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Influence of Ti microalloying on the formation of nanocrystalline structure in the 201L austenitic stainless steel during martensite thermomechanical treatment
Materials Science & Engineering A 584 (2013) 177–183
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Influence of Ti microalloying on the formation of nanocrystalline structure in the 201L austenitic stainless steel during martensite thermomechanical treatment S. Sadeghpour a,n, A. Kermanpur a, A. Najafizadeh b a b
Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran Fould Institute of Technology, Fouldshar, Isfahan, Iran
art ic l e i nf o
a b s t r a c t
Article history: Received 26 April 2013 Received in revised form 27 June 2013 Accepted 11 July 2013 Available online 19 July 2013
The martensite thermomechanical treatment was used for the formation of nano/ultrafine grain structure in the 201L austenitic stainless steel containing 0.12 wt% Ti microalloying element. The initial microstructure was provided through homogenizing, hot rolling and solution annealing of the as-cast ingots. The specimens were then cold rolled between 5% and 90% thickness reduction and subsequently annealed at 750–900 1C for various times. The results showed a promoting effect of Ti on the formation of strain-induced martensite (SIM). A nanocrystalline austenitic structure with average grain size of 45 nm was achieved by annealing at 900 1C for 60 s through a diffusional transformation mechanism. It was found that precipitation of nanosized TiC particles during the reversion annealing could retard the reversion process and suppressed grain growth in further annealing times. The tensile testing of the thermomechanically processed specimens showed a good combination of high yield strength ( 1000 MPa, six times higher than that of the initial coarse-grained steel) and excellent ductility (42% total elongation) for the Ti microalloyed 201L steel due to the SIM formation during the deformation and impressions of the nanosized Ti carbides distributed within the nano/ultrafine grain structure. Published by Elsevier B.V.
Keywords: 201L stainless steel Ti microalloying Nano/ultrafine grained structure Martensite thermomechanical treatment Strain-induced martensite Reversion annealing
1. Introduction Extensive research has been carried out on grain refinement of metals and alloys down to the nano-grain (NG, d o100 nm) and ultrafine-grain (UFG, do 500 nm) scales in order to improve their mechanical properties. Whereas significant increases in hardness and strength have been reported for nanostructured materials, they often exhibit a very low tensile elongation to failure especially uniform elongation (strain before necking) [1,2]. The reason for the low ductility in NG/UFG metals is plastic instability due to the limited strain hardening capacity that arises from the fact that extremely small grains cannot store dislocations to increase their density by order of magnitude as normally possible in the coarsegrained (CG) metals [3]. It has confirmed that any improvement of the strain hardening ability as a stabilizing mechanism to bring the plastic instability under control will be beneficial to enhancing the homogeneous plastic deformation for NG and UFG materials [4]. Several approaches have been developed in terms of microstructural design for improving ductility of the NG and UFG metals
n
Corresponding author. Tel.: +98 0936 9159 669; fax: +98 0311 3912 753. E-mail addresses: [email protected], [email protected] (S. Sadeghpour). 0921-5093/$ - see front matter Published by Elsevier B.V. http://dx.doi.org/10.1016/j.msea.2013.07.014
without sacrificing their strength. These include dispersion of nanoparticles and nano-precipitates in the nanostructured metal matrix [5], to use transformation induced plasticity (TRIP) [6] and twinning induced plasticity (TWIP) [7] effects, enhancing the fraction of high angle grain boundaries and lowering the density of dislocations [8], and improving strain hardening rate by tailoring the stacking fault energy (SFE) [9]. A recent thermomechanical route of the formation of NG/UFG structures in metastable austenitic steels called martensite treatment is characterized by strain induced martensitic transformation and its reversion to austenite [10–13]. In this approach, severe deformation of austenite at room temperature leads to strain induced transformation of austenite (γ) to martensite (α′) and upon annealing, this heavily deformed strain induced martensite (SIM) transforms back to austenite leading to a noticeable grain refinement. Because of low SFE, austenitic stainless steels (ASSs) are susceptible to deformation induced martensitic transformation. When ASSs are deformed, strain induced ε-martensite with the hexagonal close packed (hcp) structure and α′-martensite with body centered cubic (bcc) structure are formed from the metastable austenite. The formation of SIM is related closely to shear bands which are planar defects associated with overlapping of stacking faults on {111}γ. Shear band intersections act as the nucleation sites for
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α′-martensite, and the formation of the shear bands is a precursor for the strain induced α′-martensite transformation [14]. Several factors including grain size, strain rate, deformation temperature, mode of deformation, etc. may influence the martensitic transformation [15]. In addition to these factors, steel chemistry will also bear heavily on martensitic transformation. This has been explained in terms of the variation in the chemical free energy difference between austenite and α′-martensite phase as the chemical driving force. But Talonen and Hanninen [14] have suggested that the stability of austenite against martensite transformation is influenced more by SFE than the chemical driving force. It is well known that the SFE is strongly dependent on the chemical composition [16,17] and temperature [18,19]. Many authors have investigated the influence of temperature on the formation of SIM in the austenitic steels [14,20]. The effect of composition on the formation of SIM was first investigated by Angel [21], who correlated elemental compositions with the temperature at which an amount of 50% austenite will be transformed to martensite by the application of 0.30 true strain ( 35% engineering strain), denoted by Md30. Nohara et al. [22] modified Angel′s equation to include additional elements and grain size as follows (in wt%): M d30 ¼ 551462 %C9:2 %Si8:1 %Mn 13:7 %Cr29 %Ni18:5 %Mo29 %Cu 68 %Nb462 %N1:42 ðGS 8:0Þ where GS is the ASTM grain size number. A low Md30 temperature indicates high austenite stability against the formation of SIM. Moreover according to Talonen et al. [14], the temporal and compositional dependence values of the austenite stability are accounted by variation in the SFE. As mentioned before, the severe cold deformed SIM transforms back to austenite upon annealing. The α′ to γ reversion transformation has been studied by a number of authors [23–27]. This reversion may proceed either through shear or diffusion-controlled mechanisms. The alloy chemical composition plays a key role in determining the α′ to γ reversion mechanism. Takaki et al. [28] reported that the reverse transformation of martensite to austenite occurs by diffusionless mechanism at the low ratio of Cr to Ni, whereas diffusive reverse transformation happens at the high ratio of Cr to Ni in metastable Fe– Cr–Ni ternary alloys. In steels which undergo martensitic shear reversion, reversion completes during heating and reversed austenite contains a high density of dislocations just after reversion [29]. During successive annealing, grain refining proceeds through the process of recrystallization of reverse austenite. On the other hand, diffusional reversion occurs primarily by nucleation and growth of fine austenite grains at martensite lath boundaries [28]. Diffusion-type mechanism is characterized by (i) a wide annealing temperature range; (ii) the formation of defect free equiaxed austenitic grains which grow in size with time; (iii) a wide austenite grain size distribution; and (iv) possible formation of secondary phase precipitates [30]. In recent years the martensite thermomechanical process using SIM and its reverse transformation has been extensively investigated to produce NG/UFG structure in the conventional ASSs especially 300 series [10–12,31–33]. The martensite treatment has been modified to obtain finer grain sizes and better mechanical properties with combinations of recrystallization, warm deformation, and repetition of the process [34–36]. In addition Lee et al. [37] reported that the nanosized Nb particles increased the strength without reducing ductility of the UFG steel processed by martensite treatment. Nevertheless, the influence of the precipitation on the SIM formation, reversion process and grain refinement has not been investigated systematically. In this manner it is technically possible to control grain growth during the reversion annealing and to supplement the uniform elongation by precipitation of second phases in NG/UFG stainless steel.
