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A nanotwinned surface layer generated by high strain-rate deformation in a TRIP steel
Materials and Design 80 (2015) 144–151
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A nanotwinned surface layer generated by high strain-rate deformation in a TRIP steel P. Xie a, C.L. Wu a,⇑, Y. Chen a, J.H. Chen a,b, X.B. Yang a, S.Y. Duan a, N. Yan b, X.A. Zhang c, J.Y. Fang c a
Center of High Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, Changsha 410082, China Advanced Research Center of Central South University, Changsha 410083, China c College of Science, National University of Defense Technology, Changsha 410073, China b
a r t i c l e
i n f o
Article history: Received 19 February 2015 Revised 27 April 2015 Accepted 6 May 2015 Available online 7 May 2015 Keywords: Surface strength Strain-rate TRIP effect TWIP effect Thermal stability
a b s t r a c t Gradient nanotwinned layers (GNTLs) with high hardness of approximately 500 HV and good thermal stability were produced in a Fe–20Mn–3Al–3Si TRIP steel by means of a high strain-rate surface mechanical grinding treatment (SMGT). The effect of the strain-rate on the plasticity-enhancing mechanisms in steels was investigated. It is found that although it primarily exhibits transformation-induced plasticity (TRIP) when cold-rolled, this steel shows twinning-induced plasticity (TWIP) during the SMGT process, illustrating that the steel may undergo a transition from TRIP to TWIP under high-strain-rate deformation. The martensite induced by deformation is thermally unstable and can easily transform back to austenite during annealing. In contrast, the deformation twins (DTs) are thermally much stable, since de-twinning, which plays a key role in the recovery of DTs, is more difficult to occur during thermal annealing. As such, undergoing the same annealing at 600 °C for 1 h, the GNTLs containing a great many of DTs maintain high hardness, whereas the cold-rolled counterpart samples containing deformation-induced martensite softens drastically. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Recently, it has been reported that a gradient nanotwinned layer (GNTL) with a thickness of about 500 lm can be generated on the surface of a Fe–Mn austenitic steel by means of surface mechanical grinding treatment (SMGT) [1]. The hardness in the topmost surface treated by SMGT was as high as 5.3 GPa, and the GNTL enhances the material strength [1]. Besides, a gradient hierarchical nanotwinned structure, which was generated by pre-torsion, can make the yield strength of FeMnC samples to be doubled at no reduction in ductility [2]. So it is reasonable to believe that a gradient hierarchical nanotwinned structure or GNTL is suited to improve the surface strength of mechanical tools and parts. The formation of the GNTL is dependent on deformation mechanisms. Deformation mechanisms are strongly affected by the stacking fault energy (SFE) [3–8], deformation temperature [9–11], strain rate [10,12–16] and so on. For Fe–Mn–Al–Si TRIP/TWIP steels, it is believed that the twinning induced plasticity (TWIP) effect occurs when the SFE is in the range of 18–45 mJ m2, whereas the transformation induced plasticity (TRIP) effect ⇑ Corresponding author. E-mail addresses: [email protected] (C.L. Wu), [email protected] (J.H. Chen). http://dx.doi.org/10.1016/j.matdes.2015.05.017 0261-3069/Ó 2015 Elsevier Ltd. All rights reserved.
appears when the SFE is below 18 mJ m2 [6–8]. The SFE of a Fe–Mn TRIP/TWIP steel usually depends on the chemical composition [4,5,17]. It is found that the SFE values of Fe–Mn–Al–Si TRIP/TWIP steels increases with the Mn content [17]. Therefore, it is known that a Fe–20Mn–3Al–3Si alloy primarily presents the TRIP effect during deformation at room temperature [18], whereas a Fe–30Mn–3Al–3Si steel exhibits the TWIP effect [19]. On the other hand, temperature also affects the SFE [6,9–11,20,21], and further affects the deformation mechanisms of TRIP/TWIP steel. A Fe–20Mn–3Al–3Si alloy presents the TRIP effect at room temperature but exhibits the TWIP effect at 160 °C during straining [9]. The strain-rate also has an influence on the deformation mechanisms. For a Fe–12Mn–0.6C steel, deformation at low strain rates favors the e-martensite transformation while high strain-rate deformation results in mechanical twinning [16]. In the present study, SMGT has been employed to produce the GNTLs in a Fe–20Mn–3Al–3Si TRIP steel, in which the martensitic transformation is dominant when the steel is subjected to a conventional deformation process at room temperature. In our previous work on the same alloy, it has been shown that a gradient phase-transformation-strengthened surface layer can be formed by surface mechanical grinding with low-strain-rate deformation at room temperature, owing to the c ? e ? a martensitic transformation [22]. The gradient phase-transformation-strengthened
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surface layer is usually thin and easily ground off. Here, we report that when the strain-rate increases in SMGT, a GNTL with high hardness and good thermal stability will be generated, replacing the phase-transformation-strengthened surface layer. Although it has been reported [1] that a GNTL can be generated on the surface of a Fe–28Mn–3Al–3Si austenitic TWIP steel by means of SMGT, the GNTL produced in a Fe–20Mn–3Al–3Si TRIP steel has never been reported. Interesting is the transition of the steel from TRIP to TWIP, which occurs only under high-strain-rate deformation and has not previously been investigated in detail. In addition, the thermal stabilities of such a GNTL are important to know for applications. And furthermore, associated with this issue, the recovery mechanism of deformation twins (DTs) embedded in coarse grains needs to be studied, though the mechanical properties and thermal stability of DTs have been investigated [22,23]. In the present study, we demonstrate that in a Fe–20Mn–3Al– 3Si TRIP steel, the TRIP effect occurs in a cold-rolling process while the TWIP effect appears in a high-strain-rate SMGT process, the latter forms a GNTL on the surface of the steel. The hardness of the cold-rolled sample is about 405 HV, but sharply decreases to 240 HV during annealing at 600 °C for 1 h. The surface hardness of the SMGT samples with GNTLs is about 500 HV, and decreases only slightly to 400 HV during annealing at 600 °C for 1 h. De-twinning plays an important role in the recovery of DTs embedded in the original austenitic grains, and affects the thermal stabilities of the GNTLs. 2. Experimental procedures 2.1. As-received material The material selected for this study was a Fe–20Mn–3Al–3Si TRIP steel containing 20.3 wt.% Mn, 2.64 wt.% Al, 2.64 wt.% Si, 0.006 wt.% C, 0.0012 wt.% S. Using weak beam dark-field (WBDF) imaging in transmission electron microscopy (TEM) [17,25], the SFE value of materials can be estimated by measuring the dissociation width of Shockley partial dislocation pairs. In this work, the SFE value of the alloy was determined to be 8 ± 2 mJ m2 by measuring the dissociation widths of several Shockley partial
