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Study on the crystallographic orientation relationship and formation mechanism of reversed austenite in economical Cr12 super martensitic stainless steel
Materials Characterization 109 (2015) 100–106
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Study on the crystallographic orientation relationship and formation mechanism of reversed austenite in economical Cr12 super martensitic stainless steel Dong Ye a, Shaohong Li a, Jun Li a, Wen Jiang a, Jie Su b, Kunyu Zhao a,⁎ a b
Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China Institute for Structural Materials, Central Iron and Steel Research Institute, Beijing 100081, China
a r t i c l e
i n f o
Article history: Received 25 March 2015 Received in revised form 16 September 2015 Accepted 20 September 2015 Available online 26 September 2015 Keywords: Super martensitic stainless steel Reversed austenite Crystallographic orientation Electron backscatter diffraction
a b s t r a c t Effect of carbides and crystallographic orientation relationship on the formation mechanism of reversed austenite of economical Cr12 super martensitic stainless steel (SMSS) has been investigated mainly by transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The results indicate that the M23C6 precipitation and the formation of the reversed austenite have the interaction effect during tempering process in SMSS. The reversed austenite forms intensively at the sub-block boundary and the lath boundary within a misorientation range of 0–60°. M23C6 has the same crystallographic orientation relationship with reversed austenite. There are two different kinds of formation modes for reversed austenite. One is a nondiffusional shear reversion; the other is a diffusion transformation. Both are strictly limited by crystallographic orientation relationship. The austenite variants are limited to two kinds within one packet and five kinds within one prior austenite grain. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Super martensitic stainless steels (SMSSs) are a series of ultra low carbon martensitic stainless steels with a chromium content in the range 13–17 wt.% which demonstrate both high mechanical strength and good corrosion resistance, and they are much cheaper than duplex stainless steels [1]. The enhanced properties are achieved by substantial lowering of the carbon content below 0.07% and raising the Ni content to the range of 3–6% [2,3]. They have found an ample range of applications in variety of fields such as petroleum and natural gas transmission pipelines, shipbuilding, offshore drilling platforms, and hydraulic turbine blade industry [4–7]. Microstructure of super martensitic stainless steel (SMSS) consists of lath martensite and reversed austenite. This structure gives them an excellent combination of good mechanical properties and high corrosion resistance [1]. It is well known that the mechanical properties of this kind of steel, particularly the yield stress, are strongly dependent on the volume fraction of reversed austenite, which is very sensitive to the heat treatment [2,5,8–11]. Tolchard et al. [1] also indicated that austenite has a higher hydrogen solubility and greater resistance to embrittlement than martensite. Kimura et al. [12] investigated the impact of retained
⁎ Corresponding author at: Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China. E-mail address: [email protected] (K. Zhao).
http://dx.doi.org/10.1016/j.matchar.2015.09.013 1044-5803/© 2015 Elsevier Inc. All rights reserved.
austenite on corrosion rate and pitting susceptibility of modified 13Cr steel. These authors found that retained austenite had a small positive effect on the corrosion resistance of the 13Cr SMSS [1,12–15]. Today, the major issues and challenges of SMSSs have to do with the mechanical and corrosion performance and their weldability, as well as the development of fast, reliable, economic welding processes that minimize or even dispense with expensive and time-consuming post-weld heat treatments [16–18]. The recent studies indicate that the amount, morphology and distribution of the reversed austenite with sub-micron/ nano size in SMSS are the main factors that influence its comprehensive properties, so it is essential to study microstructural characteristics, volume fraction and formation mechanism of the reversed austenite in SMSS. In our previous studies [19–21], a series of SMSS was researched and we concluded that the reversed austenite with low dislocation density is formed at high temperature tempering processing in SMSSs. The transformation of the martensite to reversed austenite is a diffusion transformation, and the growth of the reversed austenite is closely related to the diffusion process of Ni. The volume fraction of reversed austenite has great influence on the mechanical properties, especially ductility for the steel. In the present work, the further formation mechanism of the reversed austenite of SMSS after high temperature tempering was investigated by using transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The research was focused on the crystallographic orientation of reversed austenite in SMSS. In addition, the interplay of M23C6 on the formation of reversed austenite was analyzed in details.
