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α to γ transformation in the nanostructured surface layer of pearlitic steels near room temperature
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Scripta Materialia 59 (2008) 806–809 www.elsevier.com/locate/scriptamat
a to c transformation in the nanostructured surface layer of pearlitic steels near room temperature Na Min,a,b Wei Lia and Xuejun Jina,* a
School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, China b School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China Received 27 March 2008; revised 26 May 2008; accepted 28 May 2008 Available online 7 June 2008
An unusual low-temperature structural transformation from ferrite (a) to austenite (c) is observed near room temperature through surface mechanical attrition treatment (SMAT) by means of transmission electron microscopy (TEM) and back-scattering Mo¨ssbauer spectroscopy (BSMS). Associated with the thermodynamic calculation, two factors that lead to the appearance of austenite after SMAT near room temperature are discussed. One factor is the formation of supersaturated nanocrystalline ferrite. The other is the compression stress at the boundaries of nano-size grains after SMAT. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Pearlitic steels; Surface mechanical attrition treatment; Nanocrystalline; a ? c transformation
Properties of nanocrystalline materials have aroused much interest over the past decade, including high hardness and strength, enhanced physical properties and improved tribological properties, which are fundamentally different from their conventional coarsegrained polycrystalline counterparts [1,2]. So far various methods have been developed to obtain nanocrystalline materials such as equal-channel angular pressing (ECAP), high-pressure torsion (HPT) and cold rolling [3–5]. Hence, a better understanding of the mechanism of strain-induced grain refinement is of academic significance. It has recently been reported that a transformation of body-centered cubic (bcc) ferrite to facecentered cubic (fcc) austenite occurred in pearlitic steels by high-pressure torsion (HPT) through the observation of high-resolution transmission electron microscopy (HRTEM) [6–8]. The following factors may affect the phase equilibrium between parent and new phase: chemical composition, hydrostatic stress, defects in the microstructure of parent phase [9]. Thus, the observations of a ? c transformation phenomena were attributed to the nanostructure and shear stress [8]. However, it is well known that the strain can induce the dissolution of iron carbide with small dimensions under severe plastic deformation [10–12]. Carbon concentration in the ferrite matrix increases due to the dissolution of cementite lamellae.
* Corresponding author. Tel.:/fax: +86 21 54745560; e-mail: [email protected]
Some experimental observations have proved that 1.0– 1.7 at.% carbon concentration supersaturated nanoferrite appeared in pearlitic steel under heavily cold drawn and ball milling [13–15]. Moreover, it has been noted that the carbon concentration depended on the dimension of ferrite phase; the carbon concentration was higher in thinner ferrite [13]. The increase of carbon concentration in the nanocrystalline ferrite phase due to the dissolution of cementite under severe plastic deformation (SPD) was not taken into account in the previous discussion. Surface mechanical attrition treatment (SMAT) is a recently developed technique that can provide a unique opportunity to investigate grain refinement process by severe plastic deformation [16,17]. Much experimental evidence has shown that extensive plastic deformation with high strain rates, such as SMAT, may refine grains to the nanometer scale (5–30 nm) [18,19]. So in this work, SMAT is applied to study the phase transformation process of nanocrystalline multiphase system with consideration of supersaturated carbon in nanoscale. The present work characterized the microstructure evolution of a pearlitic steel subjected to SMAT and discussed phase stability after severe plastic deformation with the aid of thermodynamic calculation. A Fe–0.80 wt.%C steel containing 0.61 wt.% of Mn, 0.28 wt.% of Si, 0.15 wt.%V with Ni, Cr, S, Al and N less than 0.01 wt.% was used in this study. Vanadium was added to raise the recrystallization temperature. Prior to the SMAT, the sample was austenitized at 1223 K for 15 min, then quenching into a salt bath to
1359-6462/$ - see front matter Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2008.05.038