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dislocation pairs. Fig. 1 shows a WBDF image of a Shockley partial dislocation pair in a sample deformed to 1% at room temperature. The detailed principles for the calculation of SFE values can be found in the literature [25]. 2.2. Preparation of samples and heat-treatments The Fe–20Mn–3Al–3Si TRIP steel was forged at 1200 °C, and then cut into cylindrical samples or sheets by electrical discharge machining. All samples were annealed at 1000 °C for 1 h, forming equiaxed austenite grains with an average size of approximately 55 lm. The annealed sheets were first hot rolled and then cold rolled into thin plates with a strain of approximately 0.65. The cylindrical specimens of 15 mm in diameter were processed by means of SMGT at room temperature. The detailed principles of SMGT have been described in the literature [1]. The main parameters of the SMGT, including the WC/Co tip diameter (dt), moving velocity of the tip (vt), rotation velocity of the sample (vs) and preset indentation depth (d), are as follows: dt = 8 mm, vt = 18 mm min1, vs = 403 r min1, and d is variable. In SMGT, a water-based coolant is employed to cool the samples. The SMGT samples produced at d = 32.5 lm were cut into disks with a thickness of approximately 2 mm. Cold-rolled (CR) and SMGT samples (with d = 32.5 lm) were annealed at 450, 500, 525, 550, 575, 600, 625, 650, 675 and 700 °C for 1 h, respectively, and then air-cooled. 2.3. Characterization methods A Siemens D5000 X-ray diffractometer (XRD) with Cu/Ka radiation, an FEI Quanta 200 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detector, a JEOL 3010 transmission electron microscope (TEM) and a FEI Tecnai G2 F20 TEM were used for microstructural examinations. All samples for characterization were first electro-polished using a solution of 10 vol.% perchloric acid and 90 vol.% ethanol to remove the stress layer produced during sample preparation. The samples for SEM observation were finally etched with 5 vol.% nital solution. TEM samples at different depths were prepared by focused ion beam in an FEI Quanta 3D FEG Omniprobe microscope. To study the phase structure of the SMGT samples at different depths using XRD, the specimens were examined at different depths by electro-polishing off sequential layers. Quantifying deformation twins in the SMGT samples was performed by using electron channeling contrast imaging (ECCI) in a scanning electron microscope. The detailed principles for ECCI have been described in the literature [25–27]. The hardness of the CR samples was examined using a HXD-1000T Vickers hardness tester with an applied load of 500 g, and that of the SMGT samples was tested with an applied load of 50 g. For the CR samples and the top-most deformation layers of SMGT samples, at least ten indentations were conducted to obtain the average hardness value and assess the error. 3. Results and discussion 3.1. Hardness
Fig. 1. WBDF TEM image of a Shockley partial dislocations pair in a sample deformed to 1% at room temperature. The partial dislocation pair is in the (1 1 1) 0] g-vector. The label ‘‘d’’ habit plane normal to the [1 1 1] beam, imaged with a [2 2 represents a dissociation width of the partial dislocation pair. The effective dissociation width of each partial dislocation pair is the average value of all of the ‘‘d’’.
In SMGT, plastic deformation takes place on the surface of cylindrical samples, and is affected by the preset indentation depth d. In the present work, cylindrical samples were processed by the SMGT at values of d = 15, 20 and 32.5 lm, respectively. Fig. 2a shows the hardness distribution of the surface deformation layer (SDL) as a function of the depth from the surface. For all SMGT samples, the surface hardness of the top-most deformation layer is as high as 500 ± 20 HV, and the hardness gradually decreases to approximately 230 HV (the hardness of the matrix) with increasing depth.
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presents the hardness values of SDLs at different depths (about 15, 40, and 150 lm in depth, respectively) and that of the cold-rolled samples as a function of annealing temperature. The hardness of the CR samples is approximately 405 ± 8 HV, similar to that of the sub-layer of the SMGT samples. After annealing at 450 °C for 1 h, the hardness of the CR samples decreases slightly because of recovery. The hardness decreases sharply from approximately 400 to 240 HV when the annealing temperature increases from 450 to 600 °C, and this decrement is attributed to recrystallization. However, the hardness decreasing of SDLs shifts towards higher temperatures in the process of thermal annealing. For the SDL, the hardness decrease is not obvious when annealed below 575 °C. It is interesting to note that the hardness of the outer layer annealed at 625 °C for 1 h is still larger than 400 HV, and therefore the thermal stability of the SDLs is better than that of the CR samples. 3.2. Microstructure
Fig. 2. (a) Hardness distribution vs. depth from the surface of SMGT samples with different preset indentation depth d and (b) hardness profiles vs. annealing temperature, the hardness of the SDL at different depths is tested on the same SMGT sample (at d = 32.5 lm).
However, the thickness of the hardened layer increases with increasing d. The thicknesses of the SDLs are about 300, 330 and 900 lm for d = 15, 20 and 32.5 lm, correspondingly. Fig. 2b
3.2.1. Samples treated by surface mechanical grinding The SDL microstructures of the SMGT samples are similar, so only the microstructure of those formed at d = 32.5 lm is shown here. Fig. 3 shows the depth-dependent gradient microstructure of the SDL. The top-most layer mainly consists of equiaxed nano-grains, as confirmed by TEM observation. Thus, the top-most layer is called the nano-grained layer (NGL). Below the NGL, there are many lamellar structures produced by the SMGT process. Closer to the surface, more distorted grains can be observed. Fig. 3b shows the grains and lamellar structures embedded in severely distorted grains. The lamellar structures are entangled with numerous defects. For depths larger than 200 lm, the shapes and sizes of grains remain unchanged, though many lamellar structures exist (as shown in Fig. 3a and c). In addition, Fig. 3c shows that two groups of crossed lamellar structures can form in an austenitic grain. To investigate the SDLs, a series of TEM samples at different depths were examined. Fig. 4a shows that there
Fig. 3. Microstructure of the SMGT samples. (a) Cross-sectional SEM image and (b and c) showing EBSD inverse pole image of the SDL at depths of 45 lm and 250 lm, respectively.
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Fig. 4. TEM images of the nano-grained layer (NGL). (a) Dark-field TEM image of the NGL at a depth of 25 lm, the inset is the corresponding selected area electron diffraction patterns (SAEDP) and (b) bright-field TEM image of nanotwins at a depth of 25 lm, the inset shows a high resolution TEM image of the area marked by the arrow in Fig. 4b.
are a great number of equiaxed nano-grains in the NGL. Some grains contain nano-twins, as shown in Fig. 4b. Statistical analysis of large numbers of TEM images indicates that the average size of the nano-grains at a depth of 25 lm is approximately 100 nm. The average grain size increases with depth. TEM observation shows that the lamellar structures below the NGL are mainly DTs and austenitic laths divided by DTs (as shown in Fig. 5a), and no martensitic laths can be found. Because the DTs are very thin (less than 20 nm in thickness) and contain large numbers of dislocations that cannot be detected by EBSD/SEM. Some austenitic laths form secondary DTs (as shown in Fig. 5b). Moreover, the interaction
between the primary DTs and the secondary DTs causes the austenitic laths to twist, forming dislocation walls, as shown in Fig. 5c and d. At a depth of 600 lm, the number of nanotwinned bundles decreases drastically, but the interaction between twins and dislocations is also observed. Because no martensite can be found using TEM and EBSD, XRD was employed to evaluate the phase structure of the SDLs. Fig. 6 shows the XRD patterns of the SDLs at different depth. Except for the occurrence of the austenitic (c) peaks, no martensitic peaks can be found in any of the SDLs, so no martensitic transformation occurs in the process of SMGT.