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2. Experimental procedure The experimental steel was melted with extra low impurity contents in a vacuum induction melting furnace (rated capacity: 25 kg, rated power: 100 kw, limited vacuum degree: 6 × 10− 3Pa, boosting rate: 0.05 Pa/min, maximum temperature: 1700°°C, cooling water pressure: 0.35 MPa). The ingots were cast and hot forged into round rods with diameter of 15 mm. The chemical composition of the experimental steel is shown in Table 1. The heat treatment processes including quenching and high temperature tempering were determined by both equilibrium phase diagram calculated by thermo-calc software and previous research data. Solution annealing was carried out at 1050°°C for half an hour using a vertical type furnace, and the samples were then quenched in oil. Tempering was carried out at the temperature range from 400°°C to 750°°C for 2 h. Thermal expansion experiments of the experimental steel were carried out using Formastor-F II full-automatic transformation measuring apparatus. The samples were cylindrical, with height 10 mm and diameter 3 mm, and there is a hole with diameter 2 mm on the end. First, the samples were heated at a rate of 0.05°°C/s. Then, they were held at 1100°°C for 10 min to homogenize their composition and temperature. After that, the samples were cooled to room temperature at a rate of 100°°C/s. Phase transition points were measured, including austenite reversed transformation start point (As), austenite reversed transformation finish point (Af), martensite transformation start point (Ms) and martensite transformation finish point (Mf). The volume fraction of austenite in different heat treatment was determined by PHILIPS APD-10 full-automatic X-ray diffraction instrument using Co radiation from 40 to 120° at a step interval of 0.02°, and the volume fraction was calculated by the following formula, I F110 ϕA ¼ 1= 1 þ K ; K ¼ I0A111 =I0 F110 IA111
ð1Þ
Where I0A111 and I0F110 are pure austenite and pure ferrite integrated intensity, respectively. The morphology and distribution of reversed austenite and the precipitation phase of the samples were investigated by Philip EM420 transmission electron microscope (TEM) and JES 2100 high resolution transmission electron microscope (HRTEM) after jet polishing of a thin foil in a solution of 6% perchloric acid and 94% anhydrous ethanol. The crystallographic orientation relationship between reversed austenite and martensite matrix was determined by electron diffraction patterns. Crystallographic characterization was also carried out by means of the EBSD method using FEI Quanta 650 FEG thermal field emission scanning electron microscope with HKL Channel 5. The scanning step is 0.2 μm. The results obtained by EBSD method were analyzed by means of the Orientation Imaging Microscopy (OIM™) system.
Fig. 1. X-ray diffraction spectrum of the experimental steel after quenched and tempered.
experiments, and the results show that As = 585°°C; Af = 730°°C; Ms. = 310°°C; Mf = 170°°C. Based on calculation of Fig. 1, the variation curve in the amount of austenite with tempering temperature for 2 h after quenched at 1050°°C is shown in Fig. 2. It should be noted that no austenite is detected when the tempering temperature is lower than the As temperature (585°°C), and there is a steep increase in volume fraction of austenite above 600°°C. It is believed that this part of austenite formed after tempering is reversed austenite. In general, above 550°°C, the volume fraction of austenite increases with raised temperature and it reaches a maximum value of 6.28% at 650°°C. Subsequently, the volume fraction of austenite starts to decrease. This calculated result is in accordance with the results attained by the previous studies on 15Cr SMSS [19]. According to the previous studies, it is known that during tempering, the volume fraction of the reversed austenite is controlled by two factors: the transformation volume of austenite under high temperature and the stability of austenite in tempering cooling process [19,22,23]. So there exists an optimal temperature under which the volume fraction of reversed austenite is maximum in the temperature range from As to Af.
3. Results and discussion 3.1. The volume fraction of austenite tested by XRD Fig. 1 shows the result of XRD of the steel subjected to quenching and tempering. According to the XRD pattern of quenching (black line in Fig. 1), it reveals that the content of the retained austenite is close to zero by considering the equipment and calculation errors. It means that there is no retained austenite after quenching at 1050°°C. The microstructure is uniphase martensite after quenching. Phase transition points of the experimental steel are measured by thermal expansion Table 1 Chemical compositions of the experimental steel, (wt.%). C
Mn
Si
Cr
Ni
Mo
0.019
0.40
0.16
11.69
4.94
2.04
Fig. 2. Volume fractions of austenite in the experimental steel with different tempering temperatures.
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3.2. Microstructure analysis Fig. 3(a) shows a TEM micrograph of a specimen subjected to temping at 550°°C for 2 h after quenching at 1050°°C. As shown, the specimen has a typical lath martensitic structure with darker strips. It can be observed that there is a tiny reversed austenite formed at the lath martensite grain boundary. According to the XRD results and the As temperature, we know that the martensitic reversed transformation has occurred at tempering temperature in the range from 550°°C to 600°°C, so the reversed austenite start to form at 550°°C. It is very difficult to observe such small volume of reversed austenite at 550°°C. Fig. 3(b) shows selected area electronic diffraction pattern (SADP) for Fig. 3(a). The orientation relationship between reversed austenite and lath martensite is(111)A//(011)M and ½110A ==½111M , which meets the Kurdjumov–Sachs (K–S) relationship. This orientation relationship is formed to reduce the interfacial energy of reversed austenite nucleation [23]. Previous studies [8] have shown that the special grain orientation relationship, besides chemical compositions, plays important roles on the mechanical stability of the reversed austenite. The micrograph observed by TEM of the tested steel tempered at 650°°C for 2 h after quenching at 1050°°C is show in Fig. 4(a). The morphologies of martensite matrix, the reversed austenite and fine carbides are displayed. The reversed austenite has thin sheets and platelets with ~100 nm in the martensite matrix. According to the electronic diffraction pattern for selected area shown in Fig. 4(b), the carbides can be determined as M23C6 by calibration. This carbide has a FCC lattice with space group Fm3m, which has 92 metal atoms and 24 carbon atoms in each unit cell. Fig. 4(a) displays that M23C6 carbides are formed on the M/γ interfaces. The nucleation energy of M23C6 can be calculated using the following equation [24], ΔG ¼ α1
σ3 ðΔGV þ εÞ2
:
ð2Þ
Where α1 is the geometry coefficient, σ is specific interfacial energy, ΔGV and ε are the free enthalpy variation and the elastic strain energy per unit volume of parent phase, respectively. In our case, theoretically, there are three main types of interface: the martensite–martensite interface (α/α), the new reversed austenite– martensite interface (α/γ) and the internal austenite–austenite interface (γ/γ). However, the austenite–austenite interface (γ/γ) can be neglected here, since it does not appear as reversed austenite distributes and disperses in the martensite matrix. It is known that martensite grain boundaries differ in terms of their misorientation and energy [25,26]. Most importantly, we showed the reversed austenite and the martensite meet the K-S orientation relationship. According the study by D. Raabe et al. [25], the interfacial energies of the martensite-austenite grain boundaries are generally assumed to be equal to or lower than
