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Precipitation of σ-phase in high-nitrogen austenitic 18Cr–18Mn–2Mo–0.9N stainless steel during isothermal aging
Scripta Materialia 50 (2004) 1325–1328 www.actamat-journals.com
Precipitation of r-phase in high-nitrogen austenitic 18Cr–18Mn–2Mo–0.9N stainless steel during isothermal aging Tae-Ho Lee a
a,*
, Chang-Seok Oh a, Chang Gil Lee a, Sung-Joon Kim a, Setsuo Takaki
b
Korea Institute of Machinery and Materials, Materials Processing Department, 66 Sangnam, Changwon 641-010, South Korea b Kyushu University, 6-10-1 Hakozaki, Higashi-ku, Fukuoka 812-8581, Japan Received 28 August 2003; received in revised form 5 February 2004; accepted 10 February 2004
Abstract Precipitation mechanism of r-phase in austenitic high-nitrogen steel (HNS) upon aging was investigated using electron microscopy and thermodynamic calculation. The precipitation sequence was found to be grain boundary Cr2 N, cellular Cr2 N and r-phase with aging time. It was concluded that the nitrogen-depleted zone near Cr2 N induces the nucleation of r-phase. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Stainless steels; Precipitation; CALPHAD; Nitrides; Sigma (r)-phase
1. Introduction The beneficial effects of nitrogen on the properties of high-alloyed steels have led to a widespread development of HNS owing to recent advances in processing technologies [1]. Nitrogen in solid solution is a beneficial alloying element to increase strength level without significant loss of ductility and toughness. Nitrogen is also a strong austenite-stabilizing element, and improves the resistance to localized corrosion [2,3]. Therefore, austenitic HNS constitute a group of promising structural materials that possess a favorable combination of mechanical and corrosion resistance properties. The formation mechanism of r-phase has been the subject of many investigations, motivated by its detrimental effects on toughness and corrosion resistance [4– 9]. Weiss and Stickler [6] suggested that the r-phase precipitated in the following sequence: carbide precipitation in c matrix, phase transformation from c to a, and the r-phase nucleation at the expense of a. Beckitt [7] reported that the r-phase formed directly from c after precipitation of M23 C6 carbide. However, in other studies [8,9], the formation of r-phase is thought to depend on the (Cr + Ni) contents in steel. Singhal and Martin [8] postulated that, in steel with (Cr + Ni) > *
Corresponding author. Tel.: +82-55-280-3434; fax: +82-55-2803498. E-mail address: [email protected] (T.-H. Lee).
47wt.%, the precipitation of a at grain boundaries preceded the formation of r-phase, and the r-phase was found to be transformed in-situ from a, whereas, at lower (Cr + Ni) contents, the r-phase precipitated directly from c, partially at the expense of M23 C6 . Later, Barcik [9] envisaged that the formation of r-phase was dependent on the stability of austenite, i.e., the Creq /Nieq ratio (Creq ¼ Cr + Mo + 1.5 Si + 0.5 Nb; Nieq ¼ Ni + 39 C + 26 N + 0.5 Mn) of the c, instead of simple (Cr + Ni) criterion, and the carbon content is one of the controlling factors determining the susceptibility to the r-phase precipitation. Although several mechanisms were proposed as previously mentioned, it is deemed that r-phase formation would be strongly dependent on chemistry of alloys and precipitation kinetics in the high-nitrogen stainless steel might be different from those in low-nitrogen or nitrogen-free stainless steels. Therefore, microstructural characterization and thermodynamic calculation were attempted to elucidate the formation mechanism of the r-phase in austenitic high-nitrogen stainless steel.