Fig. 5. TEM images of a nano-twinned layer (NTL). (a) TEM morphology at a depth of 200 lm, and insert figure is the SAEDP of the twinned bundles; (b) the austenite lath contains the secondary DTs at a depth of 200 lm; (c) TEM image of a distorted lath contain the secondary twins and a dislocation wall at a depth of 200 lm; (d) dark-field TEM image of (c) and its SAEDP, showing approximately 6.5° misorientation between the ‘‘A’’ and ‘‘B’’ areas. The SAEDPs in (a) and (d) are acquired along the [0 1 1]c zone axis.
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Fig. 6. XRD patterns of the SDLs at different depths.
3.2.2. SMGT samples annealed at different temperature Because the hardness of the annealed SDLs is still very high, the microstructures of the SDLs annealed at different temperature were examined in detail. Fig. 7a–d shows cross-sectional SEM images of the SMGT samples annealed for 1 h at 575, 600, 650 and 700 °C, respectively. In the 575 °C-annealed sample, a large number of small grains are arranged along the original twin boundaries (TBs) to form numerous stripes in a plaid pattern, and most of the DTs in the SDLs are still present. With increasing annealing temperature, the number of DTs gradually decreases, and the recrystallized grains gradually grow (Fig. 7b). Fig. 7c shows that a large number of DTs disappear after annealing at 650 °C, and
Fig. 7. Cross-sectional SEM image of the SMGT samples annealed at different annealing temperature. (a) At 575 °C; (b) at 600 °C; (c) at 650 °C and (d) at 700 °C.
the size of the recrystallized grains increases up to several microns. After annealing at 700 °C for 1 h, most DTs disappear, and the recrystallization of the SMGT samples is nearly complete, as shown in Fig. 7d. It is interesting that the nanograins in the top-most layer do not grow easily during thermal annealing, and their grain sizes are smaller than those of recrystallized grains in the sub-layer, as shown in Fig. 7. This finding means that the thermal stability of the NGL is higher than that of the deformation layer consisting of DTs and dislocations. However, the literature [24] reports that the thermal stability of nanograins is the lowest, with respect to that of DTs and dislocations. In this work, the high thermal stability of the NGLs is attributed to their low dislocation density and inclusion of embedded nanotwins (Fig. 4c). To understand the microstructure shown in Fig. 7a–c, EBSD was performed intensively. After annealing at 600 °C for 1 h, many small-angle grain boundaries (SAGBs) are evident either perpendicular or parallel to the original TBs (as shown in Fig. 8a and c), and the deformation twinned stripes are cut into many chain-like twinned cells (as shown in Fig. 8a and b). Some SAGBs are formed by dislocation walls generated during the SMGT process, while other SAGBs are attributed to dislocation reactions. The chain-like twinned cells must arise during the recovery of the primary DTs, companied by the interaction between dislocations and TBs. This interaction between dislocations and TBs causes DTs to disappear [28,29], and is consequently considered as de-twinning induced by the thermal activation. De-twinning not only reduces the volume of DTs but also eliminates dislocations. The areas around chain-like twinned cells are clean, indicating the low dislocation-density zones. The clean areas around chain-like twinned cells inhibit further de-twinning and therefore enhance the thermal stability of the DTs. These results illustrate that de-twinning must occur in the recovery of the DTs during high temperature annealing. The de-twinning plays a key role in the recovery of the DTs and determines the thermal stability of the twinned-microstructure. In addition, the recrystallization must occur within the high dislocation-density zones of the SMGT samples. Fig. 8a also shows the formation of many recrystallized grains in the lower right corner. 3.2.3. Cold-rolled and its annealed samples To understand the different deformation processes, the microstructures of the CR and its annealed samples were also investigated. The microstructure of the CR sample contains elongated and distorted grains and lamellar structures, as shown in Fig. 9a. TEM observations illustrate that the CR sample contains many of a0 -martensitic laths, distorted austenitic laths and small number of DTs, as shown in Fig. 9b and c. When annealing the CR samples at 450 °C for 1 h, some of the lamellar structures begin to recover. With the temperature increasing, the lamellar structures decrease, and the recrystallized grains gradually nucleate and grow. Fig. 9d shows the typical SEM morphology of the partially recrystallized CR sample, which contains both lamellar structures and recrystallization zones. TEM observations show that the DTs are more difficult to recover by annealing as compared to the a0 -martensite. Fig. 9f shows that the DTs do not change in the partially recrystallized CR sample, though recrystallized grains do appear around the DTs. After annealing at 600 °C for 1 h, the lamellar structures nearly disappear and recrystallization of the CR samples is nearly complete (Fig. 9e). The XRD analysis reveals the phase evolution of the CR samples during annealing, as shown in Fig. 10. The CR samples consist of many of a0 -martensite and austenite (c). The intensity of the a0 -martensitic peaks begins to diminish after annealing at 450 °C for 1 h and continuously decreases with temperature. After annealing at 600 °C for 1 h, a small number of a0 -martensitic peaks can be found. This result also indicates that the a0 -martensite formed by deformation is
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Fig. 8. EBSD analysis of the SMGT samples after annealing at 600 °C for 1 h. (a) EBSD inverse pole image of the SDL (at a depth of approximately 150 lm), the red and black lines represent high-angle and low-angle boundaries, respectively; (b) misorientation along the line ‘‘1–2’’, showing the twinning relationship and (c) misorientation along the line ‘‘3–4’’, showing the misorientation of low-angle boundaries. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 9. The microstructure of the CR samples with different annealing temperature for 1 h. (a) SEM image of the CR sample; (b) TEM image of a0 -martensitic lathes in the CR samples; (c) TEM image of DTs in the CR samples, and the SAEDP is acquired along the [0 1 1]c zone axis; (d) SEM image of the 550 °C-annealed CR samples, the labels ‘‘L’’ and ‘‘R’’ represent the lamellar structures and the recrystallization zones, respectively; (e) SEM image of the 600 °C-annealed CR samples and (f) TEM image of DTs in the partially recrystallized CR samples.
unstable, and it quickly undergoes a reverse transformation to austenite during annealing. 3.2.4. Thermal stability of the deformation twins and the a0 -martensite The volume fraction of the a0 -martensite (Va0 –M) was measured by XRD, and that of grains containing many of DTs (VDT) was statistically measured using ECCI in the SEM. Fig. 11 shows the Va0 –M of the CR samples and the VDT in the sublayer of the SMGT samples vs. annealing temperature. The Va0 –M of the CR samples is
approximately 52%, and decreases to 40% after annealing at 450 °C for 1 h. After annealing at 600 °C for 1 h, the Va0 –M decreases to approximately 5.8%, whereas the VDT in the sublayer of the SMGT samples begins to decrease. When the annealing temperature exceeds 625 °C, the VDT rapidly decreases. After annealing at 700 °C for 1 h, the VDT decreases to approximately 3.2%. This result further confirms that the DTs are more stable than the a0 -martensite. The full recrystallization of the SMGT samples occurs at 700 °C, while that of the CR samples takes place at 600 °C.