high angle martensite–martensite grain boundaries. And their research results indicated that the reversed austenite–martensite interface (α/γ) with K-S orientation relationship always has the lower interface energy. Therefore, because σα/γis lower than σα/α, the nucleation energy of M23C6 at M/γ interface is lower than at α/α interface. It could be confirmed that M23C6 is easy to form at the M/γ interface. Fig. 4 displays the precipitated M23C6 carbides have a FCC crystal structure and obey the cube-on-cube orientation relationship with the reversed austenite: [111]γ//[111]M23C6 and ð220Þγ==ð220ÞM23 C6 , which means the M23C6 carbides have the same orientation with γ, and this will contribute to phase transformation by reducing strain energy. In Fig. 4(b), we can also see that there are three diffraction spots of M23C6 between two diffraction spots of γ. This is obviously because that M23C6 has the same crystal structure with γ and its crystal lattice parameter is three times of γ [27,28]. Fig. 5 shows TEM image of reversed austenite in a specimen subjected to tempering at 650°°C for 2 h after quenching at 1050°°C. As shown in Fig. 5(a), the darker reversed austenite platelet is formed at the region with high dislocation density, and it is growing along lath martensite boundary. M/γ interface gradually migrates to lath martensite, and the reversed austenite swallows the lath martensite and continues to grow. It is found that the M/γ interface is rectilinear and nearly parallel to the martensite lath. This indicates that the nucleation and growth of the reversed austenite follow the certain orientation relationship. Fig. 5 (b) shows a thin sheet reversed austenite with some stacking faults. The stacking faults could provide more nucleation sites for reversed austenite, and contribute to the revered transformation by reducing the energy barrier that need to overcome the shear deformation. Fig. 5(b) also reveals that the dislocation density in the reversed austenite is obviously lower than that in the martensite matrix, which indicates that the reversed austenite does not inherit the high dislocation density structure of prior martensite matrix, and this is the indirect evidence that the austenite is formed by diffusion [19]. Regarding the precipitation of M23C6 and the reversed austenite, various groups have studied the nucleation and growth of the abutting phases as well as the crystallographic orientations. S.H. Zhang et al. [29] detected that the nucleation sequence of M23C6 and reversed austenite is difficult to be clarified. Y.Y. Song et al. [28] indicated that M23C6 carbides precipitated at the prior martensitic lath boundaries supplied the nucleation sites for the reversed austenite formation and supplied Ni-enrichment in the adjacent region. H.U. Hong et al. [30] reported the M23C6 carbide tends to select the preferential planes which have the lowest interfacial energy. H. Li et al. [31] observed that the atoms near the interface of carbide and matrix will redistribute during the growth of M23C6, and the Cr atoms migrate into the carbide, while the Fe and Ni atoms migrate into the matrix. From the above discussion, on the one hand, the formation of reversed austenite not only provides the nucleation sites of M23C6 at
Fig. 3. TEM micrographs of reversed austenite along the lath boundary in the experimental steel tempered at 550°°C for 2 h: (a) bright field and (b) SAD pattern.
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Fig. 4. TEM images of the experimental steel tempered at 650°°C after quenching at 1050°°C: (a) bright field; (b) the corresponding SAED pattern; (c) schematic diagram of (b).
the new formed K-S austenite–martensite interface (α/γ) with low interface energy, but also promotes the high equilibrium segregation of Cr to the boundary, leading to a high chemical driving force to the nucleation of M23C6. On the other hand, the M23C6 carbides precipitated at M/γ interface can cause the decrease of interface energy and the formation of the chromium depleted region and nickel enrichment zone. Then the diffusion of Ni in the adjacent area becomes easier. Thus the M23C6 carbides precipitation provides the energy of the nucleation, and leads to the formation and growth of reversed austenite. These indicate that M23C6 precipitation and the formation of reversed austenite have the interaction effect. 3.3. EBSD analysis In order to further clarify the orientation relationship of reversed austenite and martensite, the experimental steel tempering at 650°°C for 2 h is measured by using EBSD. Fig. 6(a) shows the SEM photo of the flat surface of sample by electrolytic polishing. Fig. 6(b) displays EBSD inverse pole figure color map of the area in Fig. 6(a). In Fig. 6(b), white color represents reversed austenite, and different other colors represent different crystallographic orientation martensites. Fig. 6(c) is the superposition graph which consists of martensite quality image and inverse pole figure color map of reversed austenite in Fig. 6(a). Fig. 6(d) is the superposition graph which consists of boundary map and inverse pole figure color map of the reversed austenite in Fig. 6(a). The dispersions of reversed austenite with different colors are shown in Fig. 6(c) and 6(d). In these maps, black
solid lines follow all kinds of boundaries, including prior austenite grain boundaries, packet boundaries and lath boundaries. The bcc/fcc map is shown in the Fig. 6(e), and the red color represents reversed austenite (fcc), and the white color represents the martensite (bcc). It clearly reveals the reversed austenites distribute in the martensite matrix. As shown in Fig. 6, there are 8 prior austenite grains in the 75∗75 μm area. Each prior austenite grain contains more than one packet (the martensite lath collection with same habit plane), and each packet contains more than one martensite block (the martensite lath collection with approaching orientation), which has the similar color and orientation sub laths. Fig. 7 represents the result of interface distribution analysis in the martensite structure. According to Fig. 6 (d) and Fig. 7, the packet and block boundaries are the interface which misorientation is more than 10°, and the block boundary is mostly twin boundary with misorientation of about 60°. The misorientation of sub-block boundary is 5 ~ 10°, and the martensite misorientation maintains a statistically random distribution. These maps reveal that reversed austenite mostly forms at the interior of martensite block, not only on the sub-block boundaries, but also on the lath boundaries. Only a few reversed austenites form on prior austenite grain boundaries, packet boundaries and block boundaries. Based on the results of Fig. 7, reversed austenite is concentrated in the martensitic microstructure phase interface with the misorientation range of 0° ~ 60°. In order to improve orientation analysis, one prior austenite grain area which contains 4 martensite packets in the upper of Fig. 8(a) and
Fig. 5. TEM micrographs of reversed austenite in the experimental steel tempered at 650°°C for 2 h.
