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Grain refinement of Cr25Ni5Mo1.5 duplex stainless steel by heat treatment
Materials Science and Engineering A363 (2003) 263–267
Grain refinement of Cr25Ni5Mo1.5 duplex stainless steel by heat treatment Zh.L. Jiang∗ , X.Y. Chen, H. Huang, X.Y. Liu Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China Received 6 December 2002; received in revised form 14 July 2003
Abstract A technology employing transformation (␣ → ␥ + ) and its reverse transformation (␥ + → ␣) to refine grain of duplex stainless steels has been devised. The experimental results indicated that the phases nucleate and grow both at the ␣/␥ phase interfaces and in the interior of the ␣ phases when the temperature was higher, whereas the phases nucleate at the ␣/␥ phase interfaces primarily when a lower temperature was used. The coarse ␣ grain decomposed into fine ␥ and grains at 750 ◦ C for 38 h (the treated), and fine ␣ grains were obtained after a revise transformation (␥ + → ␣) being performed, leading to an significant improved toughness. The highest impact toughness value was achieved by the process that transformation was carried out at 750 ◦ C for 38 h followed by a reverse transformation, heated at 1200 ◦ C for 1.5 min, at a heating rate of 13 ◦ C/s. © 2003 Elsevier B.V. All rights reserved. Keywords: Duplex stainless steel; Grain refinement and phase transformation
1. Introduction Ferrite/austenite duplex stainless steels are widely used in industries of chemistry, food, paper, pharmacy, marine and many other fields [1–4] due to their high performances. Duplex stainless steels have higher strength than austenite stainless steels, higher toughness than ferrite stainless steels, good welding ability, and high resistance to stress corrosion cracking, hydrogen embrittlement and intergranular corrosion. These noted properties of duplex stainless steels result from a unique microstructure configuration. The microstructure of ␣/␥ duplex stainless steels formed by a proper hot-working is a mixture of fine ferrite and austenite phases, referred as micro-duplex structure. However, an improper hot working causes coarse mixture of ␣ and ␥ phases that reduces the strength and toughness of the steel. This undesired coarse microstructure is difficult to be detected by an ultrasonic technology. It is therefore, important to develop a technology to refine the coarse grains of ␣/␥ duplex stainless steels, especially to refine coarse ␣ grains. ∗ Corresponding author. Tel.: +86-10-62772620; fax: +86-10-62771160. E-mail address: mmm [email protected] (Zh.L. Jiang).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0921-5093(03)00647-6
It is commonly known that alloying, cold deformation followed by a recrystallization, and solid phase transformation techniques are frequently used to refine grains of metallic materials. However, for a component or part made of duplex stainless steels, the chemical composition and the hot working process performed cannot be modified, a phase transformation by a heat treatment is only practical approach to refine grains. In fact, grain refinement by heat treatment is a common practice in production and R&D of metallic materials. However, it has not been reported that the grain refinement, by a heat treatment, of ␣/␥ duplex stainless steels has been processed. In the present study, grain refinement of a Cr25 Ni5 Mo1.5 ␣/␥ duplex stainless steels via a transformation of ␣ → + ␥ and its reverse transformation of + ␥ → ␣ was investigated. Also the precipitation characteristics of phase, the evolution of microstructure and impact toughness of the steel were studied.
2. Experiments The nominal chemical composition of the steel used in the present study is shown in Table 1, and the relevant phase diagram is shown in Fig. 1 [5]. The steel was initially heat
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Table 1 The nominal composition of the materials used (wt.%) Element
Content
C Mn Si P S Ni Cr Mo N Fe
0.12 0.64 0.97 0.031 0.011 5.5 25.8 1.7 0.08 Balance
pact specimens was performed at 750 ◦ C for 38h, and the reverse transformation was conducted with Gleeble 1500 hot-simulating machine at 1200 ◦ C, for 1–2 min at heating rates of 13–20 ◦ C/s, followed by a quenching in water. The microstructure of the impact specimens was also examined by the optical microscopy. The impact toughness of Charpy V notched specimens, denoted as Jcvn was measured at the temperatures of −25, 0, 25, 50 and 95 ◦ C, respectively, and the corresponding fracture surface was examined by a scanning electron microscopy.
