Deformation mechanisms in nanostructured bainitic steels under torsion

Journal Pre-proof Deformation mechanisms in nanostructured bainitic steels under torsion Avanish Kumar, Aparna Singh PII:

S0921-5093(19)31314-0

DOI:

https://doi.org/10.1016/j.msea.2019.138528

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MSA 138528

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Materials Science & Engineering A

Received Date: 29 August 2019 Revised Date:

6 October 2019

Accepted Date: 8 October 2019

Please cite this article as: A. Kumar, A. Singh, Deformation mechanisms in nanostructured bainitic steels under torsion, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/ j.msea.2019.138528. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Deformation mechanisms in nanostructured bainitic steels under torsion Avanish Kumar, Aparna Singh* Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai, 400076, India

Abstract: A clean high carbon steel was made using vacuum induction melting and three distinct nanostructures were obtained by austempering it at 250, 300 and 350°C respectively after a similar austenitization treatment at 950°C. A decrease in austempering temperature causes refinement of bainitic ferrite (BF) as well as retained austenite (RA). Moreover, the volume fraction of BF increases with a reduction in austempering temperature. The mechanical properties under tension and torsion at ambient temperature were investigated. The results demonstrated that the slenderness of BF plates as well as their enhanced volume fraction plays a major role in enhancing the tensile as well as shear strength. This came with a concomitant loss in ductility of the specimens austempered at lower temperatures. However, the hardness measurements on the gauge section were able to highlight the important differences in the deformation behavior between torsion and uniaxial tension. It was found that hardness increase of the specimen after torsion was much higher than uniaxial tension. The ductility is controlled by the content and stability of RA, BF as well as the newly formed martensite. X-ray diffraction studies showed that more RA was transformed under torsion than uniaxial tension. Steel specimens with the highest content of RA showed the highest ductility under torsion while the specimens with the highest BF showed the highest strength. Keywords: Steel, bainite, torsion, deformation, phase transformation, TRIP *

Corresponding author, E-mail: [email protected], Telephone: +91-22-2576-7605

1. Introduction There have been many studies on the development of high carbon carbide free bainitic steels that can be heat treated to be nanostructured due to their low bainitestart temperatures [1–4]. Silicon (~ 1.5 wt.%) is deliberately made part of the composition to retard the precipitation of cementite [5]. Fine slender plates of BF separated by films of RA provide an excellent combination of tensile strength and ductility to these steels [6,7]. The strength (> 2 GPa) is primarily attributed to the high density of interfaces and dislocation density. The carbon-enriched RA possesses the potential for transformation induced plasticity (TRIP effect) with the coarse austenite having a lower stability than the films [4,6,8–11]. When an external load is applied, coarse austenite may easily transform to martensite. Since martensite is hard and brittle, it deteriorates the toughness and ductility of these steels [7,12]. Most studies have investigated the tensile strength, ductility, fracture toughness and fatigue of nanostructured bainitic steels [8,9,21,22,13–20]. However, the deformation characteristics of these steels under torsion have not received enough attention. Many real-life applications such as rails, gear boxes etc., have shear loading driving the deformation. Malinov et al. [23] investigated the effect of RA on the torsional behavior of quenched and tempered steels. The torque versus angle of twist diagram showed jumps, which were dependent on the RA stability; a higher stability led to smoother curves. Hohenwarter et al. [24] studied severe plastic deformation of a carbide free bainitic steel through high pressure torsion. An alignment of the fragmented structure in the shear direction was observed. RA content decreased drastically after a small amount of strain and strong hardening behaviour was observed only at the lower strains. In the present work, we have conducted torsion tests on nano-bainitic steel specimens prepared by austempering at 250, 300 and 350°C which formed a range