2. Experimental The investigated material was a commercial highnitrogen austenitic P900NMo alloy (manufactured by VSG, Germany) with the following composition in wt.%: 17.94 Cr; 18.60 Mn; 2.09 Mo; 0.89 N; 0.04 C;
1359-6462/$ - see front matter Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2004.02.013
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balance Fe. Specimens were encapsulated into an evacuated quartz tube and solution-treated at 1150 °C for 30 min, followed by water quenching. Then, they were isothermally aged at 900 °C up to 168 h in argon atmosphere, and quenched into water. The isothermal aging at 900 °C was chosen because the nose temperature in TTP diagram was found to be around 900 °C, which was also in accordance with the literature [10]. The SEM observations were carried out on a JSM6700F operating at 15 kV after chemical etching in a Glyceregia reagent. Thin foils for TEM were prepared by means of twin-jet electrolytic polishing technique using an electrolyte of 15% perchloric acid in methanol. They were examined in a JEM 2010 operating at 200 kV. Thermodynamic calculations using Thermo-Calc databank [11] with TCFE2000 database were performed to obtain information on phase equilibrium and volume fractions of constituent phases at 900 °C. 3. Results and discussion 3.1. Precipitation behavior Aging the solution-treated specimens at 900 °C caused, at first, intergranular as well as cellular precipitation of M2 (C,N) (hereafter designated as Cr2 N). Fig. 1 shows a series of SEM micrographs of the specimen aged at 900 °C for various aging times. As shown in the figure, nucleation and growth of Cr2 N precipitate along
grain boundaries was noticeable in the sample aged about 103 s at 900 °C, and further aging of 105 s and resulted in complete precipitation of coverage the grain boundaries. After a certain incubation time of 103 s, the cellular precipitation of Cr2 N started from grain boundaries, and the volume fraction of cellular Cr2 N precipitates increased with increasing aging time. Although most of precipitated particles were Cr2 N, a small amount of r-phase also precipitated along grain and cell boundaries on prolonged aging at 900 °C. The morphology of the r-phase was a coarse and irregular blocky shape with darker contrast compared to Cr2 N, and the mean size as well as the volume fraction of rphase increased with aging time. 3.2. Thermodynamic calculation Fig. 2(a) and (b) are calculated isopleths for the Fe– 18Cr–18Mn–2Mo–N alloy and the variation of the distribution of phases in equilibrium with nitrogen content at 900 °C, respectively. The vertical (dashed) line denotes the chemical composition of the alloy investigated, and the horizontal (solid) line indicates the isothermal aging temperature of the present study. In the high-temperature region, the c single-phase field is enclosed with two-phase fields, d þ c and c þ Cr2 N. At 900 °C, the narrow field for c þ Cr2 N þ r in low-nitrogen content, and the wide field for c þ Cr2 N in high-nitrogen content coexist. In the low-temperature region, M23 C6 and M6 C carbides are stable phases, and are in equilibrium with c þ Cr2 N þ r and c þ Cr2 N, respectively. Information of stability of the c single-phase is important because it allows for the heat-treatment window (solution treatment, hot-rolling, etc.) to be estimated. In this study, the solution treatment was carried out at 1150 °C, based on thermodynamic calculation together with microstructure characterization. The chemical composition of the present steel falls inside the c þ Cr2 N two-phase field at 900 °C, indicating that the only stable precipitate with c is the Cr2 N phase under equilibrium condition. However, the composition of the decomposed c matrix may fall well down the c þ Cr2 N þ r=c þ Cr2 N boundary because of the nitrogen-depletion near the Cr2 N precipitation. Therefore, as predicted from Fig. 2, the precipitation of rphase can occur if the composition of the decomposed c shifts into the three-phase region for c þ Cr2 N þ r. 3.3. Electron microscopy
Fig. 1. SEM micrographs of the specimens aged at 900 °C for various times: (a) 102 s; (b) 103 s; (c) 104 s; (d) 105 s; (e) 72 h; (f) 168 h (r-phases are indicated by arrows).