3. Results and discussion treated at 1050 ◦ C for 2 h, followed by a quenching in water and then annealed at 300 ◦ C for 4 h, followed by cooling in the furnace. To search for the optimum treatment parameters, the samples having small size of 5 mm × 5 mm × 2 mm were treated in an electrical furnace at the temperatures of 600, 650, 700, 750 and 800 ◦ C, for a holding time between 10 and 110 h. A transformation of ␣ → + ␥ is expected to takes place during the heat treatment applied [6–8]. The microstructure of the samples that were ground mechanically, followed by a polish electrically in a solution of (25 g KClO4 + 500 g glacial acetic acid) with a current density of 0.2 A/mm2 for 90 s, then etched in a solution of hydrochloric acid saturated by FeCl3 , was examined by an optical microscopy. The volume fraction of phase was analyzed by X-ray diffraction. As a result, the optimum parameter for transformation was determined. The transformation and its reverse transformation for the impact specimens were processed. According to the results achieved with the samples having small size, the transformation for the im-
3.1. Microstructure before and after the σ transformation The microstructures before and after the transformation are shown in Fig. 2a–d, respectively. As expected, the heat treatment causes a significant refinement of the ␣ phases, in other words, the transformation of ␣ → + ␥ likely take places in the ␣ phases. It is found that the site of the phase precipitation is closely related to the transformation temperature. As shown in Fig. 3, the phases nucleate both at the ␣/␥ phase interfaces and in the interior of the ␣ phases when the temperature is higher, whereas the ␣ phases nucleate at the ␣/␥ phase interfaces primarily when a lower temperature is used. The precipitation of phase is mainly controlled by the diffusion of phases promoting elements Cr and Mo. At a higher temperature, the diffusion of Cr and Mo feasibly performs along phase interfaces and in the interior of ␣ phases, thus the phases can nucleate both at the phase interfaces and in the interior of the ␣ phases. The diffusive coefficient of Cr and Mo decreases as temperature, the formation of the phases in the interior of ␣ phases is blocked, and consequently the phases precipitate at the ␣/␥ phase interfaces because of high diffusive coefficient along the interfaces. After being heat treated for a long time, the phases coarsen substantially, as shown in Fig. 2d. The relationship between the temperature and the time at which the phase starts to coarsen is established and results are indicated in Fig. 4, the higher the temperature, the faster the coarsening is. 3.2. Volume fraction of σ phase
Fig. 1. The quasi-binary phase diagram of 70% Fe duplex stainless steel.
Dependence of the volume fraction of the phase on the temperature and time of heat treatment is exhibited in Fig. 5. Regardless of the temperature, the volume fraction of the phase increases quickly at the beginning of the treatment, and approaches a saturation value gradually. The temperature influences the change rate of volume fraction at the beginning of the treatment, as well as the saturation value of the volume fraction. The fastest change rate and the highest volume fraction of the phase are achieved at 750 ◦ C, while both the higher and lower temperature decrease the efficient rate and product rate of the transformation. This
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Fig. 2. (a) The microstructure of original sample, (b) the microstructure of sample after treatment at 700 ◦ C for 38 h, (c) the microstructure of sample after treatment at 750 ◦ C for 38 h and (d) the microstructure of sample after σ treatment at 800 ◦ C for 100 h.
result shows the nose of the C curve of the transformation is near 750 ◦ C, so the treatment at 750 ◦ C will spend the shortest time and can get the most complete transformation of ␣ → + ␥. Meanwhile it was shown that the microstructures of the sample treated at 750 ◦ C is finer than that at 800 ◦ C and is as fine as that at 700 ◦ C. It is logical that the finer the microstructure of the sample and the higher the volume fraction of phase after transformation, the finer the microstructure of the sample after reverse transformation. Since sample treated at 750 ◦ C has finer microstructure, the highest volume fraction of phase and expense the shortest treatment time, so all specimens for the impact testing were treated at 750 ◦ C for 38 h.
3.3. Microstructure of impact specimens after the reverse treatment The variables of microstructure for the impact specimens are listed in Table 2. The results indicate that resultant grain size of the specimen is dependent on an individual process, even though all processes used can refine the ␣ phase. The process A results in a fine ␣ phase, the process B produces a very fine and the process C generates an ultra ␣ phase, shown in Fig. 6a and b. The average grain sizes of various processes are also listed in Table 2. The difference between the process A and B is the different temperature used for the solution treatment, the former is 1300 ◦ C, and the latter is
Fig. 3. (a) The microstructure of sample after treatment at 800 ◦ C for 10 h and (b) the microstructure of sample after treatment at 700 ◦ C for 10 h.
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Fig. 4. Coarsening velocity of phase.
