Next, the remaining solute N content of austenite in HR200 before quenching may be still higher than that in HR200-NF. The higher N content of austenite should lower Ms temperature so that less austenite transformed to martensite during cooling, generating lower internal stress in HR200 than HR200-NF, leading to smaller strain hardening rate and then lower UTS; meanwhile, the stability of RA grains in HR200 is enhanced due to higher N content, contributing to more sustainable austenite-to-martensite transformation and improved strain hardening during the deformation of HR200. As the consequence, the two influences compensate for each other, leading to the similar strain hardening behavior. The exact influence of N alloying, however, requires more studies in future.
In this paper, a Cr/N both alloyed medium Mn martensitic steel was manufactured after the tempering for different temperatures, and the resultant microstructures and mechanical properties were studied. The following conclusions could be drawn.
(1) After the tempering of Cr/N alloyed medium Mn martensitic steel at a temperature no higher than 400°C, the resultant microstructures consist of tempered martensite, ultrafine RA grains and carbonitrides; whilst the tempering at 650°C caused the soft ferrite formed, which caused YS and UTS both rapidly decrease to minimum values.
(2) The tempering at 100°C resulted in the best strength-ductility synergy, including ultrahigh UTS of 2080 MPa and TE of 15%, due to the prominent strain hardening resulting from both the austenite-to-martensite transformation and the gradual release of internal stress. The increase of tempering temperature from 100°C to 400°C led to the rapid increase of YS by ~500 MPa and the best ductility. The former is because the internal tensile stress was relieved and dislocations annihilated, and the latter because RA grains in HR400 have the highest mechanical stability due to the highest C concentration so that the austenite-to-martensite transformation can be sustained over the large plastic straining.
(3) The N alloying in studied Cr-alloyed medium Mn martensitic steel leads to the increase of YS by 130 MPa but just has a marginal influence on UTS and TE, which is not so pronounced as expected.
(4) The contribution of precipitation hardening to YS of studied Cr/N alloyed medium Mn steel is estimated only 55–72 MPa after the tempering at 100–400°C. In particular, both UTS and TE can be improved by enhancing the strain hardening capacity via tailoring the internal stress generated due to martensite transformation and the mechanical stability of RA grains with a proper tempering treatment.
Haiwen Luo and Bin Hu acknowledge financial support from National Natural Science Foundation of China (Nos. 51831002, 51904028, and 52233018) and Fundamental Research Funds for the Central Universities (No. FRF-EYIT-23-08).
