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γ orientation relationship on VC interphase precipitation in low-carbon steels
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Scripta Materialia 69 (2013) 17–20 www.elsevier.com/locate/scriptamat
Effects of a/c orientation relationship on VC interphase precipitation in low-carbon steels Y.-J. Zhang,a,⇑ G. Miyamoto,b K. Shinbob and T. Furuharab a
Department of Metallurgy, Graduate School of Engineering, Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai, Miyagi 980-8577, Japan b Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai, Miyagi 980-8577, Japan Received 4 March 2013; revised 18 March 2013; accepted 18 March 2013 Available online 22 March 2013
The effects of the ferrite/austenite orientation relationship (OR) on interphase precipitation in Fe–0.1C–1.5Mn–0.4V (mass%) alloy were investigated by using electron backscattering diffraction and three-dimensional atom probe. VC interphase precipitation in both sheet-like and random dispersions is obtained in ferrite without a near Kurdjumov–Sachs (K-S) OR with adjacent austenite into which the ferrite grows, while the number density of VC in ferrite with a K-S OR is far lower, indicating that a large deviation from the K-S OR is necessary for interphase precipitation in low-carbon steels. Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Interphase precipitation; Electron backscattering diffraction (EBSD); Orientation relationship; Atom probe tomography; Steel
When low-alloy low-carbon steels containing Nb, Ti or V are transformed from austenite (c) into ferrite (a), alloy carbides are precipitated in parallel rows as a result of periodic nucleation at the migrating a/c interface, which is called interphase precipitation [1]. Nanosized carbides formed through interphase precipitation have been recently used to strengthen low-carbon steels [2]. This approach plays an important role in weight saving in automobiles and thus in improving fuel efficiency. Interphase precipitation has been studied since the 1960s and several mechanisms have been proposed by different researchers [3,4]. Among these, the best-known one is the ledge mechanism proposed by Honeycombe [3]. According to him, the a/c interface is composed of immobile coherent terraces with mobile incoherent steps. Alloy carbides can be nucleated only on a coherent interface due to its lower mobility which provides enough time for nucleation. In the ledge mechanism, a Kurdjumov–Sachs (K-S) orientation relationship (OR) 1 0 1c ==½ 1 1 1a Þ is asbetween a and c ((1 1 1)c//(0 1 1)a, ½ sumed and the sheet planes are taken to be parallel close-packed planes, i.e. (0 1 1)a//(1 1 1)c. Unfortunately, this model cannot be used to explain the curved sheets which are also frequently observed. Therefore, another
⇑ Corresponding
author. [email protected]
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+81
0222152049; e-mail:
mechanism, the quasi-ledge model, was proposed by Ricks and Howell [4]. Based on this model, there is no need for a and c to have any rational OR for interphase precipitation to occur. The model is characterized by a bulging process, i.e. an originally mobile interface is pinned by alloy carbides formed on the interface and thus immobilized. Then the interface between carbides are bulged and pinned again by nucleation of new carbide particles, which leads to the formation of another sheet. This process is repeated and finally interphase precipitation in curved sheets occurs. The sheet planes can thus be parallel to any crystal planes, unlike the configuration found in the conventional ledge mechanism. However, if the volume fraction of carbide is low enough, bulging and pinning will take place irregularly and randomly dispersed carbides of lower number density will then be formed, which is the so-called bowing mechanism [3]. Recently, Okamoto et al. [5] and Yen et al. [6] studied interphase precipitation in Nb-bearing and Ti–Mo-bearing low-carbon steels, respectively. According to their experimental results, the sheet plane of alloy carbides deviates from (0 1 1)a, which cannot be explained by the ledge model proposed by Honeycombe [3]. As for a/c OR, Howell et al. [7] applied selected-area diffraction (SAD) analysis on retained c and proeutectoid a in transmission electron microscopy (TEM) to