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Fig. 10. XRD patterns of the CR samples with different annealing temperature for 1 h.
Fig. 11. The volume fraction of deformation twins (VDT) in the sublayer of the SMGT samples and that of a0 -martensite (Va0 –M) of the CR samples as a function of annealing temperature.
3.3. Discussion As shown above, for all SMGT samples, the surface hardness of the top-most layer is as high as 500 ± 20 HV, which is not affected by the preset indentation depth (d). However, the thickness of the hardened layer increases with increasing d. The mechanical properties depend on the microstructure of the SDL formed during the SMGT process. The top-most layers of all SMGT samples consist of equiaxed nano-grains, some of which contain nano-twins. The surface hardness of the SMGT samples with a similar NGL must be much the same. Below the NGLs, there are many distorted grains, deformation twins with nano-thickness and dislocation structures. The number of these distorted grains, DTs and dislocation structures in the SDLs decrease with increasing the depth, so that the hardness of the SDLs changes correspondingly. Besides, the thickness of the gradient deformation layer is strongly affected by the preset indentation depth (d). In the present work, the thickness of the SDL is approximately 300, 330 and 900 lm for d = 15, 20 and 32.5 lm, correspondingly. No martensite can be found in the SMGT samples, but the martensite is the main component in the CR samples in which the TRIP effect is dominant. In the SMGT process, a very high-rate shear deformation is applied to the surface layer of the samples, where the estimated strain rate varies from 103–104 (at the top-most surface) to 102–103 (80 lm below the treated surface) [30]. Meanwhile the cumulative equivalent strains of the SMGT are approximately 15–30, which is much higher than that of cold rolling (approximately 0.65). High strain rate promotes
deformation twinning [10,29,31], so the SDLs of the SMGT samples contain large numbers of DTs. Moreover, the literature [11] reported that the temperature rise due to deformation heating will increase with the increases of strain rate and strain in Fe–23Mn–2Al–0.2C TWIP steel. The literature [32,33] reported that temperature increases as high as 450 °C at the shear band center when carbon steels are tested at strain rates in the range of 1000–5000 s1. When materials undergo high-strain-rate deformation, especially for cumulative deformation such as the SMGT in this work, large amounts of plastic work are performed in a very short time, which leads to an increase in temperature of the deformation zones. The samples were cooled by a water-based coolant, but we suspect that the temperature of the SDLs might increase instantaneously at the moment of high-rate shear deformation, after which the SDLs are quickly cooled by the applied water. Some phenomena also indicate the temperature rise due to deformation heating. For example, the low dislocation density and equiaxed nanograins in the NGL should be attributed to elevated temperature and dynamic recrystallization. According to the literature [6,20,21], the SFE value increases markedly with elevated temperature. Here, the SFE value of the SDLs probably increases instantaneously, which leads to the TWIP effect and suppresses the TRIP effect. In this work, the thermal stability of deformed samples is studied. The hardness of the CR samples, which contain a lot of a0 -martensite laths and elongated austenitic grains, is about 405 ± 8 HV. However, the a0 -martensite is relatively unstable. The Va0 –M of the CR samples quickly decreases from 40% to 5.8% when the annealing temperature increases from 450 to 600 °C, and the hardness sharply decreases from 400 to 240 HV correspondingly. It is worthy of noting that the hardness decrease of the SMGT samples is not obvious during annealing below 575 °C, and the hardness of the outer layer annealed at 625 °C for 1 h is more than 400 HV. The outstanding thermal stabilities of the GNTLs are ascribed to the DTs with good thermal stability. The DTs in the interior of coarse austenitic grains were investigated. In the SMGT process, some coarse austenite grains can forms many of multiple deformation twins, such as primary DTs and secondary DTs. These primary DTs and secondary DTs present two groups of crossed lamellar structures, as shown in Fig. 3c. Dislocations emitted from the secondary DTs must interact with the primary DTs, and leave many partial dislocations and stair-rod dislocations at the TBs [29]. When the SMGT sample is heated to a high temperature, these partial dislocations and stair-rod dislocations are activated and slip on the TBs, decreasing the thickness of the twins. De-twinning is easy at the beginning of the recovery of crossed DTs embedded in the interior of grains, which lowers the thermal stability of the DTs. When the deformation twinned stripes are cut into many chain-like twinned cells by de-twinning, the dislocation density must decrease, and de-twinning becomes difficult, which improves the thermal stability of DTs. This is to say that de-twinning plays an important role in the recovery of the DTs embedded in grains, and affects the thermal stability of DTs. In a word, ultrahard and ultrastable GNTLs can be generated on high manganese TRIP steels by means of a high strain-rate SMGT processing. Compared to conventional surface technique, the SMGT processing is simple and cost-effective. It is believed that the SMGT processing will become a new surface strengthening technique in industrial manufacturing processes.
4. Conclusion In summary, the gradient nanotwinned layers that contain many nanotwins and nanograins were produced on a Fe–20Mn–3
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Al–3Si TRIP steel by means of surface mechanical grinding treatment (SMGT). From the obtained results, the following can be concluded. (1) The gradient nanotwinned layers have high hardness, and the surface hardness can reach a value up to 500 HV. The hardness decreases from 500 to 230 HV with increasing depth, where the latter value is actually the hardness of the matrix. The thickness of the hardened layer increases with increasing the preset indentation depth d, and the thickness of the hardened layer can reach up to 900 lm at d = 32.5 lm. (2) In the SMGT process, a very high-rate shear deformation is applied to the surface layer of the samples, which promotes deformation twinning and suppresses martensitic transformation. (3) The martensite induced by deformation is unstable. The martensite content of the cold-rolled samples can be reduced from 52% to 40% after annealing at 450 °C for 1 h, and remains only by 5.8% after annealing at 600 °C for 1 h, indicating that the recrystallization in the annealed cold-rolled samples is nearly complete. (4) De-twinning, occurring during thermal annealing, plays the key role in the recovery of deformation twins embedded in grains and affects the thermal stability of deformation twins. The thermal stability of deformation twins embedded in grains is higher than that of martensite induced by deformation, but lower than that of equiaxial nano-grains in the nano-grained layers. For the SMGT samples, the volume fraction of the deformation twins begins to decrease at 600 °C for 1 h, and the hardness of the top-most layers can remain at a value above 400 HV even after annealing at 625 °C for 1 h, whereas the hardness of the cold-rolled counterpart samples decreases drastically to less than 240 HV.