The concentration distribution of nitrogen across the intergranular as well as cellular Cr2 N precipitates was measured several times in SEM. One of the typical concentration profiles is shown in Fig. 3 for the specimen aged at 900 °C for 168 h. The dashed line denotes the trace of the concentration profile of nitrogen, and
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Fig. 2. The results of thermodynamic calculations: (a) calculated isopleth of the Fe–18Cr–18Mn–2Mo–N alloy system, and (b) calculated isothermal section at 900 °C, illustrating the formation of Cr2 N, r-phase and phase separation of a and c, respectively.
Fig. 3. (a) SEM micrograph of the Cr2 N precipitates for the specimen aged at 900 °C for 168 h, and (b) EDS line profile of nitrogen across the Cr2 N.
the solid line indicates the reference position of the Cr2 N. Fig. 3(a) and (b) show the secondary electron
(SE) image of the Cr2 N precipitate, and line profile of nitrogen across the Cr2 N. The nitrogen-enriched region in Cr2 N precipitate, and the nitrogen-depleted region adjacent to the Cr2 N precipitate are clearly shown in Fig. 3(b). The cellular precipitation of Cr2 N is known to occur by the decomposition of the supersaturated austenite matrix (c1 ) to a lamellar structure consisting of an austenite matrix with equilibrated composition (c2 ) and Cr2 N, i.e., c1 ! c2 þ Cr2 N. The chemical composition of the c and Cr2 N phases, as determined by thermodynamic calculation, is given in Table 1. Comparing Fig. 2 and Table 1, the equilibrium content of nitrogen in c adjacent the Cr2 N precipitates is 0.26 wt.% after the precipitation of Cr2 N, and the nitrogen concentration of phase boundary for c þ Cr2 N þ r=c þ Cr2 N at 900 °C is calculated to be 0.70 wt.%. Thus, the composition of the decomposed c matrix shifts into the c þ Cr2 N þ r three-phase region because of the nitrogen-depletion near the precipitation of Cr2 N, and the r-phase can precipitate in this nitrogen-depleted region at the c/Cr2 N boundary. Thus, the cell boundaries, where the diffusion of nitrogen, chromium and molybdenum is faster, are thought to be the preferential sites for the nucleation of r-phase. The r-phase formation has been reported to occur in c near discontinuous precipitation of chromium nitride colonies [10,12]. Because r-phase has no solubility of nitrogen, its precipitation should occur in the nitrogenimpoverished regions. The precipitation of the r-phase must be evidence of the nitrogen gradient formation
Table 1 The calculated composition of the austenite and M2 (C,N) phases at 900 °C (unit: wt.%) Nominal Austenite M2 (C,N)
Cr
Mn
Mo
Si
C
N
17.94 15.04 55.55
18.60 18.54 19.39
2.09 1.93 4.12
0.77 0.83 –
0.038 0.032 0.111
0.890 0.257 9.104
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tion of r-phase at the c/Cr2 N boundary, as predicted in Fig. 2. Therefore, the formation of r-phase can be explained, on the bases of thermodynamic calculation and electron microscopy, by the mechanism that the formation of a nitrogen-depleted zone near the Cr2 N phase induced nucleation of the r-phase.
4. Conclusion During the isothermal aging at 900 °C, the Cr2 N phase precipitated at grain boundaries, followed by cellular precipitation. On prolonged aging, coarse rphase precipitated along grain and cell boundaries. The formation of r-phase could be explained by the mechanism that the formation of a nitrogen-depleted zone near the Cr2 N phase induced nucleation of the r-phase, which was supported by electron microscopy and thermodynamic calculation. Fig. 4. TEM micrographs of the specimen aged at 900 °C for 168 h: (a) BF image; (b) CDF image using (3 3 0)r reflection showing r-phase formed at the interface between the Cr2 N and c; (c) SAD pattern (z ¼ ð½ 111c ==½110r ); (d) schematic illustration of (c).