Fig. 7. Impact toughness of Cr25 Ni5 Mo1.5 treated by different processes.
influences the refinement of the ␣ phase. X-ray diffraction analysis indicates that a certain amount of the phase is remained after being processed by the process C, and the phase can be eliminated primarily through the process B. 3.4. Toughness and fracture surface
Fig. 5. The content of phase after different heat treatments.
1050 ◦ C. The lower solution treatment temperature results in small initial grains, subsequent the finer ␣ phase can be obtained after the reverse treatment. Different heating rates were used for the process B and C, the former is 13 ◦ C/s, and the latter is 20 ◦ C/s. Therefore, the heating rate considerably
Toughness of the specimens treated by various processes is shown in Fig. 7. As a result of the refinement treatment, impact toughness increases significantly. In the testing temperature range, toughness of the coarse-grained specimen is relatively low, and no plateau in the temperature–toughness curve appears. For the process A, a significant increase in the toughness is observed, however, the increase occurs above the room temperature only. The process B causes a considerable increase in the toughness throughout the testing temperature range. The process C can leads to an increase below the room temperature, like the process B. However, the process C cannot increase the toughness pronouncedly above the room temperature compared with the process A and B. Therefore, it seems that the solution temperature, that is, the initial grain size of the ␣ phase primarily affects the toughness below the room temperature, whereas the one above the
Fig. 6. (a) The microstructure of specimen treated by process B and (b) the microstructure of specimen treated by process C.
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Table 2 Heat treatments and the results of Cr25 Ni5 Mo1.5 Process number Solid solution
treatment
Anti- treatment
Average grain size (m)
Heating rate (◦ C) Heating temperature (◦ C) Holding time (min) A B C
1300 ◦ C for 120 min 750 ◦ C for 38 h 13 1050 ◦ C for 120 min 750 ◦ C for 38 h 13 1050 ◦ C for 120 min 750 ◦ C for 38 h 20
1200 1200 1200
1.5 1.5 1.5
16.0 10.7 6.1
Fig. 8. The morphology of impact fracture surface of (a) the original specimen, impacted at 50 ◦ C and (b) the specimen after treated by process B, impacted at 50 ◦ C.
room temperature is mainly influenced by the refined grain size of the ␣ phase. The residual phase, due to an incomplete reverse transformation is very harmful to the impact toughness above the room temperature. SEM examination shows a significant difference in the fracture surface of the specimens treated by different processes. Generally, the fracture surface of the impact specimens consists of cleavage and dimple feature. The coarse-grained specimen has a larger cleavage area, and the fine-grained specimen has an increased dimple area, i.e. the fracture surface of the initial specimen tested at 50 ◦ C shows big cleavage feature primarily, whereas that of the refined specimen processed by the process B exhibits fine dimples completely, shown in Fig. 8a and b, respectively. As a conclusion, the duplex stainless steel changes its fracture mechanism from cleavage to dimple due to grain refinement, subsequently the impact toughness increases significantly.
4. Conclusion 1. The ␣ ferrite in a duplex stainless steel decomposed into ␥ and phases during a transformation. The desired transformation, which caused the ␣ phase to decompose into fine mixture of ␥ and phases at a high rate,-
occurred during an isothermal treatment of 750 ◦ C for 38 h. 2. The coarse ␣ ferrite was refined via a transformation followed its reverse transformation. This grain refinement led to an increase in fracture surface of the duplex stainless steel, subsequent a significant increase in the impact toughness. According to the measured impact toughness value in a temperature range from −25 to 95 ◦ C, the reverse transformation was optimized as a process of heating specimens to 1200 ◦ C at a heating rate of 15 ◦ C/s for 1.5 min.
References [1] Y.H. Park, Z.H. Lee, Mater. Sci. Eng. A 297 (2001) 78–84. [2] H. Eriksson, S. Bernhardsson, Corrosion 47 (1991) 719–725. [3] J.K.L. Lai, K.W. Wong, D.J. Li, Mater. Sci. Eng. A 203 (1995) 356– 364. [4] K.M. Lee, H.S. Cho, D.C. Choi, J. Alloy Comp. 285 (1999) 156– 161. [5] M. WU, Duplex Stainless Steel, Metallurgy Book Press, Beijing, 1999, p. 9 (in Chinese). [6] C.H. Shek, K.W. Wong, J.K.L. Lai, D.J. Li, Mater. Sci. Eng. A 231 (1997) 42–47. [7] Y.S. Sato, H. Kokawa, Scripta Materialia 40 (1999) 659–663. [8] T.H. Chen, J.R. Yang, Mater. Sci. Eng. A 311 (2001) 28–41.