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A nanotwinned surface layer generated by high strain-rate deformation in a TRIP steel P. Xie a, C.L. Wu a,⇑, Y. Chen a, J.H. Chen a,b, X.B. Yang a, S.Y. Duan a, N. Yan b, X.A. Zhang c, J.Y. Fang c a
Center of High Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, Changsha 410082, China Advanced Research Center of Central South University, Changsha 410083, China c College of Science, National University of Defense Technology, Changsha 410073, China b
a r t i c l e
i n f o
Article history: Received 19 February 2015 Revised 27 April 2015 Accepted 6 May 2015 Available online 7 May 2015 Keywords: Surface strength Strain-rate TRIP effect TWIP effect Thermal stability
a b s t r a c t Gradient nanotwinned layers (GNTLs) with high hardness of approximately 500 HV and good thermal stability were produced in a Fe–20Mn–3Al–3Si TRIP steel by means of a high strain-rate surface mechanical grinding treatment (SMGT). The effect of the strain-rate on the plasticity-enhancing mechanisms in steels was investigated. It is found that although it primarily exhibits transformation-induced plasticity (TRIP) when cold-rolled, this steel shows twinning-induced plasticity (TWIP) during the SMGT process, illustrating that the steel may undergo a transition from TRIP to TWIP under high-strain-rate deformation. The martensite induced by deformation is thermally unstable and can easily transform back to austenite during annealing. In contrast, the deformation twins (DTs) are thermally much stable, since de-twinning, which plays a key role in the recovery of DTs, is more difficult to occur during thermal annealing. As such, undergoing the same annealing at 600 °C for 1 h, the GNTLs containing a great many of DTs maintain high hardness, whereas the cold-rolled counterpart samples containing deformation-induced martensite softens drastically. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction Recently, it has been reported that a gradient nanotwinned layer (GNTL) with a thickness of about 500 lm can be generated on the surface of a Fe–Mn austenitic steel by means of surface mechanical grinding treatment (SMGT) [1]. The hardness in the topmost surface treated by SMGT was as high as 5.3 GPa, and the GNTL enhances the material strength [1]. Besides, a gradient hierarchical nanotwinned structure, which was generated by pre-torsion, can make the yield strength of FeMnC samples to be doubled at no reduction in ductility [2]. So it is reasonable to believe that a gradient hierarchical nanotwinned structure or GNTL is suited to improve the surface strength of mechanical tools and parts. The formation of the GNTL is dependent on deformation mechanisms. Deformation mechanisms are strongly affected by the stacking fault energy (SFE) [3–8], deformation temperature [9–11], strain rate [10,12–16] and so on. For Fe–Mn–Al–Si TRIP/TWIP steels, it is believed that the twinning induced plasticity (TWIP) effect occurs when the SFE is in the range of 18–45 mJ m2, whereas the transformation induced plasticity (TRIP) effect ⇑ Corresponding author. E-mail addresses: [email protected] (C.L. Wu), [email protected] (J.H. Chen). http://dx.doi.org/10.1016/j.matdes.2015.05.017 0261-3069/Ó 2015 Elsevier Ltd. All rights reserved.
appears when the SFE is below 18 mJ m2 [6–8]. The SFE of a Fe–Mn TRIP/TWIP steel usually depends on the chemical composition [4,5,17]. It is found that the SFE values of Fe–Mn–Al–Si TRIP/TWIP steels increases with the Mn content [17]. Therefore, it is known that a Fe–20Mn–3Al–3Si alloy primarily presents the TRIP effect during deformation at room temperature [18], whereas a Fe–30Mn–3Al–3Si steel exhibits the TWIP effect [19]. On the other hand, temperature also affects the SFE [6,9–11,20,21], and further affects the deformation mechanisms of TRIP/TWIP steel. A Fe–20Mn–3Al–3Si alloy presents the TRIP effect at room temperature but exhibits the TWIP effect at 160 °C during straining [9]. The strain-rate also has an influence on the deformation mechanisms. For a Fe–12Mn–0.6C steel, deformation at low strain rates favors the e-martensite transformation while high strain-rate deformation results in mechanical twinning [16]. In the present study, SMGT has been employed to produce the GNTLs in a Fe–20Mn–3Al–3Si TRIP steel, in which the martensitic transformation is dominant when the steel is subjected to a conventional deformation process at room temperature. In our previous work on the same alloy, it has been shown that a gradient phase-transformation-strengthened surface layer can be formed by surface mechanical grinding with low-strain-rate deformation at room temperature, owing to the c ? e ? a martensitic transformation [22]. The gradient phase-transformation-strengthened
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surface layer is usually thin and easily ground off. Here, we report that when the strain-rate increases in SMGT, a GNTL with high hardness and good thermal stability will be generated, replacing the phase-transformation-strengthened surface layer. Although it has been reported [1] that a GNTL can be generated on the surface of a Fe–28Mn–3Al–3Si austenitic TWIP steel by means of SMGT, the GNTL produced in a Fe–20Mn–3Al–3Si TRIP steel has never been reported. Interesting is the transition of the steel from TRIP to TWIP, which occurs only under high-strain-rate deformation and has not previously been investigated in detail. In addition, the thermal stabilities of such a GNTL are important to know for applications. And furthermore, associated with this issue, the recovery mechanism of deformation twins (DTs) embedded in coarse grains needs to be studied, though the mechanical properties and thermal stability of DTs have been investigated [22,23]. In the present study, we demonstrate that in a Fe–20Mn–3Al– 3Si TRIP steel, the TRIP effect occurs in a cold-rolling process while the TWIP effect appears in a high-strain-rate SMGT process, the latter forms a GNTL on the surface of the steel. The hardness of the cold-rolled sample is about 405 HV, but sharply decreases to 240 HV during annealing at 600 °C for 1 h. The surface hardness of the SMGT samples with GNTLs is about 500 HV, and decreases only slightly to 400 HV during annealing at 600 °C for 1 h. De-twinning plays an important role in the recovery of DTs embedded in the original austenitic grains, and affects the thermal stabilities of the GNTLs. 2. Experimental procedures 2.1. As-received material The material selected for this study was a Fe–20Mn–3Al–3Si TRIP steel containing 20.3 wt.% Mn, 2.64 wt.% Al, 2.64 wt.% Si, 0.006 wt.% C, 0.0012 wt.% S. Using weak beam dark-field (WBDF) imaging in transmission electron microscopy (TEM) [17,25], the SFE value of materials can be estimated by measuring the dissociation width of Shockley partial dislocation pairs. In this work, the SFE value of the alloy was determined to be 8 ± 2 mJ m2 by measuring the dissociation widths of several Shockley partial