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Fig. 6. Crystallographic orientation maps of the experimental steel tempered at 650°°C for 2 h. (a) SEM micrograph of an area in the experimental steel; (b) inverse pole figure of the martensite matrix and reversed austenite of (a); (c) inverse pole figure of the reversed austenite; (d) boundary map; (e) the bcc/fcc map.
8(b) is studied. Fig. 8(c) and 8(d) are {100} pole figures showing orientations of the lath martensite and reversed austenite corresponding to the EBSD orientation map, respectively. In Fig. 8(c), it can be concluded that the martensite orientation relationship conforms to the K-S relationship variant orientation. The three open circle symbols (○) in Fig. 8(c) and 8(d) represent {100} orientation projection of the prior austenite grain. Fig. 8(d) also shows the green symbols and three open circle symbols (○) coincide with each other. This means that the “green” reversed austenite (the green symbols) has the same orientation with the prior austenite. Whereas the prior austenite presents the K-S orientation relationship with martensite matrix, the “green” reversed austenite presents the identical orientation
Fig. 7. Interface distribution in the martensite structure.
relationship with martensite matrix. And the close packed plane and direction of the two are parallel to each other, just like (011)α//(111)γ and ½111α==½101γ. It can be concluded that the nucleation and growth of the “green” reversed austenite, which is strictly limited by crystallographic orientation, is the reversible and nondiffusional shear deformation process (i.e., the shear reversion mechanism) [32]. This kind of reversed austenite transforms from bcc into fcc by two times shearing. It is found that when the reversed austenite grains are formed on the lath boundary plane, the α′/γ interfacial energy can be minimized and the activation energy of nucleation can be reduced. The close packed plane which is nearly parallel to the lath boundary is chose as the habit plane of the nucleation of the reversed austenite. That is to say, the habit plane is the lath boundary plane which is close to {011}α. {110} and {111} pole figures for all reversed austenite are shown in Fig. 8(e). Compared with Fig. 8(d), reversed austenites with other colors have essentially the same orientation with the surrounding martensite, and their distribution characteristic also meet with K-S polar diagram pattern. However, there are different close packed plane and different close packed direction between the other color reversed austenite with the “green” reversed austenite. It means that there is no habit plane among the other color reversed austenite and surrounding martensite matrix. Therefore we consider that the nucleation position and growing mechanism of the other color reversed austenite are different with the “green” reversed austenite. In our previous studies [19], we deeply discussed the growing mechanism of this kind of reversed austenite of a series of SMSSs. It is a diffused phase transformation, and the growth of these reversed austenites is closely related to the diffusion process of Ni. As determined in Fig. 8, although all reversed austenite which are inside the prior austenite have different crystallographic orientation relationships, the reversed austenite variants are limited to two kinds within one packet and five kinds within one prior austenite grain. That is, they are in agreement with “the austenite limitation” which many studies carried out [32].
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Fig. 8. Crystallographic orientation maps and pole figures of one prior austenite. (a) Inverse pole figure of the martensite matrix and reversed austenite; (b) inverse pole figure of reversed austenite; (c){100} pole figure showing orientations of the lath martensite corresponding to the EBSD orientation map; (d) {100} pole figure showing orientations of the reversed austenite; (e) {110} and {111} pole figure of reversed austenite.