1359-6462/$ - see front matter Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.scriptamat.2013.03.020
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Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
relate a/c crystallography to Cr23C6 precipitation and concluded that interphase precipitation only occurs on the K-S interface. In contrast, Law et al. [8] found that VC interphase precipitation takes place irrespective of a/c OR based on SAD analysis. However, determining orientation based on SAD analysis is not so accurate, especially for a small amount of retained c. Moreover, due to the small region of TEM observation, it is difficult to decide the growth direction of proeutectoid a. Miyamoto et al. [9] have recently developed a method to calculate c orientation from the orientation of martensite measured by electron backscattering diffraction (EBSD) analysis. This method can be applied to steels containing little retained c. They investigated the effects of a/c OR on VC interphase precipitation in a mediumcarbon steel and found that even in a without a K-S OR with adjacent c into which it is growing, sheet-like interphase precipitation occurs [10]. However, the characteristics of the a/c interface where carbides are nucleated have not been clarified in low-carbon steels until now. Therefore, the present study aims to clarify the effects of a/c OR on VC interphase precipitation in low-carbon steels. In order to characterize interphase precipitated VC, three-dimensional atom probe (3DAP) is used, which is known to be a powerful technique for quantitative analysis of nanosized precipitates [11,12] owing to its unique features including three-dimensional atom-byatom positional analysis and elemental identification based on mass-to-charge ratio. A V-added low-carbon steel, i.e. Fe–0.1C–1.5Mn– 0.4V (mass%), was used in the present study. The as-received alloy was homogenized at 1423 K for 345.6 ks and then machined into pieces (10 5 3 mm3). After austenitization at 1473 K for 0.6 ks, the specimens were cooled in a salt bath at a rate faster than 100 K s1 and isothermally transformed at 923 K for 60 s, followed by water quenching. Afterwards, the microstructure in the specimen etched by 3% nital was observed by optical microscopy (OM). EBSD measurements were carried out on specimens electropolished by a solution composed of 6% perchloric acid in ethanol. The dispersions of the nanosized VC precipitates in the same region as EBSD were characterized by applying a focused ion beam (FIB) microsampling method for the preparation of 3DAP specimens. The 3DAP analyses were carried out on a LEAP-4000 HR at 50 K with 20% pulse fraction and 200 kHz pulse rate. Figure 1a shows the optical microstructure of the specimen isothermally transformed at 923 K for 60 s, with prior c grain boundaries (PAGBs) indicated by dashed lines while martensite (M(c)) corresponds to untransformed c. It is found that allotriomorphic a (AF) is formed along PAGBs accompanied by a small amount of Widmanstatten a (WF). The a orientation map of the same region measured by EBSD is shown in Figure 1b. Based on the a orientation map of EBSD analysis, c orientation was calculated from martensite orientations according to the previous study [13]. Deviation angles of a/c OR from the exact K-S OR were then calculated. The circles and triangles indicate K-S and non-K-S a/c OR, which is defined by whether or not the deviation angle is less than 5 degrees in the pres-
Figure 1. (a) OM image and (b) a orientation map of the specimen partially transformed at 923 K for 60 s. AF, allotriomorphic a; WF, Widmansttaten a; M(c), martensite; PAGB, prior c grain boundary; circles and triangles indicate that K-S and non-K-S ORs are held between a and c, respectively.
ent study. According to Figure 1b, WF tends to grow into the upper c grain with which it holds the K-S OR. In contrast, most of AFs grow into the non-K-S side, although there still exist some AFs growing into the K-S side. To eliminate the influence of variant selection of a along certain PAGBs, 84 a grains at 18 PAGBs between different c grains were analyzed. The schematic of the a/ c OR analysis is shown in Figure 2a. 4h1 and 4h2 are
Figure 2. (a) Schematic for 4h1 and 4h2 for grain boundary a; (b) distributions of 4h1 and 4h2. 4h1 and 4h2: deviation angles from the exact K-S OR between a and adjacent c on each side. (4h2 > 4h1).
Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
the deviation angles from the exact K-S OR between a and c on each side, where 4h1 is smaller than 4h2. The distribution of deviation angles is summarized in Figure 2b. It is found that among all the a grains analyzed, 62% hold the K-S OR with at least one of the adjacent c grains, in which 13% hold the K-S OR with adjacent c grains on both sides. Meanwhile, 50–60% of the a/c interface can migrate during a transformation irrespective of the a/c OR. Figure 3 shows the three-dimensional V atom maps superimposed with 2 at.% V isoconcentration surfaces reconstructed from 3DAP measurements. Since carbon is also enriched in the V-rich regions, the clusters in these figures are considered to be VC precipitates. Figure 3a and c correspond to WF and AF indicated by the asterisk and diamond in Figure 1b, whose deviation angles from the exact K-S OR (4h) are 0.8° and 16.7°, respectively. Figure 3b and d show AFs taken from another field of view. It is found that almost no VC precipitate can be observed in WF, which is migrating into the K-S side. In contrast, in AF growing into the non-K-S side, the number density of VC precipitates is far higher. It is well known that interphase precipitation is generally formed on parallel sheets [3]. That is to say, such a sheetlike structure can be only observed along the direction parallel to the sheet planes. However, no matter how the reconstruction in Figure 3c is rotated, VC precipitates appear to be randomly dispersed. This type of random dispersion is also observed in other AF grains. Meanwhile, VC precipitates aligned in parallel rows are observed in some AF grains growing into the nonK-S side as shown in Figure 3d. There also exist some AF grains where both sheet-like and randomly dispersed VC are formed.
Figure 3. Three-dimensional V atom maps superimposed with 2 at.% V isoconcentration surfaces in different a grains: (a) WF (4h = 0.8°), (b) AF (4h = 3.3°), (c) AF (4h = 16.9°) and (d) AF (4h = 19.2°). (a and c) correspond to WF and AF indicated by the asterisk and diamond in Figure 1b, respectively.
19
Through cluster analysis, VC precipitation is quantified by using maximum separation method [14]. Any two solute atoms closer than a predetermined distance, dmax, are taken to be a part of the same cluster. A single cluster is then defined as consisting of all the solute atoms connected in this way. The clusters containing atoms less than another predetermined value, Nmin, are taken to be the result of random fluctuations and ignored. In the present study, only V atoms are taken into account. After trial and error, 1.0 nm and 10 are selected to be the values for dmax and Nmin, respectively. The number density of VC is obtained by simply dividing the number of clusters by the volume analyzed, while the mean cluster volume is calculated from mean number of V atoms in each V cluster, corrected by taking the detection efficiency of 3DAP (36%) and the lattice constant of VC (0.415 nm) into account. Figure 4a and b show the variations in number density and size of VC precipitates with the deviation angles from the exact K-S OR (4h) for 63 measurements, where circles and diamonds indicate sheet-like and random dispersions. Triangles indicate undefined dispersion due to low number density. According to these figures, as the a/c OR deviates from the exact K-S OR, the number density of VC increases significantly at first and remains almost constant at 4h larger than 5°. In contrast, the effect of 4h on the size of VC appears to be comparatively trivial. It should be noted that AF migrating into the K-S side also shows a relatively lower number density of VC, which means that interphase precipitation behavior occurs irrespective of a morphology. Although there exist three possible precipitation modes of VC, i.e. precipitation in c (before transformation) and a (after transformation) and interphase precipitation (during transformation), the fact that almost no precipitation of VC in a with 4h less than 5° indicates that the former
Figure 4. Variations in (a) number density and (b) size of VC precipitates with 4h.
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Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
two modes can be excluded. Therefore, it is supposed that the randomly dispersed VC precipitates in Figure 3c are formed through interphase precipitation as well. In addition, the similar number density and size of the sheet-like and random dispersions of VC precipitates also supports this suggestion. In agreement with the previous study in a mediumcarbon steel [10], VC interphase precipitation is observed on the non-K-S side in the present study. Therefore, it is obvious that the conventional ledge model proposed by Honeycombe [3] is to some extent deficient. Consequently, it is necessary for the ledge model to be extended to the non-K-S interface as suggested by Miyamoto et al. [10]. According to them, even for a a/ c interface that does not have the K-S OR, there still exist some partially coherent planes with low mobility where sheets of interphase precipitation take place. Meanwhile, they also suggest that the sheet spacing of interphase precipitation is determined by the dynamic segregation of V atoms on a migrating a/c interface. If bulging of such a non-K-S interface and subsequent nucleation of carbides occur irregularly, the dispersion of VC precipitates will be changed into a random one. The promotion of VC nucleation at a non-K-S a/c interface rather than a K-S interface might be caused by the following factors. (1) A a/c interface with a non-K-S OR generally has a higher interfacial energy than that with a K-S OR, leading to reduction in energy for the formation of critical nucleus. (2) It is well known that impurity atoms are more easily segregated at an incoherent interface than at a coherent one. Therefore, greater segregation of V may occur at a non-K-S interface, which promotes the formation of VC interphase precipitation. (3) The higher diffusivity of V at a non-K-S interface with low degree of coherency also cannot be ignored. Howell et al. found that in Fe–Cr–C alloy, the M23C6-type interphase precipitated carbide holds a rational OR with both a and c [7]. In contrast, according to previous research by Davenport et al. [15] and Khalid et al. [16], VC precipitates tend to hold the Backer– Nutting OR with a. Therefore, irrational the a/c OR observed in the present study indicates an irrational OR between c and VC as well.
The present study has investigated the effects of the a/ c OR on VC interphase precipitation in a low-carbon steel. It was found that both sheet-like and randomly dispersed interphase precipitation occurs at the non-KS interface, while the number density of VC nucleated at the K-S interface is far lower, and their dispersion is hard to identify. Therefore, it is concluded that a large deviation from the exact K-S OR is necessary for interphase precipitation to occur in low-carbon steels. The work was carried out as part of “Creation of Innovative Functions of Intelligent Materials on the Basis of Element Strategy” by CREST Basic Research Program. [1] A.T. Davenport, F.G. Berry, R.W.K. Honeycombe, Met. Sci. J. (1968) 104. [2] Y. Funakawa, T. Shiozaki, K. Tomita, T. Yamamoto, E. Maeda, ISIJ Int. 44 (2004) 1945. [3] R.W.K. Honeycombe, Metall. Mater. Trans. A 7 (1976) 915. [4] R.A. Ricks, P.R. Howell, Acta Metall. 31 (1983) 853. [5] R. Okamoto, A. Borgenstam, J. Agren, Acta Mater. 58 (2010) 4783. [6] H.W. Yen, P.Y. Chen, C.Y. Huang, J.R. Yang, Acta Mater. 59 (2011) 6264. [7] P.R. Howell, J.V. Bee, R.W.K. Honeycombe, Metall. Trans. A 10 (1979) 1213. [8] N.C. Law, S.A. Parsons, P.R. Howell, D.V. Edmonds, Mater. Sci. Technol. 3 (1987) 642. [9] G. Miyamoto, N. Takayama, T. Furuhara, Scripta Mater. 60 (2009) 1113. [10] G. Miyamoto, R. Hori, B. Poorganji, T. Furuhara, Metall. Mater. Trans. A, in press. [11] E.V. Pereloma, I.B. Timokhina, K.F. Russell, M.K. Miller, Scripta Mater. 54 (2006) 471. [12] I.B. Timokhina, P.D. Hodgson, S.P. Ringer, R.K. Zheng, E.V. Pereloma, Scripta Mater. 56 (2007) 601. [13] G. Miyamoto, N. Iwata, N. Takayama, T. Furuhara, Acta Mater. 58 (2010) 6393. [14] D. Vaumousse, A. Cerezo, P.J. Warren, Ultramicroscopy 95 (2003) 215. [15] A.T. Davenport, R.W.K. Honeycombe, Proc. R. Soc. Lond. A 322 (1971) 191. [16] F.A. Khalid, D.V. Edmonds, Mater. Sci. Technol. 9 (1993) 384.
Scripta Materialia 69 (2013) 17–20 www.elsevier.com/locate/scriptamat
Effects of a/c orientation relationship on VC interphase precipitation in low-carbon steels Y.-J. Zhang,a,⇑ G. Miyamoto,b K. Shinbob and T. Furuharab a
Department of Metallurgy, Graduate School of Engineering, Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai, Miyagi 980-8577, Japan b Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai, Miyagi 980-8577, Japan Received 4 March 2013; revised 18 March 2013; accepted 18 March 2013 Available online 22 March 2013
The effects of the ferrite/austenite orientation relationship (OR) on interphase precipitation in Fe–0.1C–1.5Mn–0.4V (mass%) alloy were investigated by using electron backscattering diffraction and three-dimensional atom probe. VC interphase precipitation in both sheet-like and random dispersions is obtained in ferrite without a near Kurdjumov–Sachs (K-S) OR with adjacent austenite into which the ferrite grows, while the number density of VC in ferrite with a K-S OR is far lower, indicating that a large deviation from the K-S OR is necessary for interphase precipitation in low-carbon steels. Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Interphase precipitation; Electron backscattering diffraction (EBSD); Orientation relationship; Atom probe tomography; Steel
When low-alloy low-carbon steels containing Nb, Ti or V are transformed from austenite (c) into ferrite (a), alloy carbides are precipitated in parallel rows as a result of periodic nucleation at the migrating a/c interface, which is called interphase precipitation [1]. Nanosized carbides formed through interphase precipitation have been recently used to strengthen low-carbon steels [2]. This approach plays an important role in weight saving in automobiles and thus in improving fuel efficiency. Interphase precipitation has been studied since the 1960s and several mechanisms have been proposed by different researchers [3,4]. Among these, the best-known one is the ledge mechanism proposed by Honeycombe [3]. According to him, the a/c interface is composed of immobile coherent terraces with mobile incoherent steps. Alloy carbides can be nucleated only on a coherent interface due to its lower mobility which provides enough time for nucleation. In the ledge mechanism, a Kurdjumov–Sachs (K-S) orientation relationship (OR) 1 0 1c ==½ 1 1 1a Þ is asbetween a and c ((1 1 1)c//(0 1 1)a, ½ sumed and the sheet planes are taken to be parallel close-packed planes, i.e. (0 1 1)a//(1 1 1)c. Unfortunately, this model cannot be used to explain the curved sheets which are also frequently observed. Therefore, another
⇑ Corresponding
author. [email protected]
Tel.:
+81
0222152049; e-mail:
mechanism, the quasi-ledge model, was proposed by Ricks and Howell [4]. Based on this model, there is no need for a and c to have any rational OR for interphase precipitation to occur. The model is characterized by a bulging process, i.e. an originally mobile interface is pinned by alloy carbides formed on the interface and thus immobilized. Then the interface between carbides are bulged and pinned again by nucleation of new carbide particles, which leads to the formation of another sheet. This process is repeated and finally interphase precipitation in curved sheets occurs. The sheet planes can thus be parallel to any crystal planes, unlike the configuration found in the conventional ledge mechanism. However, if the volume fraction of carbide is low enough, bulging and pinning will take place irregularly and randomly dispersed carbides of lower number density will then be formed, which is the so-called bowing mechanism [3]. Recently, Okamoto et al. [5] and Yen et al. [6] studied interphase precipitation in Nb-bearing and Ti–Mo-bearing low-carbon steels, respectively. According to their experimental results, the sheet plane of alloy carbides deviates from (0 1 1)a, which cannot be explained by the ledge model proposed by Honeycombe [3]. As for a/c OR, Howell et al. [7] applied selected-area diffraction (SAD) analysis on retained c and proeutectoid a in transmission electron microscopy (TEM) to
1359-6462/$ - see front matter Ó 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.scriptamat.2013.03.020
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Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
relate a/c crystallography to Cr23C6 precipitation and concluded that interphase precipitation only occurs on the K-S interface. In contrast, Law et al. [8] found that VC interphase precipitation takes place irrespective of a/c OR based on SAD analysis. However, determining orientation based on SAD analysis is not so accurate, especially for a small amount of retained c. Moreover, due to the small region of TEM observation, it is difficult to decide the growth direction of proeutectoid a. Miyamoto et al. [9] have recently developed a method to calculate c orientation from the orientation of martensite measured by electron backscattering diffraction (EBSD) analysis. This method can be applied to steels containing little retained c. They investigated the effects of a/c OR on VC interphase precipitation in a mediumcarbon steel and found that even in a without a K-S OR with adjacent c into which it is growing, sheet-like interphase precipitation occurs [10]. However, the characteristics of the a/c interface where carbides are nucleated have not been clarified in low-carbon steels until now. Therefore, the present study aims to clarify the effects of a/c OR on VC interphase precipitation in low-carbon steels. In order to characterize interphase precipitated VC, three-dimensional atom probe (3DAP) is used, which is known to be a powerful technique for quantitative analysis of nanosized precipitates [11,12] owing to its unique features including three-dimensional atom-byatom positional analysis and elemental identification based on mass-to-charge ratio. A V-added low-carbon steel, i.e. Fe–0.1C–1.5Mn– 0.4V (mass%), was used in the present study. The as-received alloy was homogenized at 1423 K for 345.6 ks and then machined into pieces (10 5 3 mm3). After austenitization at 1473 K for 0.6 ks, the specimens were cooled in a salt bath at a rate faster than 100 K s1 and isothermally transformed at 923 K for 60 s, followed by water quenching. Afterwards, the microstructure in the specimen etched by 3% nital was observed by optical microscopy (OM). EBSD measurements were carried out on specimens electropolished by a solution composed of 6% perchloric acid in ethanol. The dispersions of the nanosized VC precipitates in the same region as EBSD were characterized by applying a focused ion beam (FIB) microsampling method for the preparation of 3DAP specimens. The 3DAP analyses were carried out on a LEAP-4000 HR at 50 K with 20% pulse fraction and 200 kHz pulse rate. Figure 1a shows the optical microstructure of the specimen isothermally transformed at 923 K for 60 s, with prior c grain boundaries (PAGBs) indicated by dashed lines while martensite (M(c)) corresponds to untransformed c. It is found that allotriomorphic a (AF) is formed along PAGBs accompanied by a small amount of Widmanstatten a (WF). The a orientation map of the same region measured by EBSD is shown in Figure 1b. Based on the a orientation map of EBSD analysis, c orientation was calculated from martensite orientations according to the previous study [13]. Deviation angles of a/c OR from the exact K-S OR were then calculated. The circles and triangles indicate K-S and non-K-S a/c OR, which is defined by whether or not the deviation angle is less than 5 degrees in the pres-
Figure 1. (a) OM image and (b) a orientation map of the specimen partially transformed at 923 K for 60 s. AF, allotriomorphic a; WF, Widmansttaten a; M(c), martensite; PAGB, prior c grain boundary; circles and triangles indicate that K-S and non-K-S ORs are held between a and c, respectively.
ent study. According to Figure 1b, WF tends to grow into the upper c grain with which it holds the K-S OR. In contrast, most of AFs grow into the non-K-S side, although there still exist some AFs growing into the K-S side. To eliminate the influence of variant selection of a along certain PAGBs, 84 a grains at 18 PAGBs between different c grains were analyzed. The schematic of the a/ c OR analysis is shown in Figure 2a. 4h1 and 4h2 are
Figure 2. (a) Schematic for 4h1 and 4h2 for grain boundary a; (b) distributions of 4h1 and 4h2. 4h1 and 4h2: deviation angles from the exact K-S OR between a and adjacent c on each side. (4h2 > 4h1).
Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
the deviation angles from the exact K-S OR between a and c on each side, where 4h1 is smaller than 4h2. The distribution of deviation angles is summarized in Figure 2b. It is found that among all the a grains analyzed, 62% hold the K-S OR with at least one of the adjacent c grains, in which 13% hold the K-S OR with adjacent c grains on both sides. Meanwhile, 50–60% of the a/c interface can migrate during a transformation irrespective of the a/c OR. Figure 3 shows the three-dimensional V atom maps superimposed with 2 at.% V isoconcentration surfaces reconstructed from 3DAP measurements. Since carbon is also enriched in the V-rich regions, the clusters in these figures are considered to be VC precipitates. Figure 3a and c correspond to WF and AF indicated by the asterisk and diamond in Figure 1b, whose deviation angles from the exact K-S OR (4h) are 0.8° and 16.7°, respectively. Figure 3b and d show AFs taken from another field of view. It is found that almost no VC precipitate can be observed in WF, which is migrating into the K-S side. In contrast, in AF growing into the non-K-S side, the number density of VC precipitates is far higher. It is well known that interphase precipitation is generally formed on parallel sheets [3]. That is to say, such a sheetlike structure can be only observed along the direction parallel to the sheet planes. However, no matter how the reconstruction in Figure 3c is rotated, VC precipitates appear to be randomly dispersed. This type of random dispersion is also observed in other AF grains. Meanwhile, VC precipitates aligned in parallel rows are observed in some AF grains growing into the nonK-S side as shown in Figure 3d. There also exist some AF grains where both sheet-like and randomly dispersed VC are formed.
Figure 3. Three-dimensional V atom maps superimposed with 2 at.% V isoconcentration surfaces in different a grains: (a) WF (4h = 0.8°), (b) AF (4h = 3.3°), (c) AF (4h = 16.9°) and (d) AF (4h = 19.2°). (a and c) correspond to WF and AF indicated by the asterisk and diamond in Figure 1b, respectively.
19
Through cluster analysis, VC precipitation is quantified by using maximum separation method [14]. Any two solute atoms closer than a predetermined distance, dmax, are taken to be a part of the same cluster. A single cluster is then defined as consisting of all the solute atoms connected in this way. The clusters containing atoms less than another predetermined value, Nmin, are taken to be the result of random fluctuations and ignored. In the present study, only V atoms are taken into account. After trial and error, 1.0 nm and 10 are selected to be the values for dmax and Nmin, respectively. The number density of VC is obtained by simply dividing the number of clusters by the volume analyzed, while the mean cluster volume is calculated from mean number of V atoms in each V cluster, corrected by taking the detection efficiency of 3DAP (36%) and the lattice constant of VC (0.415 nm) into account. Figure 4a and b show the variations in number density and size of VC precipitates with the deviation angles from the exact K-S OR (4h) for 63 measurements, where circles and diamonds indicate sheet-like and random dispersions. Triangles indicate undefined dispersion due to low number density. According to these figures, as the a/c OR deviates from the exact K-S OR, the number density of VC increases significantly at first and remains almost constant at 4h larger than 5°. In contrast, the effect of 4h on the size of VC appears to be comparatively trivial. It should be noted that AF migrating into the K-S side also shows a relatively lower number density of VC, which means that interphase precipitation behavior occurs irrespective of a morphology. Although there exist three possible precipitation modes of VC, i.e. precipitation in c (before transformation) and a (after transformation) and interphase precipitation (during transformation), the fact that almost no precipitation of VC in a with 4h less than 5° indicates that the former
Figure 4. Variations in (a) number density and (b) size of VC precipitates with 4h.
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Y.-J. Zhang et al. / Scripta Materialia 69 (2013) 17–20
two modes can be excluded. Therefore, it is supposed that the randomly dispersed VC precipitates in Figure 3c are formed through interphase precipitation as well. In addition, the similar number density and size of the sheet-like and random dispersions of VC precipitates also supports this suggestion. In agreement with the previous study in a mediumcarbon steel [10], VC interphase precipitation is observed on the non-K-S side in the present study. Therefore, it is obvious that the conventional ledge model proposed by Honeycombe [3] is to some extent deficient. Consequently, it is necessary for the ledge model to be extended to the non-K-S interface as suggested by Miyamoto et al. [10]. According to them, even for a a/ c interface that does not have the K-S OR, there still exist some partially coherent planes with low mobility where sheets of interphase precipitation take place. Meanwhile, they also suggest that the sheet spacing of interphase precipitation is determined by the dynamic segregation of V atoms on a migrating a/c interface. If bulging of such a non-K-S interface and subsequent nucleation of carbides occur irregularly, the dispersion of VC precipitates will be changed into a random one. The promotion of VC nucleation at a non-K-S a/c interface rather than a K-S interface might be caused by the following factors. (1) A a/c interface with a non-K-S OR generally has a higher interfacial energy than that with a K-S OR, leading to reduction in energy for the formation of critical nucleus. (2) It is well known that impurity atoms are more easily segregated at an incoherent interface than at a coherent one. Therefore, greater segregation of V may occur at a non-K-S interface, which promotes the formation of VC interphase precipitation. (3) The higher diffusivity of V at a non-K-S interface with low degree of coherency also cannot be ignored. Howell et al. found that in Fe–Cr–C alloy, the M23C6-type interphase precipitated carbide holds a rational OR with both a and c [7]. In contrast, according to previous research by Davenport et al. [15] and Khalid et al. [16], VC precipitates tend to hold the Backer– Nutting OR with a. Therefore, irrational the a/c OR observed in the present study indicates an irrational OR between c and VC as well.
The present study has investigated the effects of the a/ c OR on VC interphase precipitation in a low-carbon steel. It was found that both sheet-like and randomly dispersed interphase precipitation occurs at the non-KS interface, while the number density of VC nucleated at the K-S interface is far lower, and their dispersion is hard to identify. Therefore, it is concluded that a large deviation from the exact K-S OR is necessary for interphase precipitation to occur in low-carbon steels. The work was carried out as part of “Creation of Innovative Functions of Intelligent Materials on the Basis of Element Strategy” by CREST Basic Research Program. [1] A.T. Davenport, F.G. Berry, R.W.K. Honeycombe, Met. Sci. J. (1968) 104. [2] Y. Funakawa, T. Shiozaki, K. Tomita, T. Yamamoto, E. Maeda, ISIJ Int. 44 (2004) 1945. [3] R.W.K. Honeycombe, Metall. Mater. Trans. A 7 (1976) 915. [4] R.A. Ricks, P.R. Howell, Acta Metall. 31 (1983) 853. [5] R. Okamoto, A. Borgenstam, J. Agren, Acta Mater. 58 (2010) 4783. [6] H.W. Yen, P.Y. Chen, C.Y. Huang, J.R. Yang, Acta Mater. 59 (2011) 6264. [7] P.R. Howell, J.V. Bee, R.W.K. Honeycombe, Metall. Trans. A 10 (1979) 1213. [8] N.C. Law, S.A. Parsons, P.R. Howell, D.V. Edmonds, Mater. Sci. Technol. 3 (1987) 642. [9] G. Miyamoto, N. Takayama, T. Furuhara, Scripta Mater. 60 (2009) 1113. [10] G. Miyamoto, R. Hori, B. Poorganji, T. Furuhara, Metall. Mater. Trans. A, in press. [11] E.V. Pereloma, I.B. Timokhina, K.F. Russell, M.K. Miller, Scripta Mater. 54 (2006) 471. [12] I.B. Timokhina, P.D. Hodgson, S.P. Ringer, R.K. Zheng, E.V. Pereloma, Scripta Mater. 56 (2007) 601. [13] G. Miyamoto, N. Iwata, N. Takayama, T. Furuhara, Acta Mater. 58 (2010) 6393. [14] D. Vaumousse, A. Cerezo, P.J. Warren, Ultramicroscopy 95 (2003) 215. [15] A.T. Davenport, R.W.K. Honeycombe, Proc. R. Soc. Lond. A 322 (1971) 191. [16] F.A. Khalid, D.V. Edmonds, Mater. Sci. Technol. 9 (1993) 384.