N. Min et al. / Scripta Materialia 59 (2008) 806–809
a temperature of 843 K for 25 min. In the present work, these samples for SMAT were treated at a vibrating frequency of 50 Hz with spherical hardened steel balls under vacuum at room temperature for 15 min and 60 min. The hardened steel balls have a diameter of 8 mm and a measured composition (wt.%) of 12.25Cr, a minor amount of C and balance Fe. The back-scattering Mo¨ssbauer spectra (BSMS) were measured at room temperature in the back-scattering acceleration geometry mode with a 57Co(Pd) source. A least-squares method was used to fit the Mo¨ssbauer spectrum. Under the selected measurement conditions, microstructure information within 100 nm thickness from the top surface can be obtained [20]. The microstructures of the untreated and SMAT treated samples with their surface layer were observed by transmission electron microscopy (TEM) using a Hitachi H-800 TEM and JEM 2100F TEM respectively. TEM specimens were mechanical thinned from the untreated side to 50 lm and hen final thinning to electron transparency by argon ion milling. In the case of BSMS, the spectrum stems from a region of 10–100 nm below the surface of the samples, and thus becomes an effect tools to provide valuable information on the composition details of thin surface layers [21]. According to the different magnetic properties of austenite and ferrite phases. BSMS results could reveal the microstructure evolution in steels which is shown in Figure 1. The hyperfine interaction parameters of the spectra of the samples derived from the least square method are listed in Table 1. The Mo¨ssbauer spectra of untreated samples consist of two sextets shown by Table 1: A sextet of a-Fe corresponds to an effective field of 331 kOe [22]; B sextet with an effective field of 306 kOe is attributed to cementite phase, which deviates a bit from the hyperfine interaction parameter reported
Fig. 1. Back-scattering Mo¨ssbauer spectrum of untreated and SMAT treated pearlitic steels: (a) untreated samples, (b) samples after SMAT treated for 60 min.
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Table 1. The Mo¨ssbauer parameters of the surface layer of pearlitic steels before and after SMAT for 60 min Sample
H (kOe) I.S. Q.S. C/2 Area (mm s1) (mm s1) (mm s1)
Untreated A 331.11 sample B 306.49 SMAT A 330.78 treated sample Singlet Doublet
0
0
0.15
0.888
0.15 0
0.36 0
0.21 0.17
0.112 0.786
0.1 1.06
1.41
0.24 0.37
0.066 0.149
by Refs. [23,24]. It seems that alloy elements have a slight influence on the hyperfine field of cementite phase [25]. Based on the area of line width, the relative volume fraction of cementite is estimated to be 11.2% (±2%). Besides the ferromagnetic sextet that corresponds to the ferrite phase, surprisingly the singlet and doublet spectrum both appear in the spectra of the SMAT treated samples indicating the presence of paramagnetic phases. One of the paramagnetic components of the spectra is related to fcc austenite [26]. According to different surrounding conditions, the sites of Fe atoms can be divided into three different types: (1) Fe atoms without near-neighbor or next-near carbon atoms, (2) Fe atoms with only one near-neighbor carbon atoms, and (3) Fe atoms without near-neighbor carbon atoms but with next-near-neighbor carbon atoms [27]. The singlet of austenite phase in SMAT-treated samples corresponds to the first and third interactions and suggests that austenite phase with high carbon concentration may be evolved during severe plastic deformation. The relative volume fraction of austenite is estimated to 6.6% (±2%). However, the sextet corresponding to the cementite phase does not appear. Figure 2 shows a typical TEM image of the full pearlitic microstructure before being SMAT-treated, where the pearlitic lamellar is apparent and few dislocation is observed within the ferrite phase. After SMAT for 60 min, the surface of samples experience severe plastic deformation. The cross-sectional optical micrograph of the steel sample (Fig. 3) shows that the severe plastic deformation occurs in the surface layer of approximately 25 lm. Different from the untreated samples, the microstructure of the sample under SMAT for 15 min which endured very high density of dislocations appeared in the ferrite phase, and a cellular structure which formed in the ferrite and the cementite was bent and fragmented, as shown in Figure 4. Figure 5 shows the microstructure within the topmost layer of the sample after SMAT for 60 min. The microstructure of the SMAT-treated steel
Fig. 2. TEM image of the full pearlite in pearlitic steels.