4. Conclusions Experiments were carried out aiming to determine the influence of M23C6 on the formation of reversed austenite and crystallographic orientation of reversed austenite of SMSS, and the main conclusions of the work can be listed as follows: 1. M23C6 is main carbide precipitated at tempering. M23C6 has the same crystallographic orientation relationship with reversed austenite, so it provides a nucleation condition for the formation of reversed austenite. The M23C6 precipitation promotes the formation of the reversed
austenite during tempering process in SMSS. The M23C6 precipitation and the formation of the reversed austenite have the interaction effect during tempering process. 2. With a nucleation and growth mechanism, the reversed austenite forms within the prior austenite grain uniformly. Especially, it forms intensively at the sub-block boundary and the lath boundary within a misorientation range of 0–60°. Two kinds of austenite variants with different orientation relationships can be formed inside a packet. 3. There are two different kinds of formation modes for reversed austenite. One is the nondiffusional shear reversion; the other is the
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Study on the crystallographic orientation relationship and formation mechanism of reversed austenite in economical Cr12 super martensitic stainless steel Dong Ye a, Shaohong Li a, Jun Li a, Wen Jiang a, Jie Su b, Kunyu Zhao a,⁎ a b
Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China Institute for Structural Materials, Central Iron and Steel Research Institute, Beijing 100081, China
a r t i c l e
i n f o
Article history: Received 25 March 2015 Received in revised form 16 September 2015 Accepted 20 September 2015 Available online 26 September 2015 Keywords: Super martensitic stainless steel Reversed austenite Crystallographic orientation Electron backscatter diffraction
a b s t r a c t Effect of carbides and crystallographic orientation relationship on the formation mechanism of reversed austenite of economical Cr12 super martensitic stainless steel (SMSS) has been investigated mainly by transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The results indicate that the M23C6 precipitation and the formation of the reversed austenite have the interaction effect during tempering process in SMSS. The reversed austenite forms intensively at the sub-block boundary and the lath boundary within a misorientation range of 0–60°. M23C6 has the same crystallographic orientation relationship with reversed austenite. There are two different kinds of formation modes for reversed austenite. One is a nondiffusional shear reversion; the other is a diffusion transformation. Both are strictly limited by crystallographic orientation relationship. The austenite variants are limited to two kinds within one packet and five kinds within one prior austenite grain. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Super martensitic stainless steels (SMSSs) are a series of ultra low carbon martensitic stainless steels with a chromium content in the range 13–17 wt.% which demonstrate both high mechanical strength and good corrosion resistance, and they are much cheaper than duplex stainless steels [1]. The enhanced properties are achieved by substantial lowering of the carbon content below 0.07% and raising the Ni content to the range of 3–6% [2,3]. They have found an ample range of applications in variety of fields such as petroleum and natural gas transmission pipelines, shipbuilding, offshore drilling platforms, and hydraulic turbine blade industry [4–7]. Microstructure of super martensitic stainless steel (SMSS) consists of lath martensite and reversed austenite. This structure gives them an excellent combination of good mechanical properties and high corrosion resistance [1]. It is well known that the mechanical properties of this kind of steel, particularly the yield stress, are strongly dependent on the volume fraction of reversed austenite, which is very sensitive to the heat treatment [2,5,8–11]. Tolchard et al. [1] also indicated that austenite has a higher hydrogen solubility and greater resistance to embrittlement than martensite. Kimura et al. [12] investigated the impact of retained
⁎ Corresponding author at: Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China. E-mail address: [email protected] (K. Zhao).
http://dx.doi.org/10.1016/j.matchar.2015.09.013 1044-5803/© 2015 Elsevier Inc. All rights reserved.
austenite on corrosion rate and pitting susceptibility of modified 13Cr steel. These authors found that retained austenite had a small positive effect on the corrosion resistance of the 13Cr SMSS [1,12–15]. Today, the major issues and challenges of SMSSs have to do with the mechanical and corrosion performance and their weldability, as well as the development of fast, reliable, economic welding processes that minimize or even dispense with expensive and time-consuming post-weld heat treatments [16–18]. The recent studies indicate that the amount, morphology and distribution of the reversed austenite with sub-micron/ nano size in SMSS are the main factors that influence its comprehensive properties, so it is essential to study microstructural characteristics, volume fraction and formation mechanism of the reversed austenite in SMSS. In our previous studies [19–21], a series of SMSS was researched and we concluded that the reversed austenite with low dislocation density is formed at high temperature tempering processing in SMSSs. The transformation of the martensite to reversed austenite is a diffusion transformation, and the growth of the reversed austenite is closely related to the diffusion process of Ni. The volume fraction of reversed austenite has great influence on the mechanical properties, especially ductility for the steel. In the present work, the further formation mechanism of the reversed austenite of SMSS after high temperature tempering was investigated by using transmission electron microscopy (TEM) and electron backscatter diffraction (EBSD). The research was focused on the crystallographic orientation of reversed austenite in SMSS. In addition, the interplay of M23C6 on the formation of reversed austenite was analyzed in details.