Table 1 Chemical composition (wt%) and Md30 (1C) temperature of Ti-modified 201L stainless steel. C
Mn
Si
Cr
Ni
Ti
P
S
Fe
Md30
0.025
7.2
0.25
16.73
4.3
0.12
0.003
0.011
Balance
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Within the scope of this idea, the objective of this paper is to investigate the influence of Ti addition on the formation of NG/ UFG structure in a Ti-modified 201L stainless steel and its structural and mechanical properties. 2. Experimental procedure In the present work a low Ni, Ti microalloyed 201L ASS was selected. In this grade nickel is partially replaced by Mn. An ingot of Ti-modified 201L ASS was prepared in a vacuum induction furnace under 10 2 mbar pressure. The chemical composition and Md30 temperature of the steel are given in Table 1. The cast ingot was homogenized at 1200 1C for 12 h and then hot rolled at 1100 1C with 85% reduction in thickness. Solution treatment was carried out at a temperature of 1200 1C for different times from 2 to 12 h to dissolve δ-ferrite phase followed by water quenching to prevent carbide precipitation. The solution annealed plate was slightly ground to remove the scale on the surface and then coldrolled to different reductions of 10–90% at room temperature with the strain rate of 0.5 s 1. The reversion annealing was carried out on the 90% cold-rolled specimens by isothermal annealing in the temperature range of 750–900 1C for 60–1800 s. Following the reversion annealing, all specimens were immediately quenched in water. The standard metallographic techniques and subsequently electro-polishing in an electrolytic bath of 200 ml HClO4+800 ml ethanol were employed for the preparation of specimens. To reveal the microstructure, specimens were electrolytically etched in nitric acid (HNO3) solution. The microstructural evolution was investigated using a light microscope “Olympus” and a scanning electron microscope (SEM Philips XL30) equipped with an energy dispersive spectrometer (EDS). For high resolution images, a field emission SEM (Hitachi S4160) was used. The volume fractions of δ-ferrite and strain-induced α′-martensite phases were determined with a ferritescope (Fischer MP30E). The ferritescope readings were converted to α′-martensite volume fraction by multiplying with a correction factor of 1.7 [38]. To avoid errors in determining the volume fraction of SIM phase due to martensite formation during primary preparation, the surfaces of specimens were electropolished before each test. Phase identification was performed by x-ray diffraction (XRD) technique using a Philips X′pert diffractometer. XRD patterns were measured at room temperature in the range of 30–1001 with Cu Kα radiation. In order to determine mechanical properties of the cold rolled and reversion annealed specimens, flat tensile specimens were cut according to ASTM E8 sub-size standard parallel to the rolling direction. Uniaxial tensile tests were conducted at room temperature using a universal tensile machine (Hounsfield H50KS) with a constant cross-head speed of 5 mm/min. Hardness of specimens was measured by the Vickers hardness method (HV) applying a 10-kg force. Ten measurements were obtained from each specimen and an average value was reported.
3. Results and discussion 3.1. Primary observation Fig. 1 shows dendritic microstructure of the as-cast steel. This microstructure shows a duplex structure consisting of dendritic
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Fig. 1. As-cast microstructure showing dendritic structure of austenite and interdenritic δ-ferrite.
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microstructure obtained following the homogenizing at 1200 1C for 12 h. In this stage the grain size of 85 mm was obtained and the volume fraction of δ-ferrite was decreased to about 11%. In order to minimize the fraction of δ-ferrite, the hot-rolled specimens were annealed at 1200 1C. After annealing for 12 h, the volume fraction of δ-ferrite was significantly decreased to less than 2%. However, in contrast to the conventional ASSs without microalloying elements, longer time was required for dissolution of δ in this work. This can be attributed to the presence of titanium. Precipitation of carbides/carbonitrides may alter steel chemistry and subsequently diffusion controlled dissolution kinetics of δ-ferrite. On the other hand, a decrease in the average grain size from 85 mm to 23 mm occurred during annealing for 2 h owing to static recrystallization. However, after annealing for 12 h at 1200 1C, no significant growth was observed in the recrystallized grain structure. This slow grain growth may arise from the pinning effect of the δ-ferrite on the newly recrystallized grain boundaries, and the role of the undissolved Ti containing carbides at grain boundaries that prevent boundary migration. 3.2. Martensite formation
Fig. 2. XRD pattern of as-cast specimen confirming the presence of austenite and ferrite phases in the microstructure.
The XRD patterns of the 5–90% cold-rolled specimens are shown in Fig. 4. The microstructure of the 5% cold-rolled specimen consists of three phases: austenite, α′-martensite and εmartensite. As can be seen, weak ε (1010) and ε (1011) reflections can be detected from the diffraction pattern of the 5% cold-rolled specimen confirming the formation of ε-martensite in the earlier stages of deformation. With increasing strain, the γ and ε reflections disappeared and the intensity of α′ peaks increased. In 35% reduction in thickness, the α′ peaks are observed solely. With increase in the cold reduction to 90%, α′ peaks are broadened owing to an intense increase in crystal defects and very fine structure of martensite. Fig. 5 shows the measured volume fraction of SIM (f α′) as a function of thickness reduction (r) in comparison with 201L [13] and 304L [12] ASSs. The volume fraction of SIM is rapidly increased and becomes saturated at specific strain (εs). According to Fig. 5 this saturation occurs at the cold reduction of 30% (εs ¼ 0.36), 40% (εs ¼0.51) and 70% (εs ¼1.20) for Ti-modified 201L, 201L and 304L steels, respectively. εs plays an important role
Fig. 3. Optical microstructure of the specimen homogenized at 1200 1C for 12 h.
austenite and interdendritic delta ferrite phases. The XRD pattern shown in Fig. 2 confirms the dual phase austenitic–ferritic microstructure. The volume fraction of δ-ferrite in the as-cast specimen measured by the Ferritescope was about 19%. In Fig. 1, the dark area is related to δ-ferrite phase. The δ-ferrite forms during solidification and remains stable in the microstructure, up to room temperature. The δ ferrite is generally assumed to be detrimental due to its effects on the high temperature workability and corrosion resistance [39,40]. Furthermore, this phase does not change to martensite during deformation; hence it can be harmful for thermomechanical processing. Therefore, to improve the quality of the ASS specimens it is preferred to minimize the δ-ferrite phase by a suitable post treatment process. Fig. 3 displays optical
Fig. 4. XRD patterns of the cold-rolled (CR) specimens.