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N. Min et al. / Scripta Materialia 59 (2008) 806–809
is homogenous with nano-sized grain instead of typical lamellae structure of pearlitic colonies, as shown in the bright field image (Fig. 5a). The dark-field image of the ferrite produced using electrons from the 110 ring is shown in Figure 5b. This shows that the top surface layer consists of roughly equiaxed nanocrystalline ferrite estimated to be 10–20 nm. The selected area electron diffraction (SAED) ring shown in Figure 5c indicates the strong ferrite rings characteristic of polycrystalline structure. There are also some reflections that could be interpreted as austenite reflections (2 0 0c, 2 2 0c, 3 1 1c), as shown in Figure 5c. This further verified the results of Mo¨ssbauer spectroscopy. This phenomenon has not been observed at the steel with speroidal cemenetite subjected to SMAT [28]. The morphology of cementite (spherical or lamellae) is expected to have a significant impact on the mechanism of the strain-induced structure evolution during SMAT. Some of spots due to cementite phase can be observed, different from the result of Mo¨ssbauer spectroscopy. This may be due to the fact that the TEM observation was made in localized area, while BSMS was measured the globe result of the sample, which cannot reflect the phase with average volume composition less than 1%. In addition, there are some reflections that can be identified as graphite. This phenomenon is also mentioned in the previous work [29] and traced through sensitive Raman observation [30]. Previous studies [17–19] already indicated that due to the high vibration frequency of the system in SMAT, re-
Fig. 3. Cross-sectional optical micrograph of the pearlitic steels after SMAT treated for 60 min.
Fig. 4. TEM image of the microstructure of pearlitic steels after SMAT for 15 min.
Fig. 5. (a) Bright field of ferrite of pearlitic steels after the SMAT for 60 min,(b) darkfield of g (1 1 0)ferrite, (c) selectedareadiffractionpattern.
peated multidirectional peening at high strain rates onto the sample surface leads to severe plastic deformation in the surface layer and finally make the surface of samples nanocrystallized. During SMAT the multi-directional deformation is greatly beneficial to the production of high dislocation densities and therefore dislocation cells which refine the ferrite lamellae into nanograins. With the increase of strain, dislocation cells developed in the ferrite and cell boundaries trap more and more carbon atoms [31]. Grain boundary behavior in nanocrystalline materials cannot be ignored since the volume fraction of grain boundaries is comparable with that of regular crystal lattice [32]. The segregation of carbon atoms at the grain boundaries of nanocrystalline ferrite makes the supersaturated average carbon concentration of nanocrystalline ferrite even up to about 1.0–1.7 at.% [13–15] detected. Our previous work has also reported that the average carbon concentration in the ferrite was measured about 1.0 at.% by means of 3DAP [33]. Because of the non-uniform characterization of the deformation during SMAT, the strain rate level is estimated to be of the order of 104 s1 with a true strain of 4 or 5 in the localized region [34]. And SMAT leads to compressive residual stress with important maximum compressive stress at the surface. The maximum value reaches about 1000 MPa [35]. The compressive residual stress of about 1 GPa for the nano-sized grains of ferrite would make the thermodynamic equilibrium temperature decrease [36]. Therefore, the following will discuss the thermodynamic equilibrium conditions for a ? c transformation based on the above analysis. The amount of energy change per unit volume of supersaturated ferrite under the compressive residual stress can be expressed as [9]
N. Min et al. / Scripta Materialia 59 (2008) 806–809
DG
a!c
c
¼ G þ DGP G
a