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2. Experimental procedure The experimental steel was melted with extra low impurity contents in a vacuum induction melting furnace (rated capacity: 25 kg, rated power: 100 kw, limited vacuum degree: 6 × 10− 3Pa, boosting rate: 0.05 Pa/min, maximum temperature: 1700°°C, cooling water pressure: 0.35 MPa). The ingots were cast and hot forged into round rods with diameter of 15 mm. The chemical composition of the experimental steel is shown in Table 1. The heat treatment processes including quenching and high temperature tempering were determined by both equilibrium phase diagram calculated by thermo-calc software and previous research data. Solution annealing was carried out at 1050°°C for half an hour using a vertical type furnace, and the samples were then quenched in oil. Tempering was carried out at the temperature range from 400°°C to 750°°C for 2 h. Thermal expansion experiments of the experimental steel were carried out using Formastor-F II full-automatic transformation measuring apparatus. The samples were cylindrical, with height 10 mm and diameter 3 mm, and there is a hole with diameter 2 mm on the end. First, the samples were heated at a rate of 0.05°°C/s. Then, they were held at 1100°°C for 10 min to homogenize their composition and temperature. After that, the samples were cooled to room temperature at a rate of 100°°C/s. Phase transition points were measured, including austenite reversed transformation start point (As), austenite reversed transformation finish point (Af), martensite transformation start point (Ms) and martensite transformation finish point (Mf). The volume fraction of austenite in different heat treatment was determined by PHILIPS APD-10 full-automatic X-ray diffraction instrument using Co radiation from 40 to 120° at a step interval of 0.02°, and the volume fraction was calculated by the following formula, I F110 ϕA ¼ 1= 1 þ K ; K ¼ I0A111 =I0 F110 IA111
ð1Þ
Where I0A111 and I0F110 are pure austenite and pure ferrite integrated intensity, respectively. The morphology and distribution of reversed austenite and the precipitation phase of the samples were investigated by Philip EM420 transmission electron microscope (TEM) and JES 2100 high resolution transmission electron microscope (HRTEM) after jet polishing of a thin foil in a solution of 6% perchloric acid and 94% anhydrous ethanol. The crystallographic orientation relationship between reversed austenite and martensite matrix was determined by electron diffraction patterns. Crystallographic characterization was also carried out by means of the EBSD method using FEI Quanta 650 FEG thermal field emission scanning electron microscope with HKL Channel 5. The scanning step is 0.2 μm. The results obtained by EBSD method were analyzed by means of the Orientation Imaging Microscopy (OIM™) system.
Fig. 1. X-ray diffraction spectrum of the experimental steel after quenched and tempered.
experiments, and the results show that As = 585°°C; Af = 730°°C; Ms. = 310°°C; Mf = 170°°C. Based on calculation of Fig. 1, the variation curve in the amount of austenite with tempering temperature for 2 h after quenched at 1050°°C is shown in Fig. 2. It should be noted that no austenite is detected when the tempering temperature is lower than the As temperature (585°°C), and there is a steep increase in volume fraction of austenite above 600°°C. It is believed that this part of austenite formed after tempering is reversed austenite. In general, above 550°°C, the volume fraction of austenite increases with raised temperature and it reaches a maximum value of 6.28% at 650°°C. Subsequently, the volume fraction of austenite starts to decrease. This calculated result is in accordance with the results attained by the previous studies on 15Cr SMSS [19]. According to the previous studies, it is known that during tempering, the volume fraction of the reversed austenite is controlled by two factors: the transformation volume of austenite under high temperature and the stability of austenite in tempering cooling process [19,22,23]. So there exists an optimal temperature under which the volume fraction of reversed austenite is maximum in the temperature range from As to Af.
3. Results and discussion 3.1. The volume fraction of austenite tested by XRD Fig. 1 shows the result of XRD of the steel subjected to quenching and tempering. According to the XRD pattern of quenching (black line in Fig. 1), it reveals that the content of the retained austenite is close to zero by considering the equipment and calculation errors. It means that there is no retained austenite after quenching at 1050°°C. The microstructure is uniphase martensite after quenching. Phase transition points of the experimental steel are measured by thermal expansion Table 1 Chemical compositions of the experimental steel, (wt.%). C
Mn
Si
Cr
Ni
Mo
0.019
0.40
0.16
11.69
4.94
2.04
Fig. 2. Volume fractions of austenite in the experimental steel with different tempering temperatures.
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3.2. Microstructure analysis Fig. 3(a) shows a TEM micrograph of a specimen subjected to temping at 550°°C for 2 h after quenching at 1050°°C. As shown, the specimen has a typical lath martensitic structure with darker strips. It can be observed that there is a tiny reversed austenite formed at the lath martensite grain boundary. According to the XRD results and the As temperature, we know that the martensitic reversed transformation has occurred at tempering temperature in the range from 550°°C to 600°°C, so the reversed austenite start to form at 550°°C. It is very difficult to observe such small volume of reversed austenite at 550°°C. Fig. 3(b) shows selected area electronic diffraction pattern (SADP) for Fig. 3(a). The orientation relationship between reversed austenite and lath martensite is(111)A//(011)M and ½110A ==½111M , which meets the Kurdjumov–Sachs (K–S) relationship. This orientation relationship is formed to reduce the interfacial energy of reversed austenite nucleation [23]. Previous studies [8] have shown that the special grain orientation relationship, besides chemical compositions, plays important roles on the mechanical stability of the reversed austenite. The micrograph observed by TEM of the tested steel tempered at 650°°C for 2 h after quenching at 1050°°C is show in Fig. 4(a). The morphologies of martensite matrix, the reversed austenite and fine carbides are displayed. The reversed austenite has thin sheets and platelets with ~100 nm in the martensite matrix. According to the electronic diffraction pattern for selected area shown in Fig. 4(b), the carbides can be determined as M23C6 by calibration. This carbide has a FCC lattice with space group Fm3m, which has 92 metal atoms and 24 carbon atoms in each unit cell. Fig. 4(a) displays that M23C6 carbides are formed on the M/γ interfaces. The nucleation energy of M23C6 can be calculated using the following equation [24], ΔG ¼ α1
σ3 ðΔGV þ εÞ2
:
ð2Þ
Where α1 is the geometry coefficient, σ is specific interfacial energy, ΔGV and ε are the free enthalpy variation and the elastic strain energy per unit volume of parent phase, respectively. In our case, theoretically, there are three main types of interface: the martensite–martensite interface (α/α), the new reversed austenite– martensite interface (α/γ) and the internal austenite–austenite interface (γ/γ). However, the austenite–austenite interface (γ/γ) can be neglected here, since it does not appear as reversed austenite distributes and disperses in the martensite matrix. It is known that martensite grain boundaries differ in terms of their misorientation and energy [25,26]. Most importantly, we showed the reversed austenite and the martensite meet the K-S orientation relationship. According the study by D. Raabe et al. [25], the interfacial energies of the martensite-austenite grain boundaries are generally assumed to be equal to or lower than