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Fig. 5. Volume fraction of α′-martensite as a function of cold reduction in the present work, 201L [13] and 304L [12] steels.
in attaining NG/UFG structure. The volume fraction of the reversed austenite and amount of its refinement are also dependent upon the εs. Decreasing εs facilitates the process of grain refinement due to provision of a large number of crystal defects such as dislocation tangles, stacking faults, shear bands and mechanical twins in the structure and provision of nucleation sites for austenite during reverse transformation [23,31]. According to Fig. 5, both 201L and Ti-modified 201L steels show lower εs than 304L steel. The 200 series of ASSs contain high manganese instead of nickel as the major austenite forming element and it has been demonstrated that they are more susceptible to martensitic transformation due to lower SFE [41]. On the other hand, the Ti-modified 201L steel displays εs lower than 201L. It may possibly be attributed to the presence of Ti that can affect the stability of austenite. It has been reflected in literature that addition of Ti to the ASS 316 decreases its SFE [42]. Also it is noteworthy that, along with the reducing effect of Ti on the SFE, other effects such as precipitation may affect the austenite stability. For instance, during solution treatment at 1200 1C the primary TiC carbides formed during solidification may dissolve incompletely and remain in the form of fine carbides. The presence of TiC can therefore lead to depletion of carbon in the austenite resulting in a decrease in SFE and consequently enhancing the susceptibility to martensitic transformation. 3.3. Austenite reversion Fig. 6 shows the XRD patterns of the solution annealed, 90% cold-rolled and reversion annealed specimens at various temperatures. As it is observed, the solution treated specimen exhibits nearly single austenite phase. Following the 90% cold reduction, the microstructure of the experimental alloy changes to fully α ′-martensite due to the severe cold rolling. As annealing temperature increases to 800 1C and then 900 1C, austenite peaks appear and increase in intensity, revealing an increase in the volume fraction of γ phase. The martensite disappearing replaced by austenite is an evidence of reversion transformation from SIM to γ-austenite. In addition, the full width at half maximum (FWHM) of the peaks decreases when the annealing temperature increases. This decrease in the FWHM value can be attributed to both reduction of internal stress due to the recovery of crystal defects created during the deformation and grain growth. According to Fig. 6 after annealing at 900 1C for the same duration, the microstructure is fully γ-austenite and the reversion is completed. This result suggests that, a higher temperature can lead to faster reversion due to the diffusion-controlled nature of reversion. Fig. 7 shows the variation of martensite volume fraction measured by the Ferritescope after 90% cold rolling followed by
Fig. 6. XRD patterns of the solution annealed (SA), 90% cold-rolled (CR) and reversion annealed (RA) specimens.
Fig. 7. Variation of martensite volume fraction with annealing time at various temperatures.
reversion annealing at various temperatures vs. annealing time. It is evident from Fig. 7 that the kinetics of the martensite reversion strongly depends on the annealing conditions. Specimens annealed at 750 1C and 800 1C for 30 s show a relatively small austenite phase fraction revealing that these annealing conditions are not sufficient to drive the α′ to γ reversion. Further annealing at 750 1C and 800 1C for 600 s leads to significant α′ to γ reversion. These data indicate that the α′ to γ reversion should be driven by a diffusion-controlled mechanism where an activation energy barrier prevents austenite nucleation at low annealing temperatures and short annealing durations, but higher annealing temperatures or longer annealing durations result in austenite nucleation and growth. Undoubtedly the core of the martensite treatment is the complete reversion transformation of SIM to austenite. As it is seen in Fig. 7, at lower temperatures longer annealing time was taken to complete reversion while the reversion rate is much faster at higher temperatures and it is relatively completed at 850 1C and 900 1C after 60 s because the diffusion rate always accelerates as the temperature is raised above the transformation temperature. According to Fig. 7 maximum reversion is observed at 900 1C while for standard 201L [23] and 301 [10,24] steels this condition has been achieved at 850 1C and 750 1C, respectively. This delayed austenite reversion may be related to the lower stability of austenite in the steel studied here compared to 201L and 301LN steels, which arises from addition of Ti as mentioned before. Another interesting observation shown in Fig. 7 is that the amount
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of martensite is continuously decreased at 750 1C by increasing the annealing time while at higher temperatures the amount of martensite drops to a minimum and then increases with increase in the annealing time. There are two alternative explanations for this phenomenon. The argument that has been reported in several works [25,28,43] is that higher annealing temperatures lead to an increase in the austenite grain size, thus raising the Ms temperature. As a consequence, upon quenching in water thermally induced martensite is more likely to form for specimens annealed at higher temperatures. The steel used in this study has a calculated Ms temperature (based on equation of Eichelman and Hull [44]) of 54 1C for an average grain size above 35 mm. Given that in this study the largest austenite grain size achieved is 4 mm, we should expect an Ms temperature lower than 54 1C. Therefore it is less likely possible for thermally induced martensite to form as a result of austenite grain growth. The second mechanism takes into consideration carbide precipitation that leads to carbon depletion from the matrix and consequently increasing of Ms temperature up to room temperature. This phenomenon takes place when, upon annealing the martensite starts to release carbon, leading to formation of carbides [22]. Calculation of Ms temperature demonstrates that for the case where the austenite matrix is without carbon, the Ms is raised to 95 1C. Since Ti is a strong carbide former, it is not unexpected for TiC precipitation to occur during the reversion annealing. Carbide particles were detected in the annealed specimens. Fig. 8 shows the carbide particles of about 70 nm in the specimen annealed at 900 1C for 1800 s. Annealing for a long time leads to a significant increase in the size of carbides and consequently simplified precipitates observation. A comparison of the EDS spectra from the particles and matrix shows an increase in C and Ti content in the second phase particles (Fig. 8b and c). The Au peak in the EDS spectra is from the Au coat applied on the surface of SEM specimen. The results confirm that the precipitates may be TiC that cannot be characterized in SEM images in short annealing times due to very small size. SEM micrographs of specimens subjected to 90% cold rolling and annealed at 850 1C and 900 1C for different annealing times of 60 s, 300 s and 600 s are illustrated in Fig. 9. The α′ to γ reversion is nearly completed at these temperatures and the microstructure of specimens consists of the reverted equiaxed austenite grains. The average austenite grain size was measured as 53 nm and 45 nm in the specimens annealed at 850 1C and 900 1C for 60 s, respectively. Holding at 850 1C and 900 1C for slightly longer duration of 600 s shows some degree of grain growth. Smaller grain size in the 900 1C could be a consequence of carbide formation hindering grain growth. The variations of reverted austenite grain size with the annealing time at different temperatures are depicted in Fig. 10. The very small increase in grain size at both 850 1C and 900 1C for duration up to 600 s seems to indicate that the presence of TiC carbides, which grow and coarsen during annealing, plays an important role in suppressing grain growth. Whereas in the standard 201L stainless steel [23] an increase in the grain size from 90 nm to 2300 nm has been reported for the annealing times of 30–600 s, in this work for the same duration, the grain size reaches only to 420 nm. Therefore it can be concluded that grain growth is restricted in 900 1C due to the presence of TiC precipitates in the austenite matrix and the grains remain within the nanometer range up to 600 s. 3.4. Mechanical properties The engineering stress–strain curves of the 90% cold rolled and annealed specimens at 900 1C are shown together with that of the initial solution annealed specimen in Fig. 11. The strength and elongation of the 90% cold rolled specimen are about 1900 MPa
Fig. 8. SEM image of carbides (a), with EDS analysis of the austenitic matrix (b), and corresponding carbides in the specimen annealed at 900 1C for 1800 s.