of nanostructures. The main objective of the present study was to examine the effects of key factors like size, content, morphology and stability of phases on deformation behaviour under shear loading. In addition, hardness and phase fraction changes under shear loading were compared with uniaxial tensile loading. 2. Experimental procedures 2.1 Design of steel As-cast steel of chemical composition Fe-0.85C-1.92Mn-1.30Si-0.29Mo-2.05Co0.44Al (in wt.%) was made using vacuum induction melting. After mild homogenization of the as-cast steel at 1000°C for 48 hours in a muffle furnace, the steel was hot rolled at 1000°C to reduce the thickness from 34 mm to 14 mm in four steps. The nanostructure was generated by austenitizing at 950°C for 40 minutes followed by isothermal bainitic transformation in a salt bath at three different temperatures 350, 300 and 250°C for 20, 30 and 40 h respectively. The samples were designated according to the austempering temperatures as NB350, NB300 and NB250 respectively. 2.2 Microstructural characterization Samples for optical and electron microscopy were machined using electron discharge machining (EDM) and conventional sample preparation method was adopted for grinding and final polishing with 0.5 µm diamond paste. Well-polished samples were etched using 2 % nital solution and subsequently examined using AmScope optical microscope and Hitachi S-3400N scanning electron microscope (SEM). In order to measure the thickness of BF and RA laths, perpendicular test lines were drawn between the longer edges of a lath in electron micrographs. The true average thickness of BF and RA laths were obtained after making stereological corrections [25]. X-Ray diffraction (XRD) experiments were

performed using PANalytical X’Pert Pro MPD instrument using Cu-Kα radiation. X’Pert Highscore software was used for detailed analysis of XRD patterns and calculation of phase fraction was done using direct comparison method [26]. Calculation of dislocation density, ρ, was made by means of micro strain data obtained from XRD analysis. Micro-strain, ε and dislocation density ρ are related as



⁄

6

⁄

, where F = 1, K = 12A (A = π/2) and b is the

magnitude of Burgers vector in α ferrite [27]. XRD analysis was also done after the tension and the torsion test. The location of sample for XRD experiment in the torsion tested specimen is shown in Figure 1.

Figure 1. Location of the sample for X-ray diffraction, SEM imaging and hardness measurements after the torsion test

2.3 Specimen preparation and test parameters Standard tensile specimens of gauge length 25 mm were cut using EDM along the transverse direction of rolling. The tensile tests were done according to ASTM E8 [28] using an Instron 250 kN machine at a cross head speed of 0.5 mm/minute. Specimens for torsion tests were cylinders of diameter 3 mm and length 110 mm that were cut using EDM out of a heat-treated block along the transverse direction of rolling. Torque vs. angle of twist data was obtained by performing torsion tests on Instron 225 Nm capacity machine. The tests were conducted at an angular speed of 60° per minute. All the specimens had deformed plastically after an initial elastic deformation region. Therefore, calculation of shear stresses at the surface had to be done using elasto-plastic analysis. Shear stress(τ) and strain(γ) were

calculated from the torque (T) and angle of twist (θ) data using the method introduced by Bailey [29] and formulations are as follows:



(1) 3

(2)

where r is the radius and L is the length of specimen. It is to be noted that the reported shear stress and strain are for the surface part of the test specimen. Vickers micro-hardness tester was used to measure the hardness of as-transformed steels and mechanically tested specimens. The reported hardness values are the mean of the results obtained after a minimum 10 indentations made at onekilogram force. 3. Results 3.1 Nanostructures Figure 2 shows the austempering temperature dependence of size and volume fraction of bainitic sheaves and the bright regions show the coarse austenite but the finer films could not be resolved. However, it is evident that the microstructure gets finer and the fraction of BF increases with a reduction in austempering temperature.

Figure 2. Optical images of samples (a) NB250, (b) NB300 and (c) NB350

Figure 3 shows the corresponding SEM images that were used to determine the lath and film thickness of BF and RA. The true mean lath thickness of BF and RA are stated in Table 1. The thickness of BF as well as RA laths decreases by lowering the austempering temperature.