where DGP = e r [36] is the additional energy change per unit volume arising from the compressive residual stress, and e is the volume strain component of the a ? c transformation. In the present work, e = 0.03 (a negative value means that the a ? c transformation is accompanied with a decrease in volume) and r is a hydrostatic stress which is a negative value because of the compressive residual stress. Thus, it is estimated that GP = 214 J mol1, when the applied maximum compressive stress for nanocrystalline ferrite is estimated to be 1 GPa. Ga and Gc are the chemical free energy of a and c phase with the carbon concentration Xc at a fixed temperature T (K) [37,38], respectively. The equilibrium temperature T for the ferrite of various compositions Xc, which is calculated fromDGa?c = 0, is shown in Figure 6, which indicates that the equilibrium temperature for transformation decreases with the increase of the carbon concentration in ferrite phase. With the carbon concentration in ferrite to 1.0 at.%, the temperature for the equilibrium is 579 K.When the carbon concentration is up to 1.5 at.%, it can even reach room temperature. At the same time, the heating effect has been reported by the existence of adiabatic bands in certain heavily deformed alloys and a local temperature rise up to a few hundred degrees celsius [39]. Therefore, a ? c transformation from ferrite with supersaturated carbon to austenite can occur at the adiabatic temperature of 600 K, and even lower temperature under SMAT. On the basis of the thermodynamic calculation of Gibbs free energy of c and a phase with respect to the concentration of carbon, it is concluded that two factors lead to the appearance of austenite after severe plastic deformation near room temperature: (1) the existence of a large number of carbon atoms supersaturated in ferrite and the increase of the Gibbs free energy of ferrite due to the formation of the nanocrystalline structure; (2) the effect of compression stress at the boundaries of nano-size grains after SMAT further decreases the difference of free energy between two phases, and makes free energy of austenite be lower than that of ferrite near room temperature. The present work was financially supported by the National Natural Science Foundation of China (Nos. 50371057, 50571064). The authors would like to thank Dr. D.M. Jiang from East China Normal University
Fig. 6. The equilibrium temperature T for the ferrite of various compositions Xc, which is calculated from DGa?c = 0.
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for performing the BSMS investigations and Dr. X.D. Wang for the help in TEM experiments. Thanks also to Prof. Q.P. Meng for helpful discussion. [1] [2] [3] [4]
H. Gleiter, Prog. Mater. Sci. 33 (1989) 223. H. Gleiter, Acta Mater. 48 (2000) 1. R.Z. Valiev, Mater. Sci. Eng. A 234–236 (1997) 59. D.H. Shin, I. Kim, J. Kim, Y.S. Kim, S.L. Semiatin, Acta Mater. 51 (2003) 983. [5] J.S.C. Jang, C.C. Koch, Scripta Mater. 24 (1990) 1590. [6] Yu. Ivanisenko, I. MacLaren, R.Z. Valiev, H.-J. Fecht, Adv. Eng. Mater. 11 (2005) 1011. [7] Yu. Ivanisenko, I. MacLaren, X. Sauvage, R.Z. Valiev, H.-J. Fecht, Solid State Phenom. 114 (2006) 133. [8] Yu. Invaniseko, I. Maclaren, X. Sauvage, R.Z. Valiev, H.-J. Fecht, Acta Mater. 54 (2006) 1659. [9] E. Hornbogen, Acta Metall. 33 (1985) 595. [10] V.N. Gridnev, V.V. Nemoshkalenko, Y.Y. Meshkov, V.G. Gavrilyuk, V.G. Prokopenko, O.N. Razumov, Phys. Status Solidi (a) 201 (1975) 31. [11] Y. Ivanisenko, W. Lojkowsky, R.Z. Valiez, H.J. Fecht, Acta Mater. 51 (2003) 5555. [12] J. Languillaume, G. Kapelski, B. Baudelet, Acta Mater. 45 (1997) 1201. [13] N. Maruyama, T. Tarui, H. Tashiro, Scripta Mater. 46 (2002) 599. [14] K. Hono, M. Ohnuma, M. Murayama, S. Nishida, A. Yoshie, Scripta Mater. 44 (2001) 977. [15] M. Peet, S.S. Babu, M.K. Miller, H.K.D.H. Bhadeshia, Scripta Mater. 50 (2004) 1277. [16] K. Lu, Mater. Sci. Eng. R 16 (1996) 161. [17] K. Lu, J. Lu, Mater. Sci. Eng. A 375 (2004) 38. [18] N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603. [19] N.R. Tao, M.L. Sui, J. Lu, K. Lu, Nanostruct. Mater. 11 (1999) 433. [20] Z.C. Ni, X.W. Wang, J.Y. Wang, E. Wu, Physica B 334 (2003) 221. [21] F. Nagy, Z. Klencsar, Nucl. Instrum. Meth. Phys. Res. B 245 (2006) 528. [22] Y.H. Zhao, H.W. Sheng, K. Lu, Acta Mater. 