high angle martensite–martensite grain boundaries. And their research results indicated that the reversed austenite–martensite interface (α/γ) with K-S orientation relationship always has the lower interface energy. Therefore, because σα/γis lower than σα/α, the nucleation energy of M23C6 at M/γ interface is lower than at α/α interface. It could be confirmed that M23C6 is easy to form at the M/γ interface. Fig. 4 displays the precipitated M23C6 carbides have a FCC crystal structure and obey the cube-on-cube orientation relationship with the reversed austenite: [111]γ//[111]M23C6 and ð220Þγ==ð220ÞM23 C6 , which means the M23C6 carbides have the same orientation with γ, and this will contribute to phase transformation by reducing strain energy. In Fig. 4(b), we can also see that there are three diffraction spots of M23C6 between two diffraction spots of γ. This is obviously because that M23C6 has the same crystal structure with γ and its crystal lattice parameter is three times of γ [27,28]. Fig. 5 shows TEM image of reversed austenite in a specimen subjected to tempering at 650°°C for 2 h after quenching at 1050°°C. As shown in Fig. 5(a), the darker reversed austenite platelet is formed at the region with high dislocation density, and it is growing along lath martensite boundary. M/γ interface gradually migrates to lath martensite, and the reversed austenite swallows the lath martensite and continues to grow. It is found that the M/γ interface is rectilinear and nearly parallel to the martensite lath. This indicates that the nucleation and growth of the reversed austenite follow the certain orientation relationship. Fig. 5 (b) shows a thin sheet reversed austenite with some stacking faults. The stacking faults could provide more nucleation sites for reversed austenite, and contribute to the revered transformation by reducing the energy barrier that need to overcome the shear deformation. Fig. 5(b) also reveals that the dislocation density in the reversed austenite is obviously lower than that in the martensite matrix, which indicates that the reversed austenite does not inherit the high dislocation density structure of prior martensite matrix, and this is the indirect evidence that the austenite is formed by diffusion [19]. Regarding the precipitation of M23C6 and the reversed austenite, various groups have studied the nucleation and growth of the abutting phases as well as the crystallographic orientations. S.H. Zhang et al. [29] detected that the nucleation sequence of M23C6 and reversed austenite is difficult to be clarified. Y.Y. Song et al. [28] indicated that M23C6 carbides precipitated at the prior martensitic lath boundaries supplied the nucleation sites for the reversed austenite formation and supplied Ni-enrichment in the adjacent region. H.U. Hong et al. [30] reported the M23C6 carbide tends to select the preferential planes which have the lowest interfacial energy. H. Li et al. [31] observed that the atoms near the interface of carbide and matrix will redistribute during the growth of M23C6, and the Cr atoms migrate into the carbide, while the Fe and Ni atoms migrate into the matrix. From the above discussion, on the one hand, the formation of reversed austenite not only provides the nucleation sites of M23C6 at
Fig. 3. TEM micrographs of reversed austenite along the lath boundary in the experimental steel tempered at 550°°C for 2 h: (a) bright field and (b) SAD pattern.
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Fig. 4. TEM images of the experimental steel tempered at 650°°C after quenching at 1050°°C: (a) bright field; (b) the corresponding SAED pattern; (c) schematic diagram of (b).
the new formed K-S austenite–martensite interface (α/γ) with low interface energy, but also promotes the high equilibrium segregation of Cr to the boundary, leading to a high chemical driving force to the nucleation of M23C6. On the other hand, the M23C6 carbides precipitated at M/γ interface can cause the decrease of interface energy and the formation of the chromium depleted region and nickel enrichment zone. Then the diffusion of Ni in the adjacent area becomes easier. Thus the M23C6 carbides precipitation provides the energy of the nucleation, and leads to the formation and growth of reversed austenite. These indicate that M23C6 precipitation and the formation of reversed austenite have the interaction effect. 3.3. EBSD analysis In order to further clarify the orientation relationship of reversed austenite and martensite, the experimental steel tempering at 650°°C for 2 h is measured by using EBSD. Fig. 6(a) shows the SEM photo of the flat surface of sample by electrolytic polishing. Fig. 6(b) displays EBSD inverse pole figure color map of the area in Fig. 6(a). In Fig. 6(b), white color represents reversed austenite, and different other colors represent different crystallographic orientation martensites. Fig. 6(c) is the superposition graph which consists of martensite quality image and inverse pole figure color map of reversed austenite in Fig. 6(a). Fig. 6(d) is the superposition graph which consists of boundary map and inverse pole figure color map of the reversed austenite in Fig. 6(a). The dispersions of reversed austenite with different colors are shown in Fig. 6(c) and 6(d). In these maps, black
solid lines follow all kinds of boundaries, including prior austenite grain boundaries, packet boundaries and lath boundaries. The bcc/fcc map is shown in the Fig. 6(e), and the red color represents reversed austenite (fcc), and the white color represents the martensite (bcc). It clearly reveals the reversed austenites distribute in the martensite matrix. As shown in Fig. 6, there are 8 prior austenite grains in the 75∗75 μm area. Each prior austenite grain contains more than one packet (the martensite lath collection with same habit plane), and each packet contains more than one martensite block (the martensite lath collection with approaching orientation), which has the similar color and orientation sub laths. Fig. 7 represents the result of interface distribution analysis in the martensite structure. According to Fig. 6 (d) and Fig. 7, the packet and block boundaries are the interface which misorientation is more than 10°, and the block boundary is mostly twin boundary with misorientation of about 60°. The misorientation of sub-block boundary is 5 ~ 10°, and the martensite misorientation maintains a statistically random distribution. These maps reveal that reversed austenite mostly forms at the interior of martensite block, not only on the sub-block boundaries, but also on the lath boundaries. Only a few reversed austenites form on prior austenite grain boundaries, packet boundaries and block boundaries. Based on the results of Fig. 7, reversed austenite is concentrated in the martensitic microstructure phase interface with the misorientation range of 0° ~ 60°. In order to improve orientation analysis, one prior austenite grain area which contains 4 martensite packets in the upper of Fig. 8(a) and
Fig. 5. TEM micrographs of reversed austenite in the experimental steel tempered at 650°°C for 2 h.