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dislocation pairs. Fig. 1 shows a WBDF image of a Shockley partial dislocation pair in a sample deformed to 1% at room temperature. The detailed principles for the calculation of SFE values can be found in the literature [25]. 2.2. Preparation of samples and heat-treatments The Fe–20Mn–3Al–3Si TRIP steel was forged at 1200 °C, and then cut into cylindrical samples or sheets by electrical discharge machining. All samples were annealed at 1000 °C for 1 h, forming equiaxed austenite grains with an average size of approximately 55 lm. The annealed sheets were first hot rolled and then cold rolled into thin plates with a strain of approximately 0.65. The cylindrical specimens of 15 mm in diameter were processed by means of SMGT at room temperature. The detailed principles of SMGT have been described in the literature [1]. The main parameters of the SMGT, including the WC/Co tip diameter (dt), moving velocity of the tip (vt), rotation velocity of the sample (vs) and preset indentation depth (d), are as follows: dt = 8 mm, vt = 18 mm min1, vs = 403 r min1, and d is variable. In SMGT, a water-based coolant is employed to cool the samples. The SMGT samples produced at d = 32.5 lm were cut into disks with a thickness of approximately 2 mm. Cold-rolled (CR) and SMGT samples (with d = 32.5 lm) were annealed at 450, 500, 525, 550, 575, 600, 625, 650, 675 and 700 °C for 1 h, respectively, and then air-cooled. 2.3. Characterization methods A Siemens D5000 X-ray diffractometer (XRD) with Cu/Ka radiation, an FEI Quanta 200 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detector, a JEOL 3010 transmission electron microscope (TEM) and a FEI Tecnai G2 F20 TEM were used for microstructural examinations. All samples for characterization were first electro-polished using a solution of 10 vol.% perchloric acid and 90 vol.% ethanol to remove the stress layer produced during sample preparation. The samples for SEM observation were finally etched with 5 vol.% nital solution. TEM samples at different depths were prepared by focused ion beam in an FEI Quanta 3D FEG Omniprobe microscope. To study the phase structure of the SMGT samples at different depths using XRD, the specimens were examined at different depths by electro-polishing off sequential layers. Quantifying deformation twins in the SMGT samples was performed by using electron channeling contrast imaging (ECCI) in a scanning electron microscope. The detailed principles for ECCI have been described in the literature [25–27]. The hardness of the CR samples was examined using a HXD-1000T Vickers hardness tester with an applied load of 500 g, and that of the SMGT samples was tested with an applied load of 50 g. For the CR samples and the top-most deformation layers of SMGT samples, at least ten indentations were conducted to obtain the average hardness value and assess the error. 3. Results and discussion 3.1. Hardness
Fig. 1. WBDF TEM image of a Shockley partial dislocations pair in a sample deformed to 1% at room temperature. The partial dislocation pair is in the (1 1 1) 0] g-vector. The label ‘‘d’’ habit plane normal to the [1 1 1] beam, imaged with a [2 2 represents a dissociation width of the partial dislocation pair. The effective dissociation width of each partial dislocation pair is the average value of all of the ‘‘d’’.
In SMGT, plastic deformation takes place on the surface of cylindrical samples, and is affected by the preset indentation depth d. In the present work, cylindrical samples were processed by the SMGT at values of d = 15, 20 and 32.5 lm, respectively. Fig. 2a shows the hardness distribution of the surface deformation layer (SDL) as a function of the depth from the surface. For all SMGT samples, the surface hardness of the top-most deformation layer is as high as 500 ± 20 HV, and the hardness gradually decreases to approximately 230 HV (the hardness of the matrix) with increasing depth.
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presents the hardness values of SDLs at different depths (about 15, 40, and 150 lm in depth, respectively) and that of the cold-rolled samples as a function of annealing temperature. The hardness of the CR samples is approximately 405 ± 8 HV, similar to that of the sub-layer of the SMGT samples. After annealing at 450 °C for 1 h, the hardness of the CR samples decreases slightly because of recovery. The hardness decreases sharply from approximately 400 to 240 HV when the annealing temperature increases from 450 to 600 °C, and this decrement is attributed to recrystallization. However, the hardness decreasing of SDLs shifts towards higher temperatures in the process of thermal annealing. For the SDL, the hardness decrease is not obvious when annealed below 575 °C. It is interesting to note that the hardness of the outer layer annealed at 625 °C for 1 h is still larger than 400 HV, and therefore the thermal stability of the SDLs is better than that of the CR samples. 3.2. Microstructure
Fig. 2. (a) Hardness distribution vs. depth from the surface of SMGT samples with different preset indentation depth d and (b) hardness profiles vs. annealing temperature, the hardness of the SDL at different depths is tested on the same SMGT sample (at d = 32.5 lm).
However, the thickness of the hardened layer increases with increasing d. The thicknesses of the SDLs are about 300, 330 and 900 lm for d = 15, 20 and 32.5 lm, correspondingly. Fig. 2b
3.2.1. Samples treated by surface mechanical grinding The SDL microstructures of the SMGT samples are similar, so only the microstructure of those formed at d = 32.5 lm is shown here. Fig. 3 shows the depth-dependent gradient microstructure of the SDL. The top-most layer mainly consists of equiaxed nano-grains, as confirmed by TEM observation. Thus, the top-most layer is called the nano-grained layer (NGL). Below the NGL, there are many lamellar structures produced by the SMGT process. Closer to the surface, more distorted grains can be observed. Fig. 3b shows the grains and lamellar structures embedded in severely distorted grains. The lamellar structures are entangled with numerous defects. For depths larger than 200 lm, the shapes and sizes of grains remain unchanged, though many lamellar structures exist (as shown in Fig. 3a and c). In addition, Fig. 3c shows that two groups of crossed lamellar structures can form in an austenitic grain. To investigate the SDLs, a series of TEM samples at different depths were examined. Fig. 4a shows that there
Fig. 3. Microstructure of the SMGT samples. (a) Cross-sectional SEM image and (b and c) showing EBSD inverse pole image of the SDL at depths of 45 lm and 250 lm, respectively.
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Fig. 4. TEM images of the nano-grained layer (NGL). (a) Dark-field TEM image of the NGL at a depth of 25 lm, the inset is the corresponding selected area electron diffraction patterns (SAEDP) and (b) bright-field TEM image of nanotwins at a depth of 25 lm, the inset shows a high resolution TEM image of the area marked by the arrow in Fig. 4b.
are a great number of equiaxed nano-grains in the NGL. Some grains contain nano-twins, as shown in Fig. 4b. Statistical analysis of large numbers of TEM images indicates that the average size of the nano-grains at a depth of 25 lm is approximately 100 nm. The average grain size increases with depth. TEM observation shows that the lamellar structures below the NGL are mainly DTs and austenitic laths divided by DTs (as shown in Fig. 5a), and no martensitic laths can be found. Because the DTs are very thin (less than 20 nm in thickness) and contain large numbers of dislocations that cannot be detected by EBSD/SEM. Some austenitic laths form secondary DTs (as shown in Fig. 5b). Moreover, the interaction
between the primary DTs and the secondary DTs causes the austenitic laths to twist, forming dislocation walls, as shown in Fig. 5c and d. At a depth of 600 lm, the number of nanotwinned bundles decreases drastically, but the interaction between twins and dislocations is also observed. Because no martensite can be found using TEM and EBSD, XRD was employed to evaluate the phase structure of the SDLs. Fig. 6 shows the XRD patterns of the SDLs at different depth. Except for the occurrence of the austenitic (c) peaks, no martensitic peaks can be found in any of the SDLs, so no martensitic transformation occurs in the process of SMGT.