Grain refinement of Cr25Ni5Mo1.5 duplex stainless steel by heat treatment Zh.L. Jiang∗ , X.Y. Chen, H. Huang, X.Y. Liu Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, PR China Received 6 December 2002; received in revised form 14 July 2003
Abstract A technology employing transformation (␣ → ␥ + ) and its reverse transformation (␥ + → ␣) to refine grain of duplex stainless steels has been devised. The experimental results indicated that the phases nucleate and grow both at the ␣/␥ phase interfaces and in the interior of the ␣ phases when the temperature was higher, whereas the phases nucleate at the ␣/␥ phase interfaces primarily when a lower temperature was used. The coarse ␣ grain decomposed into fine ␥ and grains at 750 ◦ C for 38 h (the treated), and fine ␣ grains were obtained after a revise transformation (␥ + → ␣) being performed, leading to an significant improved toughness. The highest impact toughness value was achieved by the process that transformation was carried out at 750 ◦ C for 38 h followed by a reverse transformation, heated at 1200 ◦ C for 1.5 min, at a heating rate of 13 ◦ C/s. © 2003 Elsevier B.V. All rights reserved. Keywords: Duplex stainless steel; Grain refinement and phase transformation
1. Introduction Ferrite/austenite duplex stainless steels are widely used in industries of chemistry, food, paper, pharmacy, marine and many other fields [1–4] due to their high performances. Duplex stainless steels have higher strength than austenite stainless steels, higher toughness than ferrite stainless steels, good welding ability, and high resistance to stress corrosion cracking, hydrogen embrittlement and intergranular corrosion. These noted properties of duplex stainless steels result from a unique microstructure configuration. The microstructure of ␣/␥ duplex stainless steels formed by a proper hot-working is a mixture of fine ferrite and austenite phases, referred as micro-duplex structure. However, an improper hot working causes coarse mixture of ␣ and ␥ phases that reduces the strength and toughness of the steel. This undesired coarse microstructure is difficult to be detected by an ultrasonic technology. It is therefore, important to develop a technology to refine the coarse grains of ␣/␥ duplex stainless steels, especially to refine coarse ␣ grains. ∗ Corresponding author. Tel.: +86-10-62772620; fax: +86-10-62771160. E-mail address: mmm [email protected] (Zh.L. Jiang).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0921-5093(03)00647-6
It is commonly known that alloying, cold deformation followed by a recrystallization, and solid phase transformation techniques are frequently used to refine grains of metallic materials. However, for a component or part made of duplex stainless steels, the chemical composition and the hot working process performed cannot be modified, a phase transformation by a heat treatment is only practical approach to refine grains. In fact, grain refinement by heat treatment is a common practice in production and R&D of metallic materials. However, it has not been reported that the grain refinement, by a heat treatment, of ␣/␥ duplex stainless steels has been processed. In the present study, grain refinement of a Cr25 Ni5 Mo1.5 ␣/␥ duplex stainless steels via a transformation of ␣ → + ␥ and its reverse transformation of + ␥ → ␣ was investigated. Also the precipitation characteristics of phase, the evolution of microstructure and impact toughness of the steel were studied.
2. Experiments The nominal chemical composition of the steel used in the present study is shown in Table 1, and the relevant phase diagram is shown in Fig. 1 [5]. The steel was initially heat
264
Zh.L. Jiang et al. / Materials Science and Engineering A363 (2003) 263–267
Table 1 The nominal composition of the materials used (wt.%) Element
Content
C Mn Si P S Ni Cr Mo N Fe
0.12 0.64 0.97 0.031 0.011 5.5 25.8 1.7 0.08 Balance
pact specimens was performed at 750 ◦ C for 38h, and the reverse transformation was conducted with Gleeble 1500 hot-simulating machine at 1200 ◦ C, for 1–2 min at heating rates of 13–20 ◦ C/s, followed by a quenching in water. The microstructure of the impact specimens was also examined by the optical microscopy. The impact toughness of Charpy V notched specimens, denoted as Jcvn was measured at the temperatures of −25, 0, 25, 50 and 95 ◦ C, respectively, and the corresponding fracture surface was examined by a scanning electron microscopy.