near the reaction front during discontinuous precipitation. A possible reason for the nitrogen gradient is the great difference between the diffusion coefficient of nitrogen and that of chromium in the matrix [12]. Therefore, the important factor of r-phase formation in high-nitrogen steels can be considered to be the formation of nitrogen-depleted zone near the chromium nitride colonies. Moreover, in commercial austenitic stainless steels, the formation of r-phase has been attributed to the precipitation of M23 C6 carbide, which implies that the formation of carbon-depleted zone is the important factor of r-phase formation rather than other alloying elements. Fig. 4(a) through (d) show BF image, centered dark field (CDF) image using ( 3 3 0)r reflection, SADP, and computer-simulated SADP of the specimen aged at 900 °C for 168 h. The r-phase has been reported to be a tetragonal unit cell (c=a ¼ 0:52, 30 atoms per unit cell) of space group P 42 =mnm with a ¼ 0:880 and c ¼ 0:454 nm [5,9]. Fig. 4(b) shows the direct evidence of the forma-
Acknowledgements One of author (Tae-Ho Lee) would like to thank to Dr. T. Tsuchiyama in Kyushu University for helpful discussion, and acknowledge the financial support by Japan Society for the Promotion of Science through the JSPS Dissertation Ph.D. program.
References [1] Gavriljuk VG, Berns H. High nitrogen steels. Berlin: SpringerVerlag; 1999. [2] Simmons JW. Mater Sci Eng A 1996;207:159. [3] H€annien H, Romu J, Ilola R, Tervo J, Laitinen A. J Mater Sci Tech 2001;117:424. [4] Lee TH, Kim SJ. Scripta Mater 1998;39:951. [5] Lee TH, Kim SJ, Jung YC. Metall Trans A 2000;31:1713. [6] Weiss B, Stickler R. Metall Trans A 1972;3:851. [7] Beckitt FR. J Iron Steel Inst 1969;207:632. [8] Singhan LK, Martin JW. Acta Metall 1968;16:1441. [9] Barcik J. Mater Sci Tech 1988;4:5. [10] Vanderschaeve F, Taillard R, Foct J. J Mater Sci 1995;30:6035. [11] Anderson JO, Helander T, H€ oglund L, Sundman B. CALPHAD 2002;26:273. [12] Machado IF, Padilha AF. ISIJ Inter 2000;40:719.
Precipitation of r-phase in high-nitrogen austenitic 18Cr–18Mn–2Mo–0.9N stainless steel during isothermal aging Tae-Ho Lee a
a,*
, Chang-Seok Oh a, Chang Gil Lee a, Sung-Joon Kim a, Setsuo Takaki
b
Korea Institute of Machinery and Materials, Materials Processing Department, 66 Sangnam, Changwon 641-010, South Korea b Kyushu University, 6-10-1 Hakozaki, Higashi-ku, Fukuoka 812-8581, Japan Received 28 August 2003; received in revised form 5 February 2004; accepted 10 February 2004
Abstract Precipitation mechanism of r-phase in austenitic high-nitrogen steel (HNS) upon aging was investigated using electron microscopy and thermodynamic calculation. The precipitation sequence was found to be grain boundary Cr2 N, cellular Cr2 N and r-phase with aging time. It was concluded that the nitrogen-depleted zone near Cr2 N induces the nucleation of r-phase. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Stainless steels; Precipitation; CALPHAD; Nitrides; Sigma (r)-phase
1. Introduction The beneficial effects of nitrogen on the properties of high-alloyed steels have led to a widespread development of HNS owing to recent advances in processing technologies [1]. Nitrogen in solid solution is a beneficial alloying element to increase strength level without significant loss of ductility and toughness. Nitrogen is also a strong austenite-stabilizing element, and improves the resistance to localized corrosion [2,3]. Therefore, austenitic HNS constitute a group of promising structural materials that possess a favorable combination of mechanical and corrosion resistance properties. The formation mechanism of r-phase has been the subject of many investigations, motivated by its detrimental effects on toughness and corrosion resistance [4– 9]. Weiss and Stickler [6] suggested that the r-phase precipitated in the following sequence: carbide precipitation in c matrix, phase transformation from c to a, and the r-phase nucleation at the expense of a. Beckitt [7] reported that the r-phase formed directly from c after precipitation of M23 C6 carbide. However, in other studies [8,9], the formation of r-phase is thought to depend on the (Cr + Ni) contents in steel. Singhal and Martin [8] postulated that, in steel with (Cr + Ni) > *
Corresponding author. Tel.: +82-55-280-3434; fax: +82-55-2803498. E-mail address: [email protected] (T.-H. Lee).
47wt.%, the precipitation of a at grain boundaries preceded the formation of r-phase, and the r-phase was found to be transformed in-situ from a, whereas, at lower (Cr + Ni) contents, the r-phase precipitated directly from c, partially at the expense of M23 C6 . Later, Barcik [9] envisaged that the formation of r-phase was dependent on the stability of austenite, i.e., the Creq /Nieq ratio (Creq ¼ Cr + Mo + 1.5 Si + 0.5 Nb; Nieq ¼ Ni + 39 C + 26 N + 0.5 Mn) of the c, instead of simple (Cr + Ni) criterion, and the carbon content is one of the controlling factors determining the susceptibility to the r-phase precipitation. Although several mechanisms were proposed as previously mentioned, it is deemed that r-phase formation would be strongly dependent on chemistry of alloys and precipitation kinetics in the high-nitrogen stainless steel might be different from those in low-nitrogen or nitrogen-free stainless steels. Therefore, microstructural characterization and thermodynamic calculation were attempted to elucidate the formation mechanism of the r-phase in austenitic high-nitrogen stainless steel.
2. Experimental The investigated material was a commercial highnitrogen austenitic P900NMo alloy (manufactured by VSG, Germany) with the following composition in wt.%: 17.94 Cr; 18.60 Mn; 2.09 Mo; 0.89 N; 0.04 C;
1359-6462/$ - see front matter Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2004.02.013
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T.-H. Lee et al. / Scripta Materialia 50 (2004) 1325–1328
balance Fe. Specimens were encapsulated into an evacuated quartz tube and solution-treated at 1150 °C for 30 min, followed by water quenching. Then, they were isothermally aged at 900 °C up to 168 h in argon atmosphere, and quenched into water. The isothermal aging at 900 °C was chosen because the nose temperature in TTP diagram was found to be around 900 °C, which was also in accordance with the literature [10]. The SEM observations were carried out on a JSM6700F operating at 15 kV after chemical etching in a Glyceregia reagent. Thin foils for TEM were prepared by means of twin-jet electrolytic polishing technique using an electrolyte of 15% perchloric acid in methanol. They were examined in a JEM 2010 operating at 200 kV. Thermodynamic calculations using Thermo-Calc databank [11] with TCFE2000 database were performed to obtain information on phase equilibrium and volume fractions of constituent phases at 900 °C. 3. Results and discussion 3.1. Precipitation behavior Aging the solution-treated specimens at 900 °C caused, at first, intergranular as well as cellular precipitation of M2 (C,N) (hereafter designated as Cr2 N). Fig. 1 shows a series of SEM micrographs of the specimen aged at 900 °C for various aging times. As shown in the figure, nucleation and growth of Cr2 N precipitate along
grain boundaries was noticeable in the sample aged about 103 s at 900 °C, and further aging of 105 s and resulted in complete precipitation of coverage the grain boundaries. After a certain incubation time of 103 s, the cellular precipitation of Cr2 N started from grain boundaries, and the volume fraction of cellular Cr2 N precipitates increased with increasing aging time. Although most of precipitated particles were Cr2 N, a small amount of r-phase also precipitated along grain and cell boundaries on prolonged aging at 900 °C. The morphology of the r-phase was a coarse and irregular blocky shape with darker contrast compared to Cr2 N, and the mean size as well as the volume fraction of rphase increased with aging time. 3.2. Thermodynamic calculation Fig. 2(a) and (b) are calculated isopleths for the Fe– 18Cr–18Mn–2Mo–N alloy and the variation of the distribution of phases in equilibrium with nitrogen content at 900 °C, respectively. The vertical (dashed) line denotes the chemical composition of the alloy investigated, and the horizontal (solid) line indicates the isothermal aging temperature of the present study. In the high-temperature region, the c single-phase field is enclosed with two-phase fields, d þ c and c þ Cr2 N. At 900 °C, the narrow field for c þ Cr2 N þ r in low-nitrogen content, and the wide field for c þ Cr2 N in high-nitrogen content coexist. In the low-temperature region, M23 C6 and M6 C carbides are stable phases, and are in equilibrium with c þ Cr2 N þ r and c þ Cr2 N, respectively. Information of stability of the c single-phase is important because it allows for the heat-treatment window (solution treatment, hot-rolling, etc.) to be estimated. In this study, the solution treatment was carried out at 1150 °C, based on thermodynamic calculation together with microstructure characterization. The chemical composition of the present steel falls inside the c þ Cr2 N two-phase field at 900 °C, indicating that the only stable precipitate with c is the Cr2 N phase under equilibrium condition. However, the composition of the decomposed c matrix may fall well down the c þ Cr2 N þ r=c þ Cr2 N boundary because of the nitrogen-depletion near the Cr2 N precipitation. Therefore, as predicted from Fig. 2, the precipitation of rphase can occur if the composition of the decomposed c shifts into the three-phase region for c þ Cr2 N þ r. 3.3. Electron microscopy
Fig. 1. SEM micrographs of the specimens aged at 900 °C for various times: (a) 102 s; (b) 103 s; (c) 104 s; (d) 105 s; (e) 72 h; (f) 168 h (r-phases are indicated by arrows).
The concentration distribution of nitrogen across the intergranular as well as cellular Cr2 N precipitates was measured several times in SEM. One of the typical concentration profiles is shown in Fig. 3 for the specimen aged at 900 °C for 168 h. The dashed line denotes the trace of the concentration profile of nitrogen, and
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Fig. 2. The results of thermodynamic calculations: (a) calculated isopleth of the Fe–18Cr–18Mn–2Mo–N alloy system, and (b) calculated isothermal section at 900 °C, illustrating the formation of Cr2 N, r-phase and phase separation of a and c, respectively.
Fig. 3. (a) SEM micrograph of the Cr2 N precipitates for the specimen aged at 900 °C for 168 h, and (b) EDS line profile of nitrogen across the Cr2 N.
the solid line indicates the reference position of the Cr2 N. Fig. 3(a) and (b) show the secondary electron
(SE) image of the Cr2 N precipitate, and line profile of nitrogen across the Cr2 N. The nitrogen-enriched region in Cr2 N precipitate, and the nitrogen-depleted region adjacent to the Cr2 N precipitate are clearly shown in Fig. 3(b). The cellular precipitation of Cr2 N is known to occur by the decomposition of the supersaturated austenite matrix (c1 ) to a lamellar structure consisting of an austenite matrix with equilibrated composition (c2 ) and Cr2 N, i.e., c1 ! c2 þ Cr2 N. The chemical composition of the c and Cr2 N phases, as determined by thermodynamic calculation, is given in Table 1. Comparing Fig. 2 and Table 1, the equilibrium content of nitrogen in c adjacent the Cr2 N precipitates is 0.26 wt.% after the precipitation of Cr2 N, and the nitrogen concentration of phase boundary for c þ Cr2 N þ r=c þ Cr2 N at 900 °C is calculated to be 0.70 wt.%. Thus, the composition of the decomposed c matrix shifts into the c þ Cr2 N þ r three-phase region because of the nitrogen-depletion near the precipitation of Cr2 N, and the r-phase can precipitate in this nitrogen-depleted region at the c/Cr2 N boundary. Thus, the cell boundaries, where the diffusion of nitrogen, chromium and molybdenum is faster, are thought to be the preferential sites for the nucleation of r-phase. The r-phase formation has been reported to occur in c near discontinuous precipitation of chromium nitride colonies [10,12]. Because r-phase has no solubility of nitrogen, its precipitation should occur in the nitrogenimpoverished regions. The precipitation of the r-phase must be evidence of the nitrogen gradient formation
Table 1 The calculated composition of the austenite and M2 (C,N) phases at 900 °C (unit: wt.%) Nominal Austenite M2 (C,N)
Cr
Mn
Mo
Si
C
N
17.94 15.04 55.55
18.60 18.54 19.39
2.09 1.93 4.12
0.77 0.83 –
0.038 0.032 0.111
0.890 0.257 9.104
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tion of r-phase at the c/Cr2 N boundary, as predicted in Fig. 2. Therefore, the formation of r-phase can be explained, on the bases of thermodynamic calculation and electron microscopy, by the mechanism that the formation of a nitrogen-depleted zone near the Cr2 N phase induced nucleation of the r-phase.
4. Conclusion During the isothermal aging at 900 °C, the Cr2 N phase precipitated at grain boundaries, followed by cellular precipitation. On prolonged aging, coarse rphase precipitated along grain and cell boundaries. The formation of r-phase could be explained by the mechanism that the formation of a nitrogen-depleted zone near the Cr2 N phase induced nucleation of the r-phase, which was supported by electron microscopy and thermodynamic calculation. Fig. 4. TEM micrographs of the specimen aged at 900 °C for 168 h: (a) BF image; (b) CDF image using (3 3 0)r reflection showing r-phase formed at the interface between the Cr2 N and c; (c) SAD pattern (z ¼ ð½ 111c ==½110r ); (d) schematic illustration of (c).
near the reaction front during discontinuous precipitation. A possible reason for the nitrogen gradient is the great difference between the diffusion coefficient of nitrogen and that of chromium in the matrix [12]. Therefore, the important factor of r-phase formation in high-nitrogen steels can be considered to be the formation of nitrogen-depleted zone near the chromium nitride colonies. Moreover, in commercial austenitic stainless steels, the formation of r-phase has been attributed to the precipitation of M23 C6 carbide, which implies that the formation of carbon-depleted zone is the important factor of r-phase formation rather than other alloying elements. Fig. 4(a) through (d) show BF image, centered dark field (CDF) image using ( 3 3 0)r reflection, SADP, and computer-simulated SADP of the specimen aged at 900 °C for 168 h. The r-phase has been reported to be a tetragonal unit cell (c=a ¼ 0:52, 30 atoms per unit cell) of space group P 42 =mnm with a ¼ 0:880 and c ¼ 0:454 nm [5,9]. Fig. 4(b) shows the direct evidence of the forma-
Acknowledgements One of author (Tae-Ho Lee) would like to thank to Dr. T. Tsuchiyama in Kyushu University for helpful discussion, and acknowledge the financial support by Japan Society for the Promotion of Science through the JSPS Dissertation Ph.D. program.
References [1] Gavriljuk VG, Berns H. High nitrogen steels. Berlin: SpringerVerlag; 1999. [2] Simmons JW. Mater Sci Eng A 1996;207:159. [3] H€annien H, Romu J, Ilola R, Tervo J, Laitinen A. J Mater Sci Tech 2001;117:424. [4] Lee TH, Kim SJ. Scripta Mater 1998;39:951. [5] Lee TH, Kim SJ, Jung YC. Metall Trans A 2000;31:1713. [6] Weiss B, Stickler R. Metall Trans A 1972;3:851. [7] Beckitt FR. J Iron Steel Inst 1969;207:632. [8] Singhan LK, Martin JW. Acta Metall 1968;16:1441. [9] Barcik J. Mater Sci Tech 1988;4:5. [10] Vanderschaeve F, Taillard R, Foct J. J Mater Sci 1995;30:6035. [11] Anderson JO, Helander T, H€ oglund L, Sundman B. CALPHAD 2002;26:273. [12] Machado IF, Padilha AF. ISIJ Inter 2000;40:719.