Figure 3. SEM micrographs of samples (a) NB250, (b) NB300 and (c) NB350 Table 1. Summary of microstructural measurements: lath thickness-T, carbon content- C, volume fraction of phase-V and dislocation density-ρ [9] Specimen

Tα nm

Tγ nm

Cγ(wt.%)

Vα

Vγ

ρα m-2

NB250

37±9

22±7

1.87

0.82

0.18

11.1×1015

NB300

47±12

45±12

1.66

0.71

0.29

6.1×1015

NB350

101±44

177±31

1.52

0.51

0.49

3.5×1015

Volume fraction of BF and RA and the dislocation density in BF obtained from XRD analysis are reported in the Table 1. Austempering at 250°C for 40 h led to formation of more than 0.80 volume fraction of BF. However only 0.51 volume fraction of bainite could form in case of austempering at 350°C for 20 h. Microstructural refinement along with increased volume fraction of BF can be attributed to the higher Gibbs free energy available at lower transformation temperature [10]. Refinement is further aided by the fact that the strength of austenite increases with a reduction in temperature. A decrease in austempering temperature also showed higher dislocation density in BF as can be seen in Table

1. BF takes the morphology of fine laths to minimize the elastic strain associated with displacive transformation and also due to an increased resistance to movement of interface when austenite has higher strength at lower transformation temperature [10]. When the strain energy is large enough to not get accommodated elastically by the parent phase austenite, austenite deforms plastically with generation of large number of dislocations which get inherited in transformed phase BF [30]. The higher carbon content in RA of steel austempered at lower temperature can be explained by To concept [10]. Since the stability of RA towards martensitic transformation is determined by its morphology and carbon content, a lower austempering temperature yields a relatively stable RA. 3.2 Tensile properties Table 2 summarizes the tensile properties obtained in our previous work on the same heat-treated steel blocks [9]. There is an increase yield strength and ultimate tensile strength with a reduction in austempering temperature. This is owing to formation of finer bainite and that too in greater volume fraction. The ductility is enhanced with an increase in austempering temperature due to greater volume fraction of austenite which is a softer phase than ferrite and is also capable of TRIP effect to enhance the ductility even further. Table 2. Results of tensile test on produced nano-bainitic steels [9] Specimen

Yield strength

Ultimate tensile strength

(MPa)

(MPa)

NB250

1560±32

1807±156

0.072±0.001

NB300

1382±20

1676±7

0.14±0.02

NB350

1028±52

1285±27

0.25±0.03

3.3 Torsional properties

Tensile strain at fracture

Figure 4 shows the torsion test results of the specimens austempered at the different temperatures. As expected, a lower austempering temperature shows better shear strength, but there is a concomitant loss in torsional ductility. The 0.2% offset shear strength, ultimate shear strength and strain at fracture are listed in Table 3. It can be seen that the ductility for the NB350 specimens is only marginally higher than that for NB300 specimens.

Figure 4. Shear stress versus shear strain plot of produced nano-bainitic steels Table 3. Results of torsion test on nano-bainitic steels austempered at three different temperatures Specimen

Shear strength

Shear stress at

Shear strain at

(MPa)

fracture (MPa)

fracture

NB250

956±8

1344±6

0.21±0.01

NB300

720±18

1209±8

0.77±0.05

NB350

542±20

1046±25

0.81±0.006

3.4 X-ray diffraction and hardness measurements The XRD patterns in Figure 5 were obtained from the gauge section of the torsion tested specimens post testing. Table 4 compares the content of RA and Vickers hardness of un-deformed and mechanically tested specimens. It is clear that more martensite is produced in the gauge section of the torsion tested specimens as compared to the tensile tested specimens.