3. Results and discussion treated at 1050 ◦ C for 2 h, followed by a quenching in water and then annealed at 300 ◦ C for 4 h, followed by cooling in the furnace. To search for the optimum treatment parameters, the samples having small size of 5 mm × 5 mm × 2 mm were treated in an electrical furnace at the temperatures of 600, 650, 700, 750 and 800 ◦ C, for a holding time between 10 and 110 h. A transformation of ␣ → + ␥ is expected to takes place during the heat treatment applied [6–8]. The microstructure of the samples that were ground mechanically, followed by a polish electrically in a solution of (25 g KClO4 + 500 g glacial acetic acid) with a current density of 0.2 A/mm2 for 90 s, then etched in a solution of hydrochloric acid saturated by FeCl3 , was examined by an optical microscopy. The volume fraction of phase was analyzed by X-ray diffraction. As a result, the optimum parameter for transformation was determined. The transformation and its reverse transformation for the impact specimens were processed. According to the results achieved with the samples having small size, the transformation for the im-
3.1. Microstructure before and after the σ transformation The microstructures before and after the transformation are shown in Fig. 2a–d, respectively. As expected, the heat treatment causes a significant refinement of the ␣ phases, in other words, the transformation of ␣ → + ␥ likely take places in the ␣ phases. It is found that the site of the phase precipitation is closely related to the transformation temperature. As shown in Fig. 3, the phases nucleate both at the ␣/␥ phase interfaces and in the interior of the ␣ phases when the temperature is higher, whereas the ␣ phases nucleate at the ␣/␥ phase interfaces primarily when a lower temperature is used. The precipitation of phase is mainly controlled by the diffusion of phases promoting elements Cr and Mo. At a higher temperature, the diffusion of Cr and Mo feasibly performs along phase interfaces and in the interior of ␣ phases, thus the phases can nucleate both at the phase interfaces and in the interior of the ␣ phases. The diffusive coefficient of Cr and Mo decreases as temperature, the formation of the phases in the interior of ␣ phases is blocked, and consequently the phases precipitate at the ␣/␥ phase interfaces because of high diffusive coefficient along the interfaces. After being heat treated for a long time, the phases coarsen substantially, as shown in Fig. 2d. The relationship between the temperature and the time at which the phase starts to coarsen is established and results are indicated in Fig. 4, the higher the temperature, the faster the coarsening is. 3.2. Volume fraction of σ phase
Fig. 1. The quasi-binary phase diagram of 70% Fe duplex stainless steel.
Dependence of the volume fraction of the phase on the temperature and time of heat treatment is exhibited in Fig. 5. Regardless of the temperature, the volume fraction of the phase increases quickly at the beginning of the treatment, and approaches a saturation value gradually. The temperature influences the change rate of volume fraction at the beginning of the treatment, as well as the saturation value of the volume fraction. The fastest change rate and the highest volume fraction of the phase are achieved at 750 ◦ C, while both the higher and lower temperature decrease the efficient rate and product rate of the transformation. This
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265
Fig. 2. (a) The microstructure of original sample, (b) the microstructure of sample after treatment at 700 ◦ C for 38 h, (c) the microstructure of sample after treatment at 750 ◦ C for 38 h and (d) the microstructure of sample after σ treatment at 800 ◦ C for 100 h.
result shows the nose of the C curve of the transformation is near 750 ◦ C, so the treatment at 750 ◦ C will spend the shortest time and can get the most complete transformation of ␣ → + ␥. Meanwhile it was shown that the microstructures of the sample treated at 750 ◦ C is finer than that at 800 ◦ C and is as fine as that at 700 ◦ C. It is logical that the finer the microstructure of the sample and the higher the volume fraction of phase after transformation, the finer the microstructure of the sample after reverse transformation. Since sample treated at 750 ◦ C has finer microstructure, the highest volume fraction of phase and expense the shortest treatment time, so all specimens for the impact testing were treated at 750 ◦ C for 38 h.
3.3. Microstructure of impact specimens after the reverse treatment The variables of microstructure for the impact specimens are listed in Table 2. The results indicate that resultant grain size of the specimen is dependent on an individual process, even though all processes used can refine the ␣ phase. The process A results in a fine ␣ phase, the process B produces a very fine and the process C generates an ultra ␣ phase, shown in Fig. 6a and b. The average grain sizes of various processes are also listed in Table 2. The difference between the process A and B is the different temperature used for the solution treatment, the former is 1300 ◦ C, and the latter is
Fig. 3. (a) The microstructure of sample after treatment at 800 ◦ C for 10 h and (b) the microstructure of sample after treatment at 700 ◦ C for 10 h.
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Fig. 4. Coarsening velocity of phase.