and 2%, respectively. However upon the reversion and formation of NG structure, the strength is decreased and ductility is increased sharply. The NG specimen possesses a very high yield strength
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Fig. 9. SEM micrographs of specimens reversion annealed at 850 1C for (a) 60 s, (b) 300 s and (c) 600 s and 900 1C for (d) 60 s, (e) 300 s and (f) 600 s.
Fig. 10. The effect of annealing time on the grain size of reversed austenite at 850 1C and 900 1C.
( 1000 MPa), six times higher than that of the initial coarsegrained steel and a very high ultimate tensile strength ( 1330 MPa). This improvement is also accompanied by 42% total elongation for the NG specimen. For vast majority of the NG/UFG metals, ductility is nowhere close to their characteristic ductility in the coarse-grained form. However, the present results show that a significant grain refinement has taken place by the thermomechanical treatment (from 35 mm to 45 nm), and the resulting NG/UFG Ti-modified 201L stainless steel possesses high ductility. This further improvement in ductility is believed to be attributed to the precipitation of TiC. The fine precipitates are deemed to improve mechanical properties through promotion of SIM transformation due to the depletion of solute atoms near the precipitates. More work is in progress to investigate the mechanism of such behavior. Mechanical properties upon reversion annealing along with microstructural features are collected in Table 2. As indicated in Table 2, specimens annealed at low temperatures (750 1C and 800 1C) for short durations exhibited high strength but low elongation values. This behavior resulted from the high volume fraction of untransformed martensite still present. With increase in the annealing time, the strength decreased and elongation increased considerably. Upon annealing at higher temperatures (850 1C and 900 1C), significant change in tensile properties was
Fig. 11. Engineering stress–strain curves of (A) solution annealed specimen (grain size: 35 mm), (B) 90% cold-rolled specimen, and (C) reversion annealed specimen at 900 1C for 60 s (grain size: 45 nm).
observed only after 60 s. For short annealing times at 850 1C and 900 1C, fine grain size and approximately full austenitic microstructure lead to an excellent combination of strength and elongation, while a reduction in yield strength was observed with increase in the annealing time up to 600 s due to relative grain growth during the annealing, despite the increase in martensite volume fraction. This seems to indicate that in terms of the yield strength, the increase in the volume fraction of martensite for the samples annealed at 850 1C and 900 1C for 60–600 s is compensated by the grain growth. The significant effect of austenitic grain size on the strength values is evident when comparing the samples annealed at 850 1C for 300 s and 900 1C for 60 s (see Table 2). In this comparison, although the martensite content is higher in the 850 1C–300 s sample (12% vs. 4%), the yield strength is lower (867 MPa vs. 975 MPa) because of larger grain size (79 nm vs. 45 nm). Once the nucleation of the austenite grain takes place, its growth rate as a function of time and temperature determines the final grain size. Therefore, to attain a small average grain size, it is necessary to accelerate the nucleation (reversion) rate and retard the grain growth. At the lower temperatures and times, 750 1C and 800 1C, the structure is still in the initial phase of the
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Table 2 The reverted austenite volume percentage, grain size and mechanical properties at different annealing conditions. Annealing temperature (1C)
Annealing time (s)
Austenite vol%
Grain Yield size (nm) strength (MPa)
Total elongation (%)
750 750 750 800 800 800 850 850 850 850 900 900 900 900
120 300 600 120 300 600 60 120 300 600 60 120 300 600
86 91 94 90 94 92 93 92 88 82 96 94 92 89
– – 1257 30 – – 1747 28 537 12 627 14 797 19 2177 73 457 12 567 11 937 20 4217 71
– – 34 – 37 – 36 38 37 39 42 41 45 38
1176 967 856 1025 780 829 1005 884 867 830 975 870 805 746
reversion, which is confirmed by the presence of significant fraction of martensite. While the required time to complete the phase reversion at lower temperatures is longer than desired, the rate of transformation is quite fast at higher temperatures. As mentioned before, a higher annealing temperature provides more thermal energy to overcome the activation energy for nucleation. On the other hand, short annealing times prevent the reverted austenite from entering into the grain growth stage. In fact, in competition between rate of transformation and rate of grain growth at a high temperature and for a short time, the former will be dominant. Hence, annealing at 900 1C for a short time (60 s) appears to be more favorable for fine grain formation and achieving a good combination of high yield strength ( 1000 MPa) and high total elongation ( 40%). 4. Conclusions In this work the effect of Ti addition on the formation of nanocrystalline structure in the 201L ASS was studied. From this investigation the following conclusions can be drawn: 1. Under similar thermomechanical processing, the saturation strain for the formation of α′-martensite during cold rolling of Ti-modified 201L stainless steel was found at a lower rolling reduction than that of the standard 201L ASS, showing a promoting effect of Ti on the formation of SIM. 2. The highest volume fraction of the reverted austenite ( 96%) was achieved after annealing at 900 1C for 60 s in a diffusioncontrolled reversion process. At the lower annealing temperatures, the microstructure contained retained martensite, whereas for long annealing times, the thermally-induced martensite was formed during water quenching. 3. The smallest grain size of 457 12 nm was achieved in the specimen annealed at 900 1C for 60 s. It was found that the presence of TiC precipitates may retard the reversion process and suppress the grain growth in further annealing times.
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4. The nano-grained Ti-modified 201L stainless steel with the average grain size of 45 nm exhibited a good combination of high yield strength of about 1000 MPa (six times higher than that of the initial coarse-grained steel), ultimate strength of 1330 MPa and high elongation to failure of 42%.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Influence of Ti microalloying on the formation of nanocrystalline structure in the 201L austenitic stainless steel during martensite thermomechanical treatment S. Sadeghpour a,n, A. Kermanpur a, A. Najafizadeh b a b
Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran Fould Institute of Technology, Fouldshar, Isfahan, Iran
art ic l e i nf o
a b s t r a c t
Article history: Received 26 April 2013 Received in revised form 27 June 2013 Accepted 11 July 2013 Available online 19 July 2013
The martensite thermomechanical treatment was used for the formation of nano/ultrafine grain structure in the 201L austenitic stainless steel containing 0.12 wt% Ti microalloying element. The initial microstructure was provided through homogenizing, hot rolling and solution annealing of the as-cast ingots. The specimens were then cold rolled between 5% and 90% thickness reduction and subsequently annealed at 750–900 1C for various times. The results showed a promoting effect of Ti on the formation of strain-induced martensite (SIM). A nanocrystalline austenitic structure with average grain size of 45 nm was achieved by annealing at 900 1C for 60 s through a diffusional transformation mechanism. It was found that precipitation of nanosized TiC particles during the reversion annealing could retard the reversion process and suppressed grain growth in further annealing times. The tensile testing of the thermomechanically processed specimens showed a good combination of high yield strength ( 1000 MPa, six times higher than that of the initial coarse-grained steel) and excellent ductility (42% total elongation) for the Ti microalloyed 201L steel due to the SIM formation during the deformation and impressions of the nanosized Ti carbides distributed within the nano/ultrafine grain structure. Published by Elsevier B.V.
Keywords: 201L stainless steel Ti microalloying Nano/ultrafine grained structure Martensite thermomechanical treatment Strain-induced martensite Reversion annealing
1. Introduction Extensive research has been carried out on grain refinement of metals and alloys down to the nano-grain (NG, d o100 nm) and ultrafine-grain (UFG, do 500 nm) scales in order to improve their mechanical properties. Whereas significant increases in hardness and strength have been reported for nanostructured materials, they often exhibit a very low tensile elongation to failure especially uniform elongation (strain before necking) [1,2]. The reason for the low ductility in NG/UFG metals is plastic instability due to the limited strain hardening capacity that arises from the fact that extremely small grains cannot store dislocations to increase their density by order of magnitude as normally possible in the coarsegrained (CG) metals [3]. It has confirmed that any improvement of the strain hardening ability as a stabilizing mechanism to bring the plastic instability under control will be beneficial to enhancing the homogeneous plastic deformation for NG and UFG materials [4]. Several approaches have been developed in terms of microstructural design for improving ductility of the NG and UFG metals
n
Corresponding author. Tel.: +98 0936 9159 669; fax: +98 0311 3912 753. E-mail addresses: [email protected], [email protected] (S. Sadeghpour). 0921-5093/$ - see front matter Published by Elsevier B.V. http://dx.doi.org/10.1016/j.msea.2013.07.014
without sacrificing their strength. These include dispersion of nanoparticles and nano-precipitates in the nanostructured metal matrix [5], to use transformation induced plasticity (TRIP) [6] and twinning induced plasticity (TWIP) [7] effects, enhancing the fraction of high angle grain boundaries and lowering the density of dislocations [8], and improving strain hardening rate by tailoring the stacking fault energy (SFE) [9]. A recent thermomechanical route of the formation of NG/UFG structures in metastable austenitic steels called martensite treatment is characterized by strain induced martensitic transformation and its reversion to austenite [10–13]. In this approach, severe deformation of austenite at room temperature leads to strain induced transformation of austenite (γ) to martensite (α′) and upon annealing, this heavily deformed strain induced martensite (SIM) transforms back to austenite leading to a noticeable grain refinement. Because of low SFE, austenitic stainless steels (ASSs) are susceptible to deformation induced martensitic transformation. When ASSs are deformed, strain induced ε-martensite with the hexagonal close packed (hcp) structure and α′-martensite with body centered cubic (bcc) structure are formed from the metastable austenite. The formation of SIM is related closely to shear bands which are planar defects associated with overlapping of stacking faults on {111}γ. Shear band intersections act as the nucleation sites for
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α′-martensite, and the formation of the shear bands is a precursor for the strain induced α′-martensite transformation [14]. Several factors including grain size, strain rate, deformation temperature, mode of deformation, etc. may influence the martensitic transformation [15]. In addition to these factors, steel chemistry will also bear heavily on martensitic transformation. This has been explained in terms of the variation in the chemical free energy difference between austenite and α′-martensite phase as the chemical driving force. But Talonen and Hanninen [14] have suggested that the stability of austenite against martensite transformation is influenced more by SFE than the chemical driving force. It is well known that the SFE is strongly dependent on the chemical composition [16,17] and temperature [18,19]. Many authors have investigated the influence of temperature on the formation of SIM in the austenitic steels [14,20]. The effect of composition on the formation of SIM was first investigated by Angel [21], who correlated elemental compositions with the temperature at which an amount of 50% austenite will be transformed to martensite by the application of 0.30 true strain ( 35% engineering strain), denoted by Md30. Nohara et al. [22] modified Angel′s equation to include additional elements and grain size as follows (in wt%): M d30 ¼ 551462 %C9:2 %Si8:1 %Mn 13:7 %Cr29 %Ni18:5 %Mo29 %Cu 68 %Nb462 %N1:42 ðGS 8:0Þ where GS is the ASTM grain size number. A low Md30 temperature indicates high austenite stability against the formation of SIM. Moreover according to Talonen et al. [14], the temporal and compositional dependence values of the austenite stability are accounted by variation in the SFE. As mentioned before, the severe cold deformed SIM transforms back to austenite upon annealing. The α′ to γ reversion transformation has been studied by a number of authors [23–27]. This reversion may proceed either through shear or diffusion-controlled mechanisms. The alloy chemical composition plays a key role in determining the α′ to γ reversion mechanism. Takaki et al. [28] reported that the reverse transformation of martensite to austenite occurs by diffusionless mechanism at the low ratio of Cr to Ni, whereas diffusive reverse transformation happens at the high ratio of Cr to Ni in metastable Fe– Cr–Ni ternary alloys. In steels which undergo martensitic shear reversion, reversion completes during heating and reversed austenite contains a high density of dislocations just after reversion [29]. During successive annealing, grain refining proceeds through the process of recrystallization of reverse austenite. On the other hand, diffusional reversion occurs primarily by nucleation and growth of fine austenite grains at martensite lath boundaries [28]. Diffusion-type mechanism is characterized by (i) a wide annealing temperature range; (ii) the formation of defect free equiaxed austenitic grains which grow in size with time; (iii) a wide austenite grain size distribution; and (iv) possible formation of secondary phase precipitates [30]. In recent years the martensite thermomechanical process using SIM and its reverse transformation has been extensively investigated to produce NG/UFG structure in the conventional ASSs especially 300 series [10–12,31–33]. The martensite treatment has been modified to obtain finer grain sizes and better mechanical properties with combinations of recrystallization, warm deformation, and repetition of the process [34–36]. In addition Lee et al. [37] reported that the nanosized Nb particles increased the strength without reducing ductility of the UFG steel processed by martensite treatment. Nevertheless, the influence of the precipitation on the SIM formation, reversion process and grain refinement has not been investigated systematically. In this manner it is technically possible to control grain growth during the reversion annealing and to supplement the uniform elongation by precipitation of second phases in NG/UFG stainless steel.
Table 1 Chemical composition (wt%) and Md30 (1C) temperature of Ti-modified 201L stainless steel. C
Mn
Si
Cr
Ni
Ti
P
S
Fe
Md30
0.025
7.2
0.25
16.73
4.3
0.12
0.003
0.011
Balance
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Within the scope of this idea, the objective of this paper is to investigate the influence of Ti addition on the formation of NG/ UFG structure in a Ti-modified 201L stainless steel and its structural and mechanical properties. 2. Experimental procedure In the present work a low Ni, Ti microalloyed 201L ASS was selected. In this grade nickel is partially replaced by Mn. An ingot of Ti-modified 201L ASS was prepared in a vacuum induction furnace under 10 2 mbar pressure. The chemical composition and Md30 temperature of the steel are given in Table 1. The cast ingot was homogenized at 1200 1C for 12 h and then hot rolled at 1100 1C with 85% reduction in thickness. Solution treatment was carried out at a temperature of 1200 1C for different times from 2 to 12 h to dissolve δ-ferrite phase followed by water quenching to prevent carbide precipitation. The solution annealed plate was slightly ground to remove the scale on the surface and then coldrolled to different reductions of 10–90% at room temperature with the strain rate of 0.5 s 1. The reversion annealing was carried out on the 90% cold-rolled specimens by isothermal annealing in the temperature range of 750–900 1C for 60–1800 s. Following the reversion annealing, all specimens were immediately quenched in water. The standard metallographic techniques and subsequently electro-polishing in an electrolytic bath of 200 ml HClO4+800 ml ethanol were employed for the preparation of specimens. To reveal the microstructure, specimens were electrolytically etched in nitric acid (HNO3) solution. The microstructural evolution was investigated using a light microscope “Olympus” and a scanning electron microscope (SEM Philips XL30) equipped with an energy dispersive spectrometer (EDS). For high resolution images, a field emission SEM (Hitachi S4160) was used. The volume fractions of δ-ferrite and strain-induced α′-martensite phases were determined with a ferritescope (Fischer MP30E). The ferritescope readings were converted to α′-martensite volume fraction by multiplying with a correction factor of 1.7 [38]. To avoid errors in determining the volume fraction of SIM phase due to martensite formation during primary preparation, the surfaces of specimens were electropolished before each test. Phase identification was performed by x-ray diffraction (XRD) technique using a Philips X′pert diffractometer. XRD patterns were measured at room temperature in the range of 30–1001 with Cu Kα radiation. In order to determine mechanical properties of the cold rolled and reversion annealed specimens, flat tensile specimens were cut according to ASTM E8 sub-size standard parallel to the rolling direction. Uniaxial tensile tests were conducted at room temperature using a universal tensile machine (Hounsfield H50KS) with a constant cross-head speed of 5 mm/min. Hardness of specimens was measured by the Vickers hardness method (HV) applying a 10-kg force. Ten measurements were obtained from each specimen and an average value was reported.
3. Results and discussion 3.1. Primary observation Fig. 1 shows dendritic microstructure of the as-cast steel. This microstructure shows a duplex structure consisting of dendritic
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Fig. 1. As-cast microstructure showing dendritic structure of austenite and interdenritic δ-ferrite.
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microstructure obtained following the homogenizing at 1200 1C for 12 h. In this stage the grain size of 85 mm was obtained and the volume fraction of δ-ferrite was decreased to about 11%. In order to minimize the fraction of δ-ferrite, the hot-rolled specimens were annealed at 1200 1C. After annealing for 12 h, the volume fraction of δ-ferrite was significantly decreased to less than 2%. However, in contrast to the conventional ASSs without microalloying elements, longer time was required for dissolution of δ in this work. This can be attributed to the presence of titanium. Precipitation of carbides/carbonitrides may alter steel chemistry and subsequently diffusion controlled dissolution kinetics of δ-ferrite. On the other hand, a decrease in the average grain size from 85 mm to 23 mm occurred during annealing for 2 h owing to static recrystallization. However, after annealing for 12 h at 1200 1C, no significant growth was observed in the recrystallized grain structure. This slow grain growth may arise from the pinning effect of the δ-ferrite on the newly recrystallized grain boundaries, and the role of the undissolved Ti containing carbides at grain boundaries that prevent boundary migration. 3.2. Martensite formation
Fig. 2. XRD pattern of as-cast specimen confirming the presence of austenite and ferrite phases in the microstructure.
The XRD patterns of the 5–90% cold-rolled specimens are shown in Fig. 4. The microstructure of the 5% cold-rolled specimen consists of three phases: austenite, α′-martensite and εmartensite. As can be seen, weak ε (1010) and ε (1011) reflections can be detected from the diffraction pattern of the 5% cold-rolled specimen confirming the formation of ε-martensite in the earlier stages of deformation. With increasing strain, the γ and ε reflections disappeared and the intensity of α′ peaks increased. In 35% reduction in thickness, the α′ peaks are observed solely. With increase in the cold reduction to 90%, α′ peaks are broadened owing to an intense increase in crystal defects and very fine structure of martensite. Fig. 5 shows the measured volume fraction of SIM (f α′) as a function of thickness reduction (r) in comparison with 201L [13] and 304L [12] ASSs. The volume fraction of SIM is rapidly increased and becomes saturated at specific strain (εs). According to Fig. 5 this saturation occurs at the cold reduction of 30% (εs ¼ 0.36), 40% (εs ¼0.51) and 70% (εs ¼1.20) for Ti-modified 201L, 201L and 304L steels, respectively. εs plays an important role
Fig. 3. Optical microstructure of the specimen homogenized at 1200 1C for 12 h.
austenite and interdendritic delta ferrite phases. The XRD pattern shown in Fig. 2 confirms the dual phase austenitic–ferritic microstructure. The volume fraction of δ-ferrite in the as-cast specimen measured by the Ferritescope was about 19%. In Fig. 1, the dark area is related to δ-ferrite phase. The δ-ferrite forms during solidification and remains stable in the microstructure, up to room temperature. The δ ferrite is generally assumed to be detrimental due to its effects on the high temperature workability and corrosion resistance [39,40]. Furthermore, this phase does not change to martensite during deformation; hence it can be harmful for thermomechanical processing. Therefore, to improve the quality of the ASS specimens it is preferred to minimize the δ-ferrite phase by a suitable post treatment process. Fig. 3 displays optical
Fig. 4. XRD patterns of the cold-rolled (CR) specimens.