Figure 5. XRD patterns of the gauge section of the torsion tested specimens Table 4. Content of RA and Vickers hardness of un-deformed and mechanically tested specimens Content of retained austenite in fraction Specimen Un-deformed

Tensile

Torsion

tested

tested

Vickers hardness, HV1 Un-deformed

Tensile

Torsion

tested

tested

NB250

0.18

0.16

< 0.05*

NB300

0.29

0.24

0.14

NB350

0.49

0.24

0.22

650.8±7.1

695.4±9.8

558.2±15.1 611.9±12.9 421±21.1

553.4±9.5

772.4±23.8 708.7±32 617.5±34.4

Figure 6. Fractographs of torsion tested (a) NB250, (b) NB300 and (c) NB350 specimen

3.5 Fractography Figure 6 shows the fractographs of torsion tested specimens. Since shear strain varies linearly with distance from the center to the surface of a torsion specimen, different surface features were seen along the radius. Therefore, an image at low magnification has been taken to show the overall surface morphology of fractured specimen. The magnified images taken near to the center and in between the center

and the edge (marked as squares) are shown on the right hand side and left hand side respectively. 4. Discussion It has been observed in prior studies that the tensile strength of bainitic steels varies linearly with the volume fraction of BF (Vα) divided by the bainitic lath thickness (t) [4,16,31]. Figure 7 shows the yield strength and shear strength of produced steels with respect to Vα/t calculated from microstructural observations. It is evident that the shear strength also varies approximately linearly with the Vα/t.

Figure 7. Yield strength and shear strength versus volume fraction of BF divided by its lath thickness

However, Figure 4 shows that the torsional ductility degrades with a decrease in austempering temperature. The plots also show the higher work hardening rate in steel transformed at higher austempering temperature. The jerky stress-strain plots

observed in martensitic steels by Malinov et al.[23] were not seen in the present work. This suggests that RA is relatively stable and shows a gradual transformation to martensite. NB250 steel showed an extremely high shear strength (0.61×yield strength) and hence much greater than what would be predicted based on Tresca criteria. The shear strength of NB300 and NB350 is also higher than what would be predicted by applying the Tresca criteria to the tensile test results. The primary reason may be the non-uniformity of stresses in torsion when compared with tensile testing. Moreover, stress-induced martensitic transformation may be responsible for increase in shear yield strength beyond what is predicted by the Tresca criteria since hard martensite replaces the soft RA. It is well known that the mechanical stability of RA depends on its morphology and the carbon content [10,32]. A recent study has shown that lath thickness and dislocation density of BF also affects the mechanical stability of RA [33]. NB250 steel has the lowest bainitic lath thickness and highest dislocation density among the produced steels. However, it has the lowest fraction of RA available to transform to martensite during deformation. Its limited initial RA content of around 0.18 in NB250 steel may be a possible reason for not showing much strain hardening and torsional ductility. Nevertheless, NB300 and NB350 steels showed significant work hardening, but fractured at similar torsional ductility. NB300 steel showed very uniform microstructure with similar lath thickness of BF and RA. However in case of NB350 steel, microstructural measurements showed large deviation in lath thickness of BF as well as RA. Moreover, it is also to be noted that average lath thickness of RA is greater than that of BF in NB350 steel. Coarse laths with lesser content of carbon in NB350 steel compared to NB250 and NB300 steels lead to mechanical destabilization of RA. This leads to formation of plenty of brittle martensite and causes early failure of NB350 steel and hence not a significantly higher torsional ductility than NB300 in spite of having much higher RA to start

with. Thus, it means that the torsional ductility is controlled by both the content and stability of RA in the steel. Figure 5 shows the XRD patterns of the gauge section of the torsion tested specimens. Table 4 shows the RA content and hardness of steels after the tensile and the torsion test. It was found that content of RA decreased from 0.49 to 0.22 for NB350 and from 0.29 to 0.14 for NB300 steel respectively. However, it was not possible to determine the phase fraction of RA in case of NB250 specimen because of insufficient intensity of peaks corresponding to RA and hence it was assumed to be below 0.05. Table 4 also shows a greater change in content of RA in NB250 and NB300 steels after the torsion test compared to the tension test which is also supported by the change in their hardness after the tensile and torsion tests respectively. However, NB350 steel showed similar RA content after the torsion and the tensile tests. This suggests that the extent of transformation-induced plasticity during the torsion test of NB350 steel is not to the extent that could be envisaged based on the results of NB250 and NB300 steel. This has also been reflected by torsional ductility of NB350 steel being close to that of NB300 steel although NB350 had the highest content of RA prior to the test. The possible reason for early failure of NB350 can be accumulation of a large amount of brittle martensite. Change in the hardness of steel has been found to be greater for torsion test when compared to tensile test. This can be attributed to the formation of larger amount of martensite due to the greater amount of deformation before fracture in case of torsion test. Figure 8 shows the SEM micrographs near the edge of the torsion tested specimens. The location of these SEM samples are shown in the transverse section (marked as a circle) by a schematic in Figure 1. It is clear that there is no significant change in the morphology of bainite near the region where the shear