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Fig. 6. Crystallographic orientation maps of the experimental steel tempered at 650°°C for 2 h. (a) SEM micrograph of an area in the experimental steel; (b) inverse pole figure of the martensite matrix and reversed austenite of (a); (c) inverse pole figure of the reversed austenite; (d) boundary map; (e) the bcc/fcc map.
8(b) is studied. Fig. 8(c) and 8(d) are {100} pole figures showing orientations of the lath martensite and reversed austenite corresponding to the EBSD orientation map, respectively. In Fig. 8(c), it can be concluded that the martensite orientation relationship conforms to the K-S relationship variant orientation. The three open circle symbols (○) in Fig. 8(c) and 8(d) represent {100} orientation projection of the prior austenite grain. Fig. 8(d) also shows the green symbols and three open circle symbols (○) coincide with each other. This means that the “green” reversed austenite (the green symbols) has the same orientation with the prior austenite. Whereas the prior austenite presents the K-S orientation relationship with martensite matrix, the “green” reversed austenite presents the identical orientation
Fig. 7. Interface distribution in the martensite structure.
relationship with martensite matrix. And the close packed plane and direction of the two are parallel to each other, just like (011)α//(111)γ and ½111α==½101γ. It can be concluded that the nucleation and growth of the “green” reversed austenite, which is strictly limited by crystallographic orientation, is the reversible and nondiffusional shear deformation process (i.e., the shear reversion mechanism) [32]. This kind of reversed austenite transforms from bcc into fcc by two times shearing. It is found that when the reversed austenite grains are formed on the lath boundary plane, the α′/γ interfacial energy can be minimized and the activation energy of nucleation can be reduced. The close packed plane which is nearly parallel to the lath boundary is chose as the habit plane of the nucleation of the reversed austenite. That is to say, the habit plane is the lath boundary plane which is close to {011}α. {110} and {111} pole figures for all reversed austenite are shown in Fig. 8(e). Compared with Fig. 8(d), reversed austenites with other colors have essentially the same orientation with the surrounding martensite, and their distribution characteristic also meet with K-S polar diagram pattern. However, there are different close packed plane and different close packed direction between the other color reversed austenite with the “green” reversed austenite. It means that there is no habit plane among the other color reversed austenite and surrounding martensite matrix. Therefore we consider that the nucleation position and growing mechanism of the other color reversed austenite are different with the “green” reversed austenite. In our previous studies [19], we deeply discussed the growing mechanism of this kind of reversed austenite of a series of SMSSs. It is a diffused phase transformation, and the growth of these reversed austenites is closely related to the diffusion process of Ni. As determined in Fig. 8, although all reversed austenite which are inside the prior austenite have different crystallographic orientation relationships, the reversed austenite variants are limited to two kinds within one packet and five kinds within one prior austenite grain. That is, they are in agreement with “the austenite limitation” which many studies carried out [32].
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Fig. 8. Crystallographic orientation maps and pole figures of one prior austenite. (a) Inverse pole figure of the martensite matrix and reversed austenite; (b) inverse pole figure of reversed austenite; (c){100} pole figure showing orientations of the lath martensite corresponding to the EBSD orientation map; (d) {100} pole figure showing orientations of the reversed austenite; (e) {110} and {111} pole figure of reversed austenite.
4. Conclusions Experiments were carried out aiming to determine the influence of M23C6 on the formation of reversed austenite and crystallographic orientation of reversed austenite of SMSS, and the main conclusions of the work can be listed as follows: 1. M23C6 is main carbide precipitated at tempering. M23C6 has the same crystallographic orientation relationship with reversed austenite, so it provides a nucleation condition for the formation of reversed austenite. The M23C6 precipitation promotes the formation of the reversed
austenite during tempering process in SMSS. The M23C6 precipitation and the formation of the reversed austenite have the interaction effect during tempering process. 2. With a nucleation and growth mechanism, the reversed austenite forms within the prior austenite grain uniformly. Especially, it forms intensively at the sub-block boundary and the lath boundary within a misorientation range of 0–60°. Two kinds of austenite variants with different orientation relationships can be formed inside a packet. 3. There are two different kinds of formation modes for reversed austenite. One is the nondiffusional shear reversion; the other is the
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