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Fig. 5. Volume fraction of α′-martensite as a function of cold reduction in the present work, 201L [13] and 304L [12] steels.
in attaining NG/UFG structure. The volume fraction of the reversed austenite and amount of its refinement are also dependent upon the εs. Decreasing εs facilitates the process of grain refinement due to provision of a large number of crystal defects such as dislocation tangles, stacking faults, shear bands and mechanical twins in the structure and provision of nucleation sites for austenite during reverse transformation [23,31]. According to Fig. 5, both 201L and Ti-modified 201L steels show lower εs than 304L steel. The 200 series of ASSs contain high manganese instead of nickel as the major austenite forming element and it has been demonstrated that they are more susceptible to martensitic transformation due to lower SFE [41]. On the other hand, the Ti-modified 201L steel displays εs lower than 201L. It may possibly be attributed to the presence of Ti that can affect the stability of austenite. It has been reflected in literature that addition of Ti to the ASS 316 decreases its SFE [42]. Also it is noteworthy that, along with the reducing effect of Ti on the SFE, other effects such as precipitation may affect the austenite stability. For instance, during solution treatment at 1200 1C the primary TiC carbides formed during solidification may dissolve incompletely and remain in the form of fine carbides. The presence of TiC can therefore lead to depletion of carbon in the austenite resulting in a decrease in SFE and consequently enhancing the susceptibility to martensitic transformation. 3.3. Austenite reversion Fig. 6 shows the XRD patterns of the solution annealed, 90% cold-rolled and reversion annealed specimens at various temperatures. As it is observed, the solution treated specimen exhibits nearly single austenite phase. Following the 90% cold reduction, the microstructure of the experimental alloy changes to fully α ′-martensite due to the severe cold rolling. As annealing temperature increases to 800 1C and then 900 1C, austenite peaks appear and increase in intensity, revealing an increase in the volume fraction of γ phase. The martensite disappearing replaced by austenite is an evidence of reversion transformation from SIM to γ-austenite. In addition, the full width at half maximum (FWHM) of the peaks decreases when the annealing temperature increases. This decrease in the FWHM value can be attributed to both reduction of internal stress due to the recovery of crystal defects created during the deformation and grain growth. According to Fig. 6 after annealing at 900 1C for the same duration, the microstructure is fully γ-austenite and the reversion is completed. This result suggests that, a higher temperature can lead to faster reversion due to the diffusion-controlled nature of reversion. Fig. 7 shows the variation of martensite volume fraction measured by the Ferritescope after 90% cold rolling followed by
Fig. 6. XRD patterns of the solution annealed (SA), 90% cold-rolled (CR) and reversion annealed (RA) specimens.
Fig. 7. Variation of martensite volume fraction with annealing time at various temperatures.
reversion annealing at various temperatures vs. annealing time. It is evident from Fig. 7 that the kinetics of the martensite reversion strongly depends on the annealing conditions. Specimens annealed at 750 1C and 800 1C for 30 s show a relatively small austenite phase fraction revealing that these annealing conditions are not sufficient to drive the α′ to γ reversion. Further annealing at 750 1C and 800 1C for 600 s leads to significant α′ to γ reversion. These data indicate that the α′ to γ reversion should be driven by a diffusion-controlled mechanism where an activation energy barrier prevents austenite nucleation at low annealing temperatures and short annealing durations, but higher annealing temperatures or longer annealing durations result in austenite nucleation and growth. Undoubtedly the core of the martensite treatment is the complete reversion transformation of SIM to austenite. As it is seen in Fig. 7, at lower temperatures longer annealing time was taken to complete reversion while the reversion rate is much faster at higher temperatures and it is relatively completed at 850 1C and 900 1C after 60 s because the diffusion rate always accelerates as the temperature is raised above the transformation temperature. According to Fig. 7 maximum reversion is observed at 900 1C while for standard 201L [23] and 301 [10,24] steels this condition has been achieved at 850 1C and 750 1C, respectively. This delayed austenite reversion may be related to the lower stability of austenite in the steel studied here compared to 201L and 301LN steels, which arises from addition of Ti as mentioned before. Another interesting observation shown in Fig. 7 is that the amount
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of martensite is continuously decreased at 750 1C by increasing the annealing time while at higher temperatures the amount of martensite drops to a minimum and then increases with increase in the annealing time. There are two alternative explanations for this phenomenon. The argument that has been reported in several works [25,28,43] is that higher annealing temperatures lead to an increase in the austenite grain size, thus raising the Ms temperature. As a consequence, upon quenching in water thermally induced martensite is more likely to form for specimens annealed at higher temperatures. The steel used in this study has a calculated Ms temperature (based on equation of Eichelman and Hull [44]) of 54 1C for an average grain size above 35 mm. Given that in this study the largest austenite grain size achieved is 4 mm, we should expect an Ms temperature lower than 54 1C. Therefore it is less likely possible for thermally induced martensite to form as a result of austenite grain growth. The second mechanism takes into consideration carbide precipitation that leads to carbon depletion from the matrix and consequently increasing of Ms temperature up to room temperature. This phenomenon takes place when, upon annealing the martensite starts to release carbon, leading to formation of carbides [22]. Calculation of Ms temperature demonstrates that for the case where the austenite matrix is without carbon, the Ms is raised to 95 1C. Since Ti is a strong carbide former, it is not unexpected for TiC precipitation to occur during the reversion annealing. Carbide particles were detected in the annealed specimens. Fig. 8 shows the carbide particles of about 70 nm in the specimen annealed at 900 1C for 1800 s. Annealing for a long time leads to a significant increase in the size of carbides and consequently simplified precipitates observation. A comparison of the EDS spectra from the particles and matrix shows an increase in C and Ti content in the second phase particles (Fig. 8b and c). The Au peak in the EDS spectra is from the Au coat applied on the surface of SEM specimen. The results confirm that the precipitates may be TiC that cannot be characterized in SEM images in short annealing times due to very small size. SEM micrographs of specimens subjected to 90% cold rolling and annealed at 850 1C and 900 1C for different annealing times of 60 s, 300 s and 600 s are illustrated in Fig. 9. The α′ to γ reversion is nearly completed at these temperatures and the microstructure of specimens consists of the reverted equiaxed austenite grains. The average austenite grain size was measured as 53 nm and 45 nm in the specimens annealed at 850 1C and 900 1C for 60 s, respectively. Holding at 850 1C and 900 1C for slightly longer duration of 600 s shows some degree of grain growth. Smaller grain size in the 900 1C could be a consequence of carbide formation hindering grain growth. The variations of reverted austenite grain size with the annealing time at different temperatures are depicted in Fig. 10. The very small increase in grain size at both 850 1C and 900 1C for duration up to 600 s seems to indicate that the presence of TiC carbides, which grow and coarsen during annealing, plays an important role in suppressing grain growth. Whereas in the standard 201L stainless steel [23] an increase in the grain size from 90 nm to 2300 nm has been reported for the annealing times of 30–600 s, in this work for the same duration, the grain size reaches only to 420 nm. Therefore it can be concluded that grain growth is restricted in 900 1C due to the presence of TiC precipitates in the austenite matrix and the grains remain within the nanometer range up to 600 s. 3.4. Mechanical properties The engineering stress–strain curves of the 90% cold rolled and annealed specimens at 900 1C are shown together with that of the initial solution annealed specimen in Fig. 11. The strength and elongation of the 90% cold rolled specimen are about 1900 MPa
Fig. 8. SEM image of carbides (a), with EDS analysis of the austenitic matrix (b), and corresponding carbides in the specimen annealed at 900 1C for 1800 s.