49 (2001) 365. [23] V.N. Gridnev, V.G. Gavriljuk, I.Y. Dekhtyar, Y.Y. Meshkov, V.G. Prokopenko, P.S. Nizin, Phys. Status Solidi (a) 14 (1972) 689. [24] V.A. Shabashov, L.G. Korshunov, A.G. Mukoseev, V.V. Sagaradze, A.V. Makarov, V.P. Pilyugin, S.I. Novikov, N.F. Vidanova, Mater. Sci. Eng. A 346 (2003) 196. [25] V.G. Gavrilyuk, V.G. Prokopenko, O.N. Razumov, Phys. Status Solidi (a) 53 (1979) 147. [26] J. Ding, H. Huang, P.G. McCormick, R. Street, J. Magn. Magn. Mater. 139 (1995) 109. [27] J. Desimoni, R. Gregorutti, K. Laneri, J.L. Sarutti, R.C. Mercader, Metall. Mater. Trans. A 30 (1999) 2745. [28] L. Zhou, G. Liu, X.L. Ma, K. Lu, Acta Mater. 56 (2008) 78. [29] X.P. Yong, G. Liu, J. Lu, K. Lu, Acta Metall. Sin. 38 (2002) 157. [30] N. Min, W. Li, X.J. Jin, in press. [31] X. Sauvage, Y. Inanisenko, J. Mater. Sci. 42 (2007) 1615. [32] Q.P. Meng, Y.H. Rong, T.Y. Hsu, Mater. Sci. Eng. A 471 (2007) 22. [33] N. Min, W. Li, H.B. Wang, X.J. Jin, Chin. Sci. Bull., in press. [34] Z.G. Liu, H.J. Fecht, M. Umemoto, Mater. Sci. Eng. A 375–377 (2004) 839. [35] T. Roland, D. Retraint, K. Lu, J. Lu, Scripta Mater. 54 (2006) 1949. [36] J.R. Patel, M. Cohen, Acta Metall. 1 (1953) 531. [37] T.Y. Hsu, L. Li (Eds.), Thermodynamics of Materials, Science Publications, Beijing, BJ, 2005. [38] T.Y. Hsu, H.B. Chang, Acta Metall. 32 (1984) 343. [39] J. Duffy, Y.C. Chi, Mater. Sci. Eng. A 517A (1992) 195.
Scripta Materialia 59 (2008) 806–809 www.elsevier.com/locate/scriptamat
a to c transformation in the nanostructured surface layer of pearlitic steels near room temperature Na Min,a,b Wei Lia and Xuejun Jina,* a
School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, China b School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China Received 27 March 2008; revised 26 May 2008; accepted 28 May 2008 Available online 7 June 2008
An unusual low-temperature structural transformation from ferrite (a) to austenite (c) is observed near room temperature through surface mechanical attrition treatment (SMAT) by means of transmission electron microscopy (TEM) and back-scattering Mo¨ssbauer spectroscopy (BSMS). Associated with the thermodynamic calculation, two factors that lead to the appearance of austenite after SMAT near room temperature are discussed. One factor is the formation of supersaturated nanocrystalline ferrite. The other is the compression stress at the boundaries of nano-size grains after SMAT. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Pearlitic steels; Surface mechanical attrition treatment; Nanocrystalline; a ? c transformation
Properties of nanocrystalline materials have aroused much interest over the past decade, including high hardness and strength, enhanced physical properties and improved tribological properties, which are fundamentally different from their conventional coarsegrained polycrystalline counterparts [1,2]. So far various methods have been developed to obtain nanocrystalline materials such as equal-channel angular pressing (ECAP), high-pressure torsion (HPT) and cold rolling [3–5]. Hence, a better understanding of the mechanism of strain-induced grain refinement is of academic significance. It has recently been reported that a transformation of body-centered cubic (bcc) ferrite to facecentered cubic (fcc) austenite occurred in pearlitic steels by high-pressure torsion (HPT) through the observation of high-resolution transmission electron microscopy (HRTEM) [6–8]. The following factors may affect the phase equilibrium between parent and new phase: chemical composition, hydrostatic stress, defects in the microstructure of parent phase [9]. Thus, the observations of a ? c transformation phenomena were attributed to the nanostructure and shear stress [8]. However, it is well known that the strain can induce the dissolution of iron carbide with small dimensions under severe plastic deformation [10–12]. Carbon concentration in the ferrite matrix increases due to the dissolution of cementite lamellae.
* Corresponding author. Tel.:/fax: +86 21 54745560; e-mail: [email protected]
Some experimental observations have proved that 1.0– 1.7 at.% carbon concentration supersaturated nanoferrite appeared in pearlitic steel under heavily cold drawn and ball milling [13–15]. Moreover, it has been noted that the carbon concentration depended on the dimension of ferrite phase; the carbon concentration was higher in thinner ferrite [13]. The increase of carbon concentration in the nanocrystalline ferrite phase due to the dissolution of cementite under severe plastic deformation (SPD) was not taken into account in the previous discussion. Surface mechanical attrition treatment (SMAT) is a recently developed technique that can provide a unique opportunity to investigate grain refinement process by severe plastic deformation [16,17]. Much experimental evidence has shown that extensive plastic deformation with high strain rates, such as SMAT, may refine grains to the nanometer scale (5–30 nm) [18,19]. So in this work, SMAT is applied to study the phase transformation process of nanocrystalline multiphase system with consideration of supersaturated carbon in nanoscale. The present work characterized the microstructure evolution of a pearlitic steel subjected to SMAT and discussed phase stability after severe plastic deformation with the aid of thermodynamic calculation. A Fe–0.80 wt.%C steel containing 0.61 wt.% of Mn, 0.28 wt.% of Si, 0.15 wt.%V with Ni, Cr, S, Al and N less than 0.01 wt.% was used in this study. Vanadium was added to raise the recrystallization temperature. Prior to the SMAT, the sample was austenitized at 1223 K for 15 min, then quenching into a salt bath to
1359-6462/$ - see front matter Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2008.05.038
N. Min et al. / Scripta Materialia 59 (2008) 806–809
a temperature of 843 K for 25 min. In the present work, these samples for SMAT were treated at a vibrating frequency of 50 Hz with spherical hardened steel balls under vacuum at room temperature for 15 min and 60 min. The hardened steel balls have a diameter of 8 mm and a measured composition (wt.%) of 12.25Cr, a minor amount of C and balance Fe. The back-scattering Mo¨ssbauer spectra (BSMS) were measured at room temperature in the back-scattering acceleration geometry mode with a 57Co(Pd) source. A least-squares method was used to fit the Mo¨ssbauer spectrum. Under the selected measurement conditions, microstructure information within 100 nm thickness from the top surface can be obtained [20]. The microstructures of the untreated and SMAT treated samples with their surface layer were observed by transmission electron microscopy (TEM) using a Hitachi H-800 TEM and JEM 2100F TEM respectively. TEM specimens were mechanical thinned from the untreated side to 50 lm and hen final thinning to electron transparency by argon ion milling. In the case of BSMS, the spectrum stems from a region of 10–100 nm below the surface of the samples, and thus becomes an effect tools to provide valuable information on the composition details of thin surface layers [21]. According to the different magnetic properties of austenite and ferrite phases. BSMS results could reveal the microstructure evolution in steels which is shown in Figure 1. The hyperfine interaction parameters of the spectra of the samples derived from the least square method are listed in Table 1. The Mo¨ssbauer spectra of untreated samples consist of two sextets shown by Table 1: A sextet of a-Fe corresponds to an effective field of 331 kOe [22]; B sextet with an effective field of 306 kOe is attributed to cementite phase, which deviates a bit from the hyperfine interaction parameter reported
Fig. 1. Back-scattering Mo¨ssbauer spectrum of untreated and SMAT treated pearlitic steels: (a) untreated samples, (b) samples after SMAT treated for 60 min.