Fig. 5. TEM images of a nano-twinned layer (NTL). (a) TEM morphology at a depth of 200 lm, and insert figure is the SAEDP of the twinned bundles; (b) the austenite lath contains the secondary DTs at a depth of 200 lm; (c) TEM image of a distorted lath contain the secondary twins and a dislocation wall at a depth of 200 lm; (d) dark-field TEM image of (c) and its SAEDP, showing approximately 6.5° misorientation between the ‘‘A’’ and ‘‘B’’ areas. The SAEDPs in (a) and (d) are acquired along the [0 1 1]c zone axis.
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Fig. 6. XRD patterns of the SDLs at different depths.
3.2.2. SMGT samples annealed at different temperature Because the hardness of the annealed SDLs is still very high, the microstructures of the SDLs annealed at different temperature were examined in detail. Fig. 7a–d shows cross-sectional SEM images of the SMGT samples annealed for 1 h at 575, 600, 650 and 700 °C, respectively. In the 575 °C-annealed sample, a large number of small grains are arranged along the original twin boundaries (TBs) to form numerous stripes in a plaid pattern, and most of the DTs in the SDLs are still present. With increasing annealing temperature, the number of DTs gradually decreases, and the recrystallized grains gradually grow (Fig. 7b). Fig. 7c shows that a large number of DTs disappear after annealing at 650 °C, and
Fig. 7. Cross-sectional SEM image of the SMGT samples annealed at different annealing temperature. (a) At 575 °C; (b) at 600 °C; (c) at 650 °C and (d) at 700 °C.
the size of the recrystallized grains increases up to several microns. After annealing at 700 °C for 1 h, most DTs disappear, and the recrystallization of the SMGT samples is nearly complete, as shown in Fig. 7d. It is interesting that the nanograins in the top-most layer do not grow easily during thermal annealing, and their grain sizes are smaller than those of recrystallized grains in the sub-layer, as shown in Fig. 7. This finding means that the thermal stability of the NGL is higher than that of the deformation layer consisting of DTs and dislocations. However, the literature [24] reports that the thermal stability of nanograins is the lowest, with respect to that of DTs and dislocations. In this work, the high thermal stability of the NGLs is attributed to their low dislocation density and inclusion of embedded nanotwins (Fig. 4c). To understand the microstructure shown in Fig. 7a–c, EBSD was performed intensively. After annealing at 600 °C for 1 h, many small-angle grain boundaries (SAGBs) are evident either perpendicular or parallel to the original TBs (as shown in Fig. 8a and c), and the deformation twinned stripes are cut into many chain-like twinned cells (as shown in Fig. 8a and b). Some SAGBs are formed by dislocation walls generated during the SMGT process, while other SAGBs are attributed to dislocation reactions. The chain-like twinned cells must arise during the recovery of the primary DTs, companied by the interaction between dislocations and TBs. This interaction between dislocations and TBs causes DTs to disappear [28,29], and is consequently considered as de-twinning induced by the thermal activation. De-twinning not only reduces the volume of DTs but also eliminates dislocations. The areas around chain-like twinned cells are clean, indicating the low dislocation-density zones. The clean areas around chain-like twinned cells inhibit further de-twinning and therefore enhance the thermal stability of the DTs. These results illustrate that de-twinning must occur in the recovery of the DTs during high temperature annealing. The de-twinning plays a key role in the recovery of the DTs and determines the thermal stability of the twinned-microstructure. In addition, the recrystallization must occur within the high dislocation-density zones of the SMGT samples. Fig. 8a also shows the formation of many recrystallized grains in the lower right corner. 3.2.3. Cold-rolled and its annealed samples To understand the different deformation processes, the microstructures of the CR and its annealed samples were also investigated. The microstructure of the CR sample contains elongated and distorted grains and lamellar structures, as shown in Fig. 9a. TEM observations illustrate that the CR sample contains many of a0 -martensitic laths, distorted austenitic laths and small number of DTs, as shown in Fig. 9b and c. When annealing the CR samples at 450 °C for 1 h, some of the lamellar structures begin to recover. With the temperature increasing, the lamellar structures decrease, and the recrystallized grains gradually nucleate and grow. Fig. 9d shows the typical SEM morphology of the partially recrystallized CR sample, which contains both lamellar structures and recrystallization zones. TEM observations show that the DTs are more difficult to recover by annealing as compared to the a0 -martensite. Fig. 9f shows that the DTs do not change in the partially recrystallized CR sample, though recrystallized grains do appear around the DTs. After annealing at 600 °C for 1 h, the lamellar structures nearly disappear and recrystallization of the CR samples is nearly complete (Fig. 9e). The XRD analysis reveals the phase evolution of the CR samples during annealing, as shown in Fig. 10. The CR samples consist of many of a0 -martensite and austenite (c). The intensity of the a0 -martensitic peaks begins to diminish after annealing at 450 °C for 1 h and continuously decreases with temperature. After annealing at 600 °C for 1 h, a small number of a0 -martensitic peaks can be found. This result also indicates that the a0 -martensite formed by deformation is
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Fig. 8. EBSD analysis of the SMGT samples after annealing at 600 °C for 1 h. (a) EBSD inverse pole image of the SDL (at a depth of approximately 150 lm), the red and black lines represent high-angle and low-angle boundaries, respectively; (b) misorientation along the line ‘‘1–2’’, showing the twinning relationship and (c) misorientation along the line ‘‘3–4’’, showing the misorientation of low-angle boundaries. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 9. The microstructure of the CR samples with different annealing temperature for 1 h. (a) SEM image of the CR sample; (b) TEM image of a0 -martensitic lathes in the CR samples; (c) TEM image of DTs in the CR samples, and the SAEDP is acquired along the [0 1 1]c zone axis; (d) SEM image of the 550 °C-annealed CR samples, the labels ‘‘L’’ and ‘‘R’’ represent the lamellar structures and the recrystallization zones, respectively; (e) SEM image of the 600 °C-annealed CR samples and (f) TEM image of DTs in the partially recrystallized CR samples.
unstable, and it quickly undergoes a reverse transformation to austenite during annealing. 3.2.4. Thermal stability of the deformation twins and the a0 -martensite The volume fraction of the a0 -martensite (Va0 –M) was measured by XRD, and that of grains containing many of DTs (VDT) was statistically measured using ECCI in the SEM. Fig. 11 shows the Va0 –M of the CR samples and the VDT in the sublayer of the SMGT samples vs. annealing temperature. The Va0 –M of the CR samples is
approximately 52%, and decreases to 40% after annealing at 450 °C for 1 h. After annealing at 600 °C for 1 h, the Va0 –M decreases to approximately 5.8%, whereas the VDT in the sublayer of the SMGT samples begins to decrease. When the annealing temperature exceeds 625 °C, the VDT rapidly decreases. After annealing at 700 °C for 1 h, the VDT decreases to approximately 3.2%. This result further confirms that the DTs are more stable than the a0 -martensite. The full recrystallization of the SMGT samples occurs at 700 °C, while that of the CR samples takes place at 600 °C.