and 2%, respectively. However upon the reversion and formation of NG structure, the strength is decreased and ductility is increased sharply. The NG specimen possesses a very high yield strength
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Fig. 9. SEM micrographs of specimens reversion annealed at 850 1C for (a) 60 s, (b) 300 s and (c) 600 s and 900 1C for (d) 60 s, (e) 300 s and (f) 600 s.
Fig. 10. The effect of annealing time on the grain size of reversed austenite at 850 1C and 900 1C.
( 1000 MPa), six times higher than that of the initial coarsegrained steel and a very high ultimate tensile strength ( 1330 MPa). This improvement is also accompanied by 42% total elongation for the NG specimen. For vast majority of the NG/UFG metals, ductility is nowhere close to their characteristic ductility in the coarse-grained form. However, the present results show that a significant grain refinement has taken place by the thermomechanical treatment (from 35 mm to 45 nm), and the resulting NG/UFG Ti-modified 201L stainless steel possesses high ductility. This further improvement in ductility is believed to be attributed to the precipitation of TiC. The fine precipitates are deemed to improve mechanical properties through promotion of SIM transformation due to the depletion of solute atoms near the precipitates. More work is in progress to investigate the mechanism of such behavior. Mechanical properties upon reversion annealing along with microstructural features are collected in Table 2. As indicated in Table 2, specimens annealed at low temperatures (750 1C and 800 1C) for short durations exhibited high strength but low elongation values. This behavior resulted from the high volume fraction of untransformed martensite still present. With increase in the annealing time, the strength decreased and elongation increased considerably. Upon annealing at higher temperatures (850 1C and 900 1C), significant change in tensile properties was
Fig. 11. Engineering stress–strain curves of (A) solution annealed specimen (grain size: 35 mm), (B) 90% cold-rolled specimen, and (C) reversion annealed specimen at 900 1C for 60 s (grain size: 45 nm).
observed only after 60 s. For short annealing times at 850 1C and 900 1C, fine grain size and approximately full austenitic microstructure lead to an excellent combination of strength and elongation, while a reduction in yield strength was observed with increase in the annealing time up to 600 s due to relative grain growth during the annealing, despite the increase in martensite volume fraction. This seems to indicate that in terms of the yield strength, the increase in the volume fraction of martensite for the samples annealed at 850 1C and 900 1C for 60–600 s is compensated by the grain growth. The significant effect of austenitic grain size on the strength values is evident when comparing the samples annealed at 850 1C for 300 s and 900 1C for 60 s (see Table 2). In this comparison, although the martensite content is higher in the 850 1C–300 s sample (12% vs. 4%), the yield strength is lower (867 MPa vs. 975 MPa) because of larger grain size (79 nm vs. 45 nm). Once the nucleation of the austenite grain takes place, its growth rate as a function of time and temperature determines the final grain size. Therefore, to attain a small average grain size, it is necessary to accelerate the nucleation (reversion) rate and retard the grain growth. At the lower temperatures and times, 750 1C and 800 1C, the structure is still in the initial phase of the
S. Sadeghpour et al. / Materials Science & Engineering A 584 (2013) 177–183
Table 2 The reverted austenite volume percentage, grain size and mechanical properties at different annealing conditions. Annealing temperature (1C)
Annealing time (s)
Austenite vol%
Grain Yield size (nm) strength (MPa)
Total elongation (%)
750 750 750 800 800 800 850 850 850 850 900 900 900 900
120 300 600 120 300 600 60 120 300 600 60 120 300 600
86 91 94 90 94 92 93 92 88 82 96 94 92 89
– – 1257 30 – – 1747 28 537 12 627 14 797 19 2177 73 457 12 567 11 937 20 4217 71
– – 34 – 37 – 36 38 37 39 42 41 45 38
1176 967 856 1025 780 829 1005 884 867 830 975 870 805 746
reversion, which is confirmed by the presence of significant fraction of martensite. While the required time to complete the phase reversion at lower temperatures is longer than desired, the rate of transformation is quite fast at higher temperatures. As mentioned before, a higher annealing temperature provides more thermal energy to overcome the activation energy for nucleation. On the other hand, short annealing times prevent the reverted austenite from entering into the grain growth stage. In fact, in competition between rate of transformation and rate of grain growth at a high temperature and for a short time, the former will be dominant. Hence, annealing at 900 1C for a short time (60 s) appears to be more favorable for fine grain formation and achieving a good combination of high yield strength ( 1000 MPa) and high total elongation ( 40%). 4. Conclusions In this work the effect of Ti addition on the formation of nanocrystalline structure in the 201L ASS was studied. From this investigation the following conclusions can be drawn: 1. Under similar thermomechanical processing, the saturation strain for the formation of α′-martensite during cold rolling of Ti-modified 201L stainless steel was found at a lower rolling reduction than that of the standard 201L ASS, showing a promoting effect of Ti on the formation of SIM. 2. The highest volume fraction of the reverted austenite ( 96%) was achieved after annealing at 900 1C for 60 s in a diffusioncontrolled reversion process. At the lower annealing temperatures, the microstructure contained retained martensite, whereas for long annealing times, the thermally-induced martensite was formed during water quenching. 3. The smallest grain size of 457 12 nm was achieved in the specimen annealed at 900 1C for 60 s. It was found that the presence of TiC precipitates may retard the reversion process and suppress the grain growth in further annealing times.
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4. The nano-grained Ti-modified 201L stainless steel with the average grain size of 45 nm exhibited a good combination of high yield strength of about 1000 MPa (six times higher than that of the initial coarse-grained steel), ultimate strength of 1330 MPa and high elongation to failure of 42%.
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