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Table 1. The Mo¨ssbauer parameters of the surface layer of pearlitic steels before and after SMAT for 60 min Sample
H (kOe) I.S. Q.S. C/2 Area (mm s1) (mm s1) (mm s1)
Untreated A 331.11 sample B 306.49 SMAT A 330.78 treated sample Singlet Doublet
0
0
0.15
0.888
0.15 0
0.36 0
0.21 0.17
0.112 0.786
0.1 1.06
1.41
0.24 0.37
0.066 0.149
by Refs. [23,24]. It seems that alloy elements have a slight influence on the hyperfine field of cementite phase [25]. Based on the area of line width, the relative volume fraction of cementite is estimated to be 11.2% (±2%). Besides the ferromagnetic sextet that corresponds to the ferrite phase, surprisingly the singlet and doublet spectrum both appear in the spectra of the SMAT treated samples indicating the presence of paramagnetic phases. One of the paramagnetic components of the spectra is related to fcc austenite [26]. According to different surrounding conditions, the sites of Fe atoms can be divided into three different types: (1) Fe atoms without near-neighbor or next-near carbon atoms, (2) Fe atoms with only one near-neighbor carbon atoms, and (3) Fe atoms without near-neighbor carbon atoms but with next-near-neighbor carbon atoms [27]. The singlet of austenite phase in SMAT-treated samples corresponds to the first and third interactions and suggests that austenite phase with high carbon concentration may be evolved during severe plastic deformation. The relative volume fraction of austenite is estimated to 6.6% (±2%). However, the sextet corresponding to the cementite phase does not appear. Figure 2 shows a typical TEM image of the full pearlitic microstructure before being SMAT-treated, where the pearlitic lamellar is apparent and few dislocation is observed within the ferrite phase. After SMAT for 60 min, the surface of samples experience severe plastic deformation. The cross-sectional optical micrograph of the steel sample (Fig. 3) shows that the severe plastic deformation occurs in the surface layer of approximately 25 lm. Different from the untreated samples, the microstructure of the sample under SMAT for 15 min which endured very high density of dislocations appeared in the ferrite phase, and a cellular structure which formed in the ferrite and the cementite was bent and fragmented, as shown in Figure 4. Figure 5 shows the microstructure within the topmost layer of the sample after SMAT for 60 min. The microstructure of the SMAT-treated steel
Fig. 2. TEM image of the full pearlite in pearlitic steels.
808
N. Min et al. / Scripta Materialia 59 (2008) 806–809
is homogenous with nano-sized grain instead of typical lamellae structure of pearlitic colonies, as shown in the bright field image (Fig. 5a). The dark-field image of the ferrite produced using electrons from the 110 ring is shown in Figure 5b. This shows that the top surface layer consists of roughly equiaxed nanocrystalline ferrite estimated to be 10–20 nm. The selected area electron diffraction (SAED) ring shown in Figure 5c indicates the strong ferrite rings characteristic of polycrystalline structure. There are also some reflections that could be interpreted as austenite reflections (2 0 0c, 2 2 0c, 3 1 1c), as shown in Figure 5c. This further verified the results of Mo¨ssbauer spectroscopy. This phenomenon has not been observed at the steel with speroidal cemenetite subjected to SMAT [28]. The morphology of cementite (spherical or lamellae) is expected to have a significant impact on the mechanism of the strain-induced structure evolution during SMAT. Some of spots due to cementite phase can be observed, different from the result of Mo¨ssbauer spectroscopy. This may be due to the fact that the TEM observation was made in localized area, while BSMS was measured the globe result of the sample, which cannot reflect the phase with average volume composition less than 1%. In addition, there are some reflections that can be identified as graphite. This phenomenon is also mentioned in the previous work [29] and traced through sensitive Raman observation [30]. Previous studies [17–19] already indicated that due to the high vibration frequency of the system in SMAT, re-
Fig. 3. Cross-sectional optical micrograph of the pearlitic steels after SMAT treated for 60 min.
Fig. 4. TEM image of the microstructure of pearlitic steels after SMAT for 15 min.