Fig. 7. Impact toughness of Cr25 Ni5 Mo1.5 treated by different processes.
influences the refinement of the ␣ phase. X-ray diffraction analysis indicates that a certain amount of the phase is remained after being processed by the process C, and the phase can be eliminated primarily through the process B. 3.4. Toughness and fracture surface
Fig. 5. The content of phase after different heat treatments.
1050 ◦ C. The lower solution treatment temperature results in small initial grains, subsequent the finer ␣ phase can be obtained after the reverse treatment. Different heating rates were used for the process B and C, the former is 13 ◦ C/s, and the latter is 20 ◦ C/s. Therefore, the heating rate considerably
Toughness of the specimens treated by various processes is shown in Fig. 7. As a result of the refinement treatment, impact toughness increases significantly. In the testing temperature range, toughness of the coarse-grained specimen is relatively low, and no plateau in the temperature–toughness curve appears. For the process A, a significant increase in the toughness is observed, however, the increase occurs above the room temperature only. The process B causes a considerable increase in the toughness throughout the testing temperature range. The process C can leads to an increase below the room temperature, like the process B. However, the process C cannot increase the toughness pronouncedly above the room temperature compared with the process A and B. Therefore, it seems that the solution temperature, that is, the initial grain size of the ␣ phase primarily affects the toughness below the room temperature, whereas the one above the
Fig. 6. (a) The microstructure of specimen treated by process B and (b) the microstructure of specimen treated by process C.
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Table 2 Heat treatments and the results of Cr25 Ni5 Mo1.5 Process number Solid solution
treatment
Anti- treatment
Average grain size (m)
Heating rate (◦ C) Heating temperature (◦ C) Holding time (min) A B C
1300 ◦ C for 120 min 750 ◦ C for 38 h 13 1050 ◦ C for 120 min 750 ◦ C for 38 h 13 1050 ◦ C for 120 min 750 ◦ C for 38 h 20
1200 1200 1200
1.5 1.5 1.5
16.0 10.7 6.1
Fig. 8. The morphology of impact fracture surface of (a) the original specimen, impacted at 50 ◦ C and (b) the specimen after treated by process B, impacted at 50 ◦ C.
room temperature is mainly influenced by the refined grain size of the ␣ phase. The residual phase, due to an incomplete reverse transformation is very harmful to the impact toughness above the room temperature. SEM examination shows a significant difference in the fracture surface of the specimens treated by different processes. Generally, the fracture surface of the impact specimens consists of cleavage and dimple feature. The coarse-grained specimen has a larger cleavage area, and the fine-grained specimen has an increased dimple area, i.e. the fracture surface of the initial specimen tested at 50 ◦ C shows big cleavage feature primarily, whereas that of the refined specimen processed by the process B exhibits fine dimples completely, shown in Fig. 8a and b, respectively. As a conclusion, the duplex stainless steel changes its fracture mechanism from cleavage to dimple due to grain refinement, subsequently the impact toughness increases significantly.
4. Conclusion 1. The ␣ ferrite in a duplex stainless steel decomposed into ␥ and phases during a transformation. The desired transformation, which caused the ␣ phase to decompose into fine mixture of ␥ and phases at a high rate,-
occurred during an isothermal treatment of 750 ◦ C for 38 h. 2. The coarse ␣ ferrite was refined via a transformation followed its reverse transformation. This grain refinement led to an increase in fracture surface of the duplex stainless steel, subsequent a significant increase in the impact toughness. According to the measured impact toughness value in a temperature range from −25 to 95 ◦ C, the reverse transformation was optimized as a process of heating specimens to 1200 ◦ C at a heating rate of 15 ◦ C/s for 1.5 min.
References [1] Y.H. Park, Z.H. Lee, Mater. Sci. Eng. A 297 (2001) 78–84. [2] H. Eriksson, S. Bernhardsson, Corrosion 47 (1991) 719–725. [3] J.K.L. Lai, K.W. Wong, D.J. Li, Mater. Sci. Eng. A 203 (1995) 356– 364. [4] K.M. Lee, H.S. Cho, D.C. Choi, J. Alloy Comp. 285 (1999) 156– 161. [5] M. WU, Duplex Stainless Steel, Metallurgy Book Press, Beijing, 1999, p. 9 (in Chinese). [6] C.H. Shek, K.W. Wong, J.K.L. Lai, D.J. Li, Mater. Sci. Eng. A 231 (1997) 42–47. [7] Y.S. Sato, H. Kokawa, Scripta Materialia 40 (1999) 659–663. [8] T.H. Chen, J.R. Yang, Mater. Sci. Eng. A 311 (2001) 28–41.