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Fig. 10. XRD patterns of the CR samples with different annealing temperature for 1 h.
Fig. 11. The volume fraction of deformation twins (VDT) in the sublayer of the SMGT samples and that of a0 -martensite (Va0 –M) of the CR samples as a function of annealing temperature.
3.3. Discussion As shown above, for all SMGT samples, the surface hardness of the top-most layer is as high as 500 ± 20 HV, which is not affected by the preset indentation depth (d). However, the thickness of the hardened layer increases with increasing d. The mechanical properties depend on the microstructure of the SDL formed during the SMGT process. The top-most layers of all SMGT samples consist of equiaxed nano-grains, some of which contain nano-twins. The surface hardness of the SMGT samples with a similar NGL must be much the same. Below the NGLs, there are many distorted grains, deformation twins with nano-thickness and dislocation structures. The number of these distorted grains, DTs and dislocation structures in the SDLs decrease with increasing the depth, so that the hardness of the SDLs changes correspondingly. Besides, the thickness of the gradient deformation layer is strongly affected by the preset indentation depth (d). In the present work, the thickness of the SDL is approximately 300, 330 and 900 lm for d = 15, 20 and 32.5 lm, correspondingly. No martensite can be found in the SMGT samples, but the martensite is the main component in the CR samples in which the TRIP effect is dominant. In the SMGT process, a very high-rate shear deformation is applied to the surface layer of the samples, where the estimated strain rate varies from 103–104 (at the top-most surface) to 102–103 (80 lm below the treated surface) [30]. Meanwhile the cumulative equivalent strains of the SMGT are approximately 15–30, which is much higher than that of cold rolling (approximately 0.65). High strain rate promotes
deformation twinning [10,29,31], so the SDLs of the SMGT samples contain large numbers of DTs. Moreover, the literature [11] reported that the temperature rise due to deformation heating will increase with the increases of strain rate and strain in Fe–23Mn–2Al–0.2C TWIP steel. The literature [32,33] reported that temperature increases as high as 450 °C at the shear band center when carbon steels are tested at strain rates in the range of 1000–5000 s1. When materials undergo high-strain-rate deformation, especially for cumulative deformation such as the SMGT in this work, large amounts of plastic work are performed in a very short time, which leads to an increase in temperature of the deformation zones. The samples were cooled by a water-based coolant, but we suspect that the temperature of the SDLs might increase instantaneously at the moment of high-rate shear deformation, after which the SDLs are quickly cooled by the applied water. Some phenomena also indicate the temperature rise due to deformation heating. For example, the low dislocation density and equiaxed nanograins in the NGL should be attributed to elevated temperature and dynamic recrystallization. According to the literature [6,20,21], the SFE value increases markedly with elevated temperature. Here, the SFE value of the SDLs probably increases instantaneously, which leads to the TWIP effect and suppresses the TRIP effect. In this work, the thermal stability of deformed samples is studied. The hardness of the CR samples, which contain a lot of a0 -martensite laths and elongated austenitic grains, is about 405 ± 8 HV. However, the a0 -martensite is relatively unstable. The Va0 –M of the CR samples quickly decreases from 40% to 5.8% when the annealing temperature increases from 450 to 600 °C, and the hardness sharply decreases from 400 to 240 HV correspondingly. It is worthy of noting that the hardness decrease of the SMGT samples is not obvious during annealing below 575 °C, and the hardness of the outer layer annealed at 625 °C for 1 h is more than 400 HV. The outstanding thermal stabilities of the GNTLs are ascribed to the DTs with good thermal stability. The DTs in the interior of coarse austenitic grains were investigated. In the SMGT process, some coarse austenite grains can forms many of multiple deformation twins, such as primary DTs and secondary DTs. These primary DTs and secondary DTs present two groups of crossed lamellar structures, as shown in Fig. 3c. Dislocations emitted from the secondary DTs must interact with the primary DTs, and leave many partial dislocations and stair-rod dislocations at the TBs [29]. When the SMGT sample is heated to a high temperature, these partial dislocations and stair-rod dislocations are activated and slip on the TBs, decreasing the thickness of the twins. De-twinning is easy at the beginning of the recovery of crossed DTs embedded in the interior of grains, which lowers the thermal stability of the DTs. When the deformation twinned stripes are cut into many chain-like twinned cells by de-twinning, the dislocation density must decrease, and de-twinning becomes difficult, which improves the thermal stability of DTs. This is to say that de-twinning plays an important role in the recovery of the DTs embedded in grains, and affects the thermal stability of DTs. In a word, ultrahard and ultrastable GNTLs can be generated on high manganese TRIP steels by means of a high strain-rate SMGT processing. Compared to conventional surface technique, the SMGT processing is simple and cost-effective. It is believed that the SMGT processing will become a new surface strengthening technique in industrial manufacturing processes.
4. Conclusion In summary, the gradient nanotwinned layers that contain many nanotwins and nanograins were produced on a Fe–20Mn–3
P. Xie et al. / Materials and Design 80 (2015) 144–151
Al–3Si TRIP steel by means of surface mechanical grinding treatment (SMGT). From the obtained results, the following can be concluded. (1) The gradient nanotwinned layers have high hardness, and the surface hardness can reach a value up to 500 HV. The hardness decreases from 500 to 230 HV with increasing depth, where the latter value is actually the hardness of the matrix. The thickness of the hardened layer increases with increasing the preset indentation depth d, and the thickness of the hardened layer can reach up to 900 lm at d = 32.5 lm. (2) In the SMGT process, a very high-rate shear deformation is applied to the surface layer of the samples, which promotes deformation twinning and suppresses martensitic transformation. (3) The martensite induced by deformation is unstable. The martensite content of the cold-rolled samples can be reduced from 52% to 40% after annealing at 450 °C for 1 h, and remains only by 5.8% after annealing at 600 °C for 1 h, indicating that the recrystallization in the annealed cold-rolled samples is nearly complete. (4) De-twinning, occurring during thermal annealing, plays the key role in the recovery of deformation twins embedded in grains and affects the thermal stability of deformation twins. The thermal stability of deformation twins embedded in grains is higher than that of martensite induced by deformation, but lower than that of equiaxial nano-grains in the nano-grained layers. For the SMGT samples, the volume fraction of the deformation twins begins to decrease at 600 °C for 1 h, and the hardness of the top-most layers can remain at a value above 400 HV even after annealing at 625 °C for 1 h, whereas the hardness of the cold-rolled counterpart samples decreases drastically to less than 240 HV.
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