Fig. 5. (a) Bright field of ferrite of pearlitic steels after the SMAT for 60 min,(b) darkfield of g (1 1 0)ferrite, (c) selectedareadiffractionpattern.
peated multidirectional peening at high strain rates onto the sample surface leads to severe plastic deformation in the surface layer and finally make the surface of samples nanocrystallized. During SMAT the multi-directional deformation is greatly beneficial to the production of high dislocation densities and therefore dislocation cells which refine the ferrite lamellae into nanograins. With the increase of strain, dislocation cells developed in the ferrite and cell boundaries trap more and more carbon atoms [31]. Grain boundary behavior in nanocrystalline materials cannot be ignored since the volume fraction of grain boundaries is comparable with that of regular crystal lattice [32]. The segregation of carbon atoms at the grain boundaries of nanocrystalline ferrite makes the supersaturated average carbon concentration of nanocrystalline ferrite even up to about 1.0–1.7 at.% [13–15] detected. Our previous work has also reported that the average carbon concentration in the ferrite was measured about 1.0 at.% by means of 3DAP [33]. Because of the non-uniform characterization of the deformation during SMAT, the strain rate level is estimated to be of the order of 104 s1 with a true strain of 4 or 5 in the localized region [34]. And SMAT leads to compressive residual stress with important maximum compressive stress at the surface. The maximum value reaches about 1000 MPa [35]. The compressive residual stress of about 1 GPa for the nano-sized grains of ferrite would make the thermodynamic equilibrium temperature decrease [36]. Therefore, the following will discuss the thermodynamic equilibrium conditions for a ? c transformation based on the above analysis. The amount of energy change per unit volume of supersaturated ferrite under the compressive residual stress can be expressed as [9]
N. Min et al. / Scripta Materialia 59 (2008) 806–809
DG
a!c
c
¼ G þ DGP G
a
where DGP = e r [36] is the additional energy change per unit volume arising from the compressive residual stress, and e is the volume strain component of the a ? c transformation. In the present work, e = 0.03 (a negative value means that the a ? c transformation is accompanied with a decrease in volume) and r is a hydrostatic stress which is a negative value because of the compressive residual stress. Thus, it is estimated that GP = 214 J mol1, when the applied maximum compressive stress for nanocrystalline ferrite is estimated to be 1 GPa. Ga and Gc are the chemical free energy of a and c phase with the carbon concentration Xc at a fixed temperature T (K) [37,38], respectively. The equilibrium temperature T for the ferrite of various compositions Xc, which is calculated fromDGa?c = 0, is shown in Figure 6, which indicates that the equilibrium temperature for transformation decreases with the increase of the carbon concentration in ferrite phase. With the carbon concentration in ferrite to 1.0 at.%, the temperature for the equilibrium is 579 K.When the carbon concentration is up to 1.5 at.%, it can even reach room temperature. At the same time, the heating effect has been reported by the existence of adiabatic bands in certain heavily deformed alloys and a local temperature rise up to a few hundred degrees celsius [39]. Therefore, a ? c transformation from ferrite with supersaturated carbon to austenite can occur at the adiabatic temperature of 600 K, and even lower temperature under SMAT. On the basis of the thermodynamic calculation of Gibbs free energy of c and a phase with respect to the concentration of carbon, it is concluded that two factors lead to the appearance of austenite after severe plastic deformation near room temperature: (1) the existence of a large number of carbon atoms supersaturated in ferrite and the increase of the Gibbs free energy of ferrite due to the formation of the nanocrystalline structure; (2) the effect of compression stress at the boundaries of nano-size grains after SMAT further decreases the difference of free energy between two phases, and makes free energy of austenite be lower than that of ferrite near room temperature. The present work was financially supported by the National Natural Science Foundation of China (Nos. 50371057, 50571064). The authors would like to thank Dr. D.M. Jiang from East China Normal University
Fig. 6. The equilibrium temperature T for the ferrite of various compositions Xc, which is calculated from DGa?c = 0.
809
for performing the BSMS investigations and Dr. X.D. Wang for the help in TEM experiments. Thanks also to Prof. Q.P. Meng for helpful discussion. [1] [2] [3] [4]
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