strain is maximum. Since nital preferentially etches away BF which appears dark grey in the microstructure, it is difficult to make out whether fragmentation has happened or not. However, RA clearly appeared un-fragmented and intact even after the torsion test. This makes it clear that the RA was mechanically stable and a great amount of energy has been spent in transformation of RA to martensite till fracture which was also reflected by change in the hardness after the torsion tests as given in Table 4. SEM is not able to discern this transformation of austenite to martensite due to lack of resolution although the occurrence of transformation has been proved by the XRD studies (Table 4 and Figure 5). Anton Hohenwarter [24] reported strong alignment and severe fragmentation of microstructure in case of a carbide free bainitic steel that underwent high pressure torsion. These extreme levels of bending and subsequent fragmentation can be due to the severe plastic deformation (Von-Mises strain up to 16) in the aforesaid study which could be sustained without fracture due to the compressive stress that were applied in addition to torsion [24]. It was reported that all the RA was transformed to martensite in the aforementioned high pressure torsion study [24]. However, in the current study RA is not completely transformed in NB300 and NB350 but there is hardening due to partial transformation of RA to martensite. Although martensite could not be resolved in the micrographs, XRD results support the hardening behavior by showing that a great fraction of RA has reduced.

Figure 8. SEM micrographs of the torsion tested (a) NB250, (b) NB300 and (c) NB350 specimen

It can be seen in the fractographs (Figure 6) that the features appears similar to the fractographs of fracture toughness tested specimens [9]. However, there appears more tortuosity in the specimen with coarser microstructure. Images on the left hand side are taken in smoother region where elongated/stretched dimples were clearly visible. These features of a ductile shear fracture were found in all the specimens. The dimples are flatter and more stretched in case of NB250 specimen. One can clearly see the increase in size and protuberance of these dimples as we go from NB250 to NB350 specimen. The dimples in NB250 specimens are elliptical whereas they get almost equiaxed in NB350 specimen. It seems that the dimples had been formed in the crack nucleation stage. After the critical crack has been nucleated, there is a Mode III fracture (tearing mode) which propagates towards the center leading to cleavage kind of facets near the center of the specimen. 5. Conclusions The role of bainitic lath and RA size and their content in influencing the deformation behavior under torsion is investigated and compared with the response for the same kind of steels under tension. The following key conclusions can be drawn from this study: a. The shear strength under torsion improves with a decrease in austempering temperature. This is due to the increase in volume fraction of BF along with its refinement. b. The torsional ductility improves with an increase in austempering temperature till transformation at 300°C. This is due to higher volume fraction of the relatively ductile phase RA as well as TRIP effect. c. Alignment of RA films along the shear direction was not observed unlike high pressure torsion studies.

d. The increase in hardness in the gauge section under torsional deformation is much more than what has been observed for tensile deformation. This is because of great extent of martensitic transformation of RA under shear loading which continues till high strains due to absence of necking. e. The torsional ductility of the specimens austempered at 350°C does not greatly improve even though the microstructure comprised of higher austenite volume fraction when compared with the specimens austempered at 300°C. This is because of the presence of coarse retained austenite in the initial microstructure that transforms into brittle martensite which degrades the ductility during subsequent deformation. Thus, greater amount of austenite does not guarantee higher ductility under torsional loading.

Acknowledgement Authors would like thank the staff of ferrous process lab, Department of Metallurgical Engineering and Materials Science, IIT Bombay for their assistance in carrying out the experiments.

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Declaration of interests ☒ The 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. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: