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Bending fatigue study of an austenite-ferrite dual-phase FeAlMn steel
Materials Science and Engineering, A118 (1989) 25-31
25
Bending Fatigue Study of an Austenite-Ferrite Dual-phase Fe-AI-Mn Steel C. T. HU and C. L. TARN
Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, 30043 (Taiwan) (Received November 14, 1988; in revised form February 27, 1989)
Abstract
A high stress bending fatigue test has been conducted on a cold-rolled and annealed Fe-8.2wt. %Al-30.lwt.%Mn steel which exhibits a mixed dual-phase structure containing 45 vol. % austenite and 55 vol. % ferrite. The axis of the applied stress is either perpendicular (70 to or parallel (L) to the rolling direction of specimens. It is found that the tensile property of the L specimens is much superior to that of the T specimens, but their fatigue properties are comparable. Transmission electron microscopy (TEM) examinations indicate that a heterogeneous planar dislocation substructure was developed by the cyclic bending stress within the austenite region, and the tangled dislocation cell substructure was observed within the ferrite region. The former is associated with the early persistent slip bands and the easy initiation and propagation of crack. However, the latter is accompanied by the cross-slip of dislocations. From the in situ surface crack propagation study, it is found that the crack is always initiated from and moved along the Slip bands within the austenite region. The twin and grain boundaries, and austenite-ferrite interfaces are less favourable. The fatigue crack path cannot durably stay within the ferrite region (arrays) of the bending fatigue specimen. This" result implies that the ferrite phase serves as a crack arrester. One reason is that the annealed ferrite phase is softer than austenite and dislocation easily cross-slips in the ferrite structure, so that the blunting effect during fatigue testing may retard the crack propagation. The other important reason for the better fatigue properties o f the ferrite phase than the austenite is the shortage of carbon in the present material. I. Introduction
The study of fatigue of single- and multi-phase materials is an important metallurgical subject 0921-5093/89/$3.50
[1-8]. Several approaches to the problems of fatigue resistance have been reported in the literature [9]. Such concepts as slip-mode, stacking fault energy, homogenization of slip planes and microstructures have been studied by many people [9-13]. It has been inferred that the optimal microstructures to improve the crack resistance generally exhibit homogeneous slip bands, minimal inclusions, and lack of grain boundaries transverse to the axis of stress. The comparison of fatigue behaviours of several carbon steels has been investigated by Bathias and Pelloux [13]. They found that the low stacking fault energy rather than the different crystal structure can give rise to a better resistance against fatigue crack propagation. Another study of a two-phase alloy by rotating beam fatigue testing has been reported by Hayden and Floreen [14]. They observed that the volume fraction ratios of constituents to achieve the maximum fatigue strength are quite different. For example, the austenite-ferrite series (IN-744 heat) with almost equal amounts of the two phases exhibited the best fatigue strength, and both the martensite-ferrite and iron-copper series achieved the strength with composition near one end of the tie-line. However, Hayden and Floreen [14] did not compare the fatigue resistances of individual phases in their dual-phase alloys. In addition, the chemical compositions of steel in the study of Bathias and Pelloux [13] were quite different. Moreover, the carbon would mostly dissolve into the austenite phase which caused a solid solution strengthening effect to austenite. This condition raises another problem in relation to the difference of fatigue resistances of various phases. A new family of Fe-A1-Mn and Fe-A1-Mn-C alloys have drawn much attention in recent years. It is believed to be a potential candidate as substitutes for the conventional Fe-Cr-Ni stainless steels [15-19]. The partial equilibrium diagram of © Elsevier Sequoia/Printed in The Netherlands
26
Fe-A1-Mn has been studied previously [20]. By using this diagram, a mixed dual-phase structure with equal amounts of austenite and ferrite regions could be produced by an adequate combination of aluminum and manganese. The present study concentrates on the examination of fatigue crack resistances of various phases in one alloy and tries to correlate with the transmission electron microscopy (TEM) observations. A high stress bending fatigue testing has been adopted to make the crack examination convenient. In this investigation the studied dual-phase alloy contains very few carbons which generally provide a favourable solid solution strengthening effect in the austenite phase, therefore, the results give a trustworthy comparison of the fatigue crack resistances between austenite and ferrite phases.
Both tensile (S) and bending fatigue (F) test specimens were sectioned from the cold-rolled sheet; the axis of applied stress (tensile or bending) is either parallel (L) or perpendicular (T) to the rolling direction. All specimens were then electropolished at room temperature with a solution of 5% perchloric acid and 95% acetic anhydride, operated at 20 V. Tensile properties of both SL and ST specimens were tested by an Instron machine with a strain rate of 4 × 10 - 4 s - 1. Constant stress-amplitude fatigue tests of both FL and FT specimens were carried out in completely reversed plane bending with a Sonntage model SF2-U bending fatigue machine, operating at 30 Hz and 442.6 MPa (about 80% ultimate tensile strength (UTS) of SL specimens and 110% UTS of ST specimens). To distinguish austenite from ferrite during several intermediate stages of complete fatigue testing, some specimens were lightly etched with 4% Nital before bending. The fracture and free surfaces of selected fatigue specimens were examined with both optical microscopy and scanning electron microscopy (SEM). Thin foils for TEM investigation were prepared using a jet electropolisher with an electrolyte of 90% acetic acid and 10% perchloric acid at room temperature. To study the maximal stress region in the fatigue test, i.e. the region near the free surface, one surface of the specimen was painted with dye after a slight electropolish, then the electrojet polishing process was conducted through the other surface.
2. Experimental procedure An Fe-8.2wt.%A1-30. lwt.%Mn alloy was prepared from commercially pure iron and aluminum, and electrolytic grade manganese with an L-H IS 8/III VIM furnace under a protective argon atmosphere. The chemical compositions of ingot and individual phases (obtained from induction-coupled plasma spectroscopy (ICPS) and electron probe microanalysis (EPMA)) are listed in Table 1. The detected carbon content is only 0.016 wt.% in the present alloy. A loss of approximately 1.3 wt.% Mn and 0.2 wt.% A1 (total weight per cent) has been detected during melting which is attributed to the high vapour pressure of manganese and easy oxidation of aluminum. The thickness of slabs cut from an ingot was reduced from 35 mm to 10 mm by hot forging at 1200 °C. The material was subsequently homogenized at 1100 °C for 6 h, air cooled to room temperature then subjected to cold rolling to a plate 3 mm thick. The alloy was fully annealed at 950 °C for 1.5 h followed by oil quenching to prevent any second-phase precipitation. An average grain size of about 18 # m (size ranging from 5 to 35/~m) was obtained with the Heyn intercept method.
TABLE 1
3. Results and discussions The results of EPMA indicate the composition of austenite and ferrite phase to be 58wt.%Fe-7.9wt.%A1-33.8wt.%Mn and 62.4wt.% Fe-9.4wt.%A1- 28.2wt.%Mn respectively as shown in Table 1. The mechanical properties of SL, ST, FL and FT specimens at room temperature and the microhardness in various phase regions are listed in Table 2. Both the 0.2% yield stress (YS) and UTS of SL specimens are higher than those of ST by about 140%. This is possible
Chemical compositions (in weight per cent) of ingot and individual phases
Ingot (by ICP) Austenite (by EPMA) Ferrite (by EPMA)
Fe
AI
Mn
C
Si
P
S
Balance 57.97 62.35
8.2 7.93 9.42
30.1 33.75 28.15
0.016
< 0.03
0.005
0.008
27 TABLE 2
Mechanical properties of various specimens
(a) Microhardness
Microhardness (Vickers hardness number)
Annealed austenite phase
Annealed ferrite phase
169.6 + 17.9
157.6_+11.3
(b) Tensile property Specimen
SL" ST b
UTS
0.2% YS
Elongation
RA
(MPa)
(MPa)
(%)
(%)
553 397.9
376.5 253.4
30 31
57.8 50.6
(c) Fatigue property with bending fatigue stress of 442.6 MPa Specimen
A verage fatigue life (cycles)
Standard deviation 6 offatigue life
Fig. 2. TEM microstructure of the austenite phase in the annealed condition. Twin boundaries and dispersed dislocations are observed.
(cycles) FLa FT b
61600 41800
6100 7760
axis of applied stress parallel to the rolling direction. bT, axis of applied stress perpendicular to the rolling direction. aL,
Fig. 3. TEM microstructure of the ferrite phase in the annealed condition.
Fig. 1. Optical micrograph of Fe-8.2wt.%AI-30.1wt.%Mn steel after being cold rolled and annealed for 1.5 h at 950 °C; the average grain size is 18/am.
owing to a string arrangement of austenite phase formed by the cold rolling process as shown in Fig. 1. T h e annealing twins with different crystallographic orientations are observed in the austenite region, and they are conveniently used to distinguish austenite from ferrite phases. Grange [21 ] has shown that the principles of fibre c o m p o site strengthening can be applied to a limited extent in rolling pearlite structure with a reinforcing, oriented, fibrous constituent. T h e same principle is also applied to the present material. A
volume fraction of 45% austenite and 55% ferrite was determined using the linear intercept method. Transmission electron micrographs of austenite and ferrite regions before testing are illustrated in Figs. 2 and 3 respectively. Neither the precipitate nor the Fe3A1Cx-type carbides [22, 23] were observed, which is possibly caused by the shortage of carbon in this alloy. Figure 2 shows many annealing twin boundaries in the austenite region which indicate a low stacking fault energy in this phase. It was also confirmed by Kim et al. [24]. T h e average high stress bending fatigue lives of six F L and five F T specimens are 61 600 (standard deviation 6 = 6100) and 41 800 (6 = 7760) cycles respectively. T E M micrographs of high
28
Fig; 4. T E M microstructure of the austenite phase in a fatigue-failed FL specimen. Many dislocations of planar slip mode are observed.
Fig. 6. SEM micrograph on the free surface of a fatigued FL specimen. Microcracks along slip bands within the austenite region are observed, aap = 442.6 MPa; life = 5.3 x 104 cycles.
Fig. 5. T E M mierostructure of the ferrite phase in a fatiguefailed FL specimen. Many distinct dislocation cells are observed.
stress areas in austenite and ferrite regions after fatigue testing are shown in Figs. 4 and 5. Many dislocations of planar slip-mode following (111) and (i 1 i) slip planes are observed in the austenite region as shown in Fig. 4. However, Fig. 5 displays a distinct dislocation cell substructure within the fatigued ferrite region, which illustrates the effortless cross-slip character in b.c.c, metal developed by the high amplitude cyclic bending stress. According to Grosskreutz [9], the "planar" and "wavy" slip-modes are suitable to describe the two extremes of deformation in austenite and ferrite regions of this dual-phase alloy. However, no dislocation cell within the ferrite region has been observed after the tensile testing. It is suggested that the dislocation cell substructure developed during fatigue testing prevents any local concentration of strain in the ferrite region,
Fig. 7. SEM micrograph of the free surface of a fatigued FL specimen. Penetration of a microcrack through the ferrite phase and linking with another microcrack is observed. aao = 442.6 MPa; life = 6.4 × 1 0 4 cycles.
therefore it creates a prolonged bending fatigue life of FT specimens, especially loaded with an unusually high bending stress (442.4 MPa) which is larger than the UTS (397.9 MPa) of ST specimens by a factor of about 1.1. Figures 6 and 7 show the SEM examination of the free surfaces of fatigued FL specimens. The
29 straight, clear and dense persistent slip bands are first generated all over the austenite region, then the slip bands developed in ferrite are blurry, wavy and sparse. Many microcracks mostly initiate along the persistent slip bands in austenite; however, either grain boundaries or austeniteferrite interfaces are less favourable initiation sites for fatigue cracks. There is considerable evidence that the austenite region with a low stacking fault energy exhibits less initiation resistance fatigue cracking than ferrite. Further study of the propagation of surface cracks was conducted on both FL and FT specimens. The lightly etched specimens were taken from the bending machine for crack length measurement and photographic recording after every 5000 cycles during the fatigue testing. The surface crack propagation rates of both FL and FT specimens are shown in Fig. 8, and they are comparable. Figures 9 and 10 indicate the propagation process of surface cracks on FL and FT specimens respectively. The fatigue cracks in either specimen propagated along a direction roughly perpendicular to the axis of applied stress. Figure 9 shows that the string arrangement of the austenite phase in specimen FL is perpendicular to the crack propagation, and a microcrack in the austenite region has penetrated through a neighbouring ferrite phase and linked with another microcrack. Meanwhile, no initiation of microcrack is observed within the ferrite region. When the string arrangement is parallel to the crack propagation as in specimen FT, the cracks have initiated and propagated mostly 1,5 1./+ 1.3 1.2 1.1 1.0
~ ~= g -~
along the austenite phase as shown in Fig. 10. The crack penetrated into the ferrite region on some occasions, then the propagation was impeded in the ferrite phase, and the cracks turned back into austenite again. The crack blunting phenomenon was" also observed in the ferrite region at the lower right area of Fig. 10. It is believed that the ferrite phase serves as a crack arrester in this high stress amplitude bending fatigue test which is quite contrary to most fatigue behaviours of carbon steels or alloy steels. Bathias and Pelloux [13] found that austenite steel with a low stacking fault energy displayed a superior resistance in fatigue crack propagation owing to dislocations constrained to move in a more planar fashion; thus it limited local concentration of plastic deformation and suppressed
Fig. 9. Optical micrograph of an FL specimen which has been fatigued (25 000 cycles with an applied bending stress of 442.6 MPa). A microcrack has penetrated through a neighbouring ferrite phase and linked with another microcrack in the austenite region.
o:FL specimen o:FT specimen
0.9 0.8 0.7 0.~ 0.5 0.3 O.2 03 O.C cycles ( xlO 4 )
Fig. 8. Total crack lengths of FL and FT specimens vs. fatigue cycles. The crack propagation rates of both FL and FT specimens are comparable.
Fig. 10. Optical micrograph of an FT specimen which has been fatigued (25 000 cycles with an applied bending stress of 442.6 MPa). A crack penetrated into the ferrite phase then turned back to the austenite region, meanwhile, the crack blunting phenomenon occurred within the ferrite region at the lower right area.
30
fatigue damage. Wan et al. [19, 25, 26] observed that Fe-AI-Mn steels with an austenite structure are better than those with a ferrite structure in many properties, such as strength, ductility and corrosion resistance. However, their steels contain certain amounts of elemental carbon ranging from 0.1% to 1.5%. The carbon mostly dissolves in the austenite phase, or precipitates to form carbides in the ferrite phase. This effect creates a significant solid solution strengthening in austenite, and raises both the strength and the ductility which greatly contributed to the resistance of fatigue crack initiation and propagation. However, the hard and brittle carbides in the ferrite region are detrimental to the resistance to fatigue cracking by both the stress concentration on the interfaces and the breakage of carbide particles. The present study proves that once the carbon effect is removed, the different crystal structure rather than the stacking fault energy provides the important fatigue resistance, and the ferrite phase proves to be a better crack arrester than austenite by the crack blunting effect. The better resistance character in ferrite also explains the comparable crack propagation rates between FL and FT specimens as shown in Fig. 8. Since the frequencies of crack grown into the ferrite area are comparable between those two specimens, similar impediments to crack propagation hold for each sample even though the string arrangements of constituents are completely different. Finally, there is no dislocation cell substructure observed in the ferrite region after tensile testing, and ferrite cannot play the role as a crack arrester, meanwhile, ferrite is softer than austenite; therefore, both the strength (YS and UTS) and the elongation of ST specimens are low.
4. Conclusions The extent of the string arrangement in microstructure and associated tensile and bending fatigue properties in an austenite-ferrite dualphase Fe-8.2wt.%AI-30.1wt.%Mn steel have been investigated. The results are summarized as follows. ( 1 ) The carbon plays an important role in both strengthening and fatigue resistance in the austenite phase. Once the carbon effect is removed, the ferrite phase provides better resistance to fatigue crack than austenite. (2) A well-defined dislocation cell substructure developed in the ferrite region by high ampli-
rude of cyclic stress which corresponded to the observed crack blunting phenomenon and made the character as fatigue crack arrester of the ferrite phase. (3) The low stacking fault energy in the austenite phase raised a planar deformation mode which generated less resistance to fatigue crack by dislocation pile-up and stress concentration. (4) No dislocation cell substructure was observed in the ferrite region after tensile testing; therefore, both the strength and the elongation ot ST specimens were lower than those of SL.
Acknowledgment This work was supported by the National Science Council, Taiwan.
References 1 W.J.M. Rankine, Proc. Inst. Civ. Eng., 3 (1843) 105. 2 Morris E. Fine, Metall. Trans. A, 11 (1980) 365. 3 E. E. Lauffer and W. N. Roberts, Philos. Mag., 14 (1966) 65. 4 F. Garland, N. M. Parikh and Liebowitz, Fracture, 7 (1972) 841. 5 J. Lindigkeit, A. Gysler and G. LiJtjering, Metall. Trans. A, 12(1981) 1613. 6 C. R. Aita and J. Weertman, Metall. Trans. A, 10 (1979) 535. 7 J. Awatani, K. Katagiri and T. Shiraishi, Metall. Trans. A, 7(1976) 807. 8 C. Calabrese and C. Laird, Metall. Trans., 5 (1974) 1785. 9 J.C. Grosskreutz, Metall. Trans., 3 (1972) 1255. 10 J. Awatani, K. Katagiri and K. Koyanagi, Metall. Trans. A, 10(1979) 503. 11 R.W. Landgraf, A S T M Spec. Tech. PubL, 467 (1970) 3. 12 D. K. Benson, ASTMSpec. Tech. Publ., 467(1970) 228. 13 C. Bathias and R. M. Pelloux, Metall. Trans., 4 (1973) 1265. 14 H. W. Hayden and S. Floreen, Metall. Trans., 4 (1973) 561. 15 D.J. Schmatz, Trans. Am. Soc. Met., 52 (1960) 898. 16 J. Charles, A. Berghezan, A. Lutts and P. L. Dancoisne, Met. Progr., 119 (1981) 71. 17 J. L. Ham and R. E. Carins, Prod. Eng., 29 (52) (1958) 50. 18 A. Inoue, Y. Kouima, T. Minemura and T. Masumoto, Metall. Trans. A, 12(1981) 1245. 19 C. M. Wan, H. J. Lai, M. T. Jahn, C. T. Hu and H. Heh, in N. R. Comins and J. B. Clark (eds.), Specialty Steels and Hard Metals, Pergamon, Oxford, 1983, p. 221. 20 W. Koster and W. Tonn, Arch. Eisenhuettenwes., 7(1934) 365. 21 R. A. Grange, 2nd Int. Conf. on Strength of Metal and Alloys, American Society for Metals, Metals Park, OH, 1970, p. 861.
31 22 N. A. Storchak and A. G. Drachinskaya, Phys. Met. metall., 44 (2)(1978) 123. 23 G. S. Krivonogov and M. F. Aleksegenko, Phys. Met. Metall., 4 (1975) 86. 24 Y. G. Kim, Y. S. Park and J. K. Han, Metall. Trans. A, 16 (1985) 1689.
25 C. M. Wan, S. C. Chang and Y. C. Pan, Proc. Ann. Conf.
CIMS, National Cheng Kung Univ., Tainan, Taiwan, April 15, 1984, Chinese Institute of Materials Science, 1984, p. 180. 26 J. B. Duh, W. T. Tsai and J. T. Lee, Scr. Metall., 21 (1987) 95.
25
Bending Fatigue Study of an Austenite-Ferrite Dual-phase Fe-AI-Mn Steel C. T. HU and C. L. TARN
Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, 30043 (Taiwan) (Received November 14, 1988; in revised form February 27, 1989)
Abstract
A high stress bending fatigue test has been conducted on a cold-rolled and annealed Fe-8.2wt. %Al-30.lwt.%Mn steel which exhibits a mixed dual-phase structure containing 45 vol. % austenite and 55 vol. % ferrite. The axis of the applied stress is either perpendicular (70 to or parallel (L) to the rolling direction of specimens. It is found that the tensile property of the L specimens is much superior to that of the T specimens, but their fatigue properties are comparable. Transmission electron microscopy (TEM) examinations indicate that a heterogeneous planar dislocation substructure was developed by the cyclic bending stress within the austenite region, and the tangled dislocation cell substructure was observed within the ferrite region. The former is associated with the early persistent slip bands and the easy initiation and propagation of crack. However, the latter is accompanied by the cross-slip of dislocations. From the in situ surface crack propagation study, it is found that the crack is always initiated from and moved along the Slip bands within the austenite region. The twin and grain boundaries, and austenite-ferrite interfaces are less favourable. The fatigue crack path cannot durably stay within the ferrite region (arrays) of the bending fatigue specimen. This" result implies that the ferrite phase serves as a crack arrester. One reason is that the annealed ferrite phase is softer than austenite and dislocation easily cross-slips in the ferrite structure, so that the blunting effect during fatigue testing may retard the crack propagation. The other important reason for the better fatigue properties o f the ferrite phase than the austenite is the shortage of carbon in the present material. I. Introduction
The study of fatigue of single- and multi-phase materials is an important metallurgical subject 0921-5093/89/$3.50
[1-8]. Several approaches to the problems of fatigue resistance have been reported in the literature [9]. Such concepts as slip-mode, stacking fault energy, homogenization of slip planes and microstructures have been studied by many people [9-13]. It has been inferred that the optimal microstructures to improve the crack resistance generally exhibit homogeneous slip bands, minimal inclusions, and lack of grain boundaries transverse to the axis of stress. The comparison of fatigue behaviours of several carbon steels has been investigated by Bathias and Pelloux [13]. They found that the low stacking fault energy rather than the different crystal structure can give rise to a better resistance against fatigue crack propagation. Another study of a two-phase alloy by rotating beam fatigue testing has been reported by Hayden and Floreen [14]. They observed that the volume fraction ratios of constituents to achieve the maximum fatigue strength are quite different. For example, the austenite-ferrite series (IN-744 heat) with almost equal amounts of the two phases exhibited the best fatigue strength, and both the martensite-ferrite and iron-copper series achieved the strength with composition near one end of the tie-line. However, Hayden and Floreen [14] did not compare the fatigue resistances of individual phases in their dual-phase alloys. In addition, the chemical compositions of steel in the study of Bathias and Pelloux [13] were quite different. Moreover, the carbon would mostly dissolve into the austenite phase which caused a solid solution strengthening effect to austenite. This condition raises another problem in relation to the difference of fatigue resistances of various phases. A new family of Fe-A1-Mn and Fe-A1-Mn-C alloys have drawn much attention in recent years. It is believed to be a potential candidate as substitutes for the conventional Fe-Cr-Ni stainless steels [15-19]. The partial equilibrium diagram of © Elsevier Sequoia/Printed in The Netherlands
26
Fe-A1-Mn has been studied previously [20]. By using this diagram, a mixed dual-phase structure with equal amounts of austenite and ferrite regions could be produced by an adequate combination of aluminum and manganese. The present study concentrates on the examination of fatigue crack resistances of various phases in one alloy and tries to correlate with the transmission electron microscopy (TEM) observations. A high stress bending fatigue testing has been adopted to make the crack examination convenient. In this investigation the studied dual-phase alloy contains very few carbons which generally provide a favourable solid solution strengthening effect in the austenite phase, therefore, the results give a trustworthy comparison of the fatigue crack resistances between austenite and ferrite phases.
Both tensile (S) and bending fatigue (F) test specimens were sectioned from the cold-rolled sheet; the axis of applied stress (tensile or bending) is either parallel (L) or perpendicular (T) to the rolling direction. All specimens were then electropolished at room temperature with a solution of 5% perchloric acid and 95% acetic anhydride, operated at 20 V. Tensile properties of both SL and ST specimens were tested by an Instron machine with a strain rate of 4 × 10 - 4 s - 1. Constant stress-amplitude fatigue tests of both FL and FT specimens were carried out in completely reversed plane bending with a Sonntage model SF2-U bending fatigue machine, operating at 30 Hz and 442.6 MPa (about 80% ultimate tensile strength (UTS) of SL specimens and 110% UTS of ST specimens). To distinguish austenite from ferrite during several intermediate stages of complete fatigue testing, some specimens were lightly etched with 4% Nital before bending. The fracture and free surfaces of selected fatigue specimens were examined with both optical microscopy and scanning electron microscopy (SEM). Thin foils for TEM investigation were prepared using a jet electropolisher with an electrolyte of 90% acetic acid and 10% perchloric acid at room temperature. To study the maximal stress region in the fatigue test, i.e. the region near the free surface, one surface of the specimen was painted with dye after a slight electropolish, then the electrojet polishing process was conducted through the other surface.
2. Experimental procedure An Fe-8.2wt.%A1-30. lwt.%Mn alloy was prepared from commercially pure iron and aluminum, and electrolytic grade manganese with an L-H IS 8/III VIM furnace under a protective argon atmosphere. The chemical compositions of ingot and individual phases (obtained from induction-coupled plasma spectroscopy (ICPS) and electron probe microanalysis (EPMA)) are listed in Table 1. The detected carbon content is only 0.016 wt.% in the present alloy. A loss of approximately 1.3 wt.% Mn and 0.2 wt.% A1 (total weight per cent) has been detected during melting which is attributed to the high vapour pressure of manganese and easy oxidation of aluminum. The thickness of slabs cut from an ingot was reduced from 35 mm to 10 mm by hot forging at 1200 °C. The material was subsequently homogenized at 1100 °C for 6 h, air cooled to room temperature then subjected to cold rolling to a plate 3 mm thick. The alloy was fully annealed at 950 °C for 1.5 h followed by oil quenching to prevent any second-phase precipitation. An average grain size of about 18 # m (size ranging from 5 to 35/~m) was obtained with the Heyn intercept method.
TABLE 1
3. Results and discussions The results of EPMA indicate the composition of austenite and ferrite phase to be 58wt.%Fe-7.9wt.%A1-33.8wt.%Mn and 62.4wt.% Fe-9.4wt.%A1- 28.2wt.%Mn respectively as shown in Table 1. The mechanical properties of SL, ST, FL and FT specimens at room temperature and the microhardness in various phase regions are listed in Table 2. Both the 0.2% yield stress (YS) and UTS of SL specimens are higher than those of ST by about 140%. This is possible
Chemical compositions (in weight per cent) of ingot and individual phases
Ingot (by ICP) Austenite (by EPMA) Ferrite (by EPMA)
Fe
AI
Mn
C
Si
P
S
Balance 57.97 62.35
8.2 7.93 9.42
30.1 33.75 28.15
0.016
< 0.03
0.005
0.008
27 TABLE 2
Mechanical properties of various specimens
(a) Microhardness
Microhardness (Vickers hardness number)
Annealed austenite phase
Annealed ferrite phase
169.6 + 17.9
157.6_+11.3
(b) Tensile property Specimen
SL" ST b
UTS
0.2% YS
Elongation
RA
(MPa)
(MPa)
(%)
(%)
553 397.9
376.5 253.4
30 31
57.8 50.6
(c) Fatigue property with bending fatigue stress of 442.6 MPa Specimen
A verage fatigue life (cycles)
Standard deviation 6 offatigue life
Fig. 2. TEM microstructure of the austenite phase in the annealed condition. Twin boundaries and dispersed dislocations are observed.
(cycles) FLa FT b
61600 41800
6100 7760
axis of applied stress parallel to the rolling direction. bT, axis of applied stress perpendicular to the rolling direction. aL,
Fig. 3. TEM microstructure of the ferrite phase in the annealed condition.
Fig. 1. Optical micrograph of Fe-8.2wt.%AI-30.1wt.%Mn steel after being cold rolled and annealed for 1.5 h at 950 °C; the average grain size is 18/am.
owing to a string arrangement of austenite phase formed by the cold rolling process as shown in Fig. 1. T h e annealing twins with different crystallographic orientations are observed in the austenite region, and they are conveniently used to distinguish austenite from ferrite phases. Grange [21 ] has shown that the principles of fibre c o m p o site strengthening can be applied to a limited extent in rolling pearlite structure with a reinforcing, oriented, fibrous constituent. T h e same principle is also applied to the present material. A
volume fraction of 45% austenite and 55% ferrite was determined using the linear intercept method. Transmission electron micrographs of austenite and ferrite regions before testing are illustrated in Figs. 2 and 3 respectively. Neither the precipitate nor the Fe3A1Cx-type carbides [22, 23] were observed, which is possibly caused by the shortage of carbon in this alloy. Figure 2 shows many annealing twin boundaries in the austenite region which indicate a low stacking fault energy in this phase. It was also confirmed by Kim et al. [24]. T h e average high stress bending fatigue lives of six F L and five F T specimens are 61 600 (standard deviation 6 = 6100) and 41 800 (6 = 7760) cycles respectively. T E M micrographs of high
28
Fig; 4. T E M microstructure of the austenite phase in a fatigue-failed FL specimen. Many dislocations of planar slip mode are observed.
Fig. 6. SEM micrograph on the free surface of a fatigued FL specimen. Microcracks along slip bands within the austenite region are observed, aap = 442.6 MPa; life = 5.3 x 104 cycles.
Fig. 5. T E M mierostructure of the ferrite phase in a fatiguefailed FL specimen. Many distinct dislocation cells are observed.
stress areas in austenite and ferrite regions after fatigue testing are shown in Figs. 4 and 5. Many dislocations of planar slip-mode following (111) and (i 1 i) slip planes are observed in the austenite region as shown in Fig. 4. However, Fig. 5 displays a distinct dislocation cell substructure within the fatigued ferrite region, which illustrates the effortless cross-slip character in b.c.c, metal developed by the high amplitude cyclic bending stress. According to Grosskreutz [9], the "planar" and "wavy" slip-modes are suitable to describe the two extremes of deformation in austenite and ferrite regions of this dual-phase alloy. However, no dislocation cell within the ferrite region has been observed after the tensile testing. It is suggested that the dislocation cell substructure developed during fatigue testing prevents any local concentration of strain in the ferrite region,
Fig. 7. SEM micrograph of the free surface of a fatigued FL specimen. Penetration of a microcrack through the ferrite phase and linking with another microcrack is observed. aao = 442.6 MPa; life = 6.4 × 1 0 4 cycles.
therefore it creates a prolonged bending fatigue life of FT specimens, especially loaded with an unusually high bending stress (442.4 MPa) which is larger than the UTS (397.9 MPa) of ST specimens by a factor of about 1.1. Figures 6 and 7 show the SEM examination of the free surfaces of fatigued FL specimens. The
29 straight, clear and dense persistent slip bands are first generated all over the austenite region, then the slip bands developed in ferrite are blurry, wavy and sparse. Many microcracks mostly initiate along the persistent slip bands in austenite; however, either grain boundaries or austeniteferrite interfaces are less favourable initiation sites for fatigue cracks. There is considerable evidence that the austenite region with a low stacking fault energy exhibits less initiation resistance fatigue cracking than ferrite. Further study of the propagation of surface cracks was conducted on both FL and FT specimens. The lightly etched specimens were taken from the bending machine for crack length measurement and photographic recording after every 5000 cycles during the fatigue testing. The surface crack propagation rates of both FL and FT specimens are shown in Fig. 8, and they are comparable. Figures 9 and 10 indicate the propagation process of surface cracks on FL and FT specimens respectively. The fatigue cracks in either specimen propagated along a direction roughly perpendicular to the axis of applied stress. Figure 9 shows that the string arrangement of the austenite phase in specimen FL is perpendicular to the crack propagation, and a microcrack in the austenite region has penetrated through a neighbouring ferrite phase and linked with another microcrack. Meanwhile, no initiation of microcrack is observed within the ferrite region. When the string arrangement is parallel to the crack propagation as in specimen FT, the cracks have initiated and propagated mostly 1,5 1./+ 1.3 1.2 1.1 1.0
~ ~= g -~
along the austenite phase as shown in Fig. 10. The crack penetrated into the ferrite region on some occasions, then the propagation was impeded in the ferrite phase, and the cracks turned back into austenite again. The crack blunting phenomenon was" also observed in the ferrite region at the lower right area of Fig. 10. It is believed that the ferrite phase serves as a crack arrester in this high stress amplitude bending fatigue test which is quite contrary to most fatigue behaviours of carbon steels or alloy steels. Bathias and Pelloux [13] found that austenite steel with a low stacking fault energy displayed a superior resistance in fatigue crack propagation owing to dislocations constrained to move in a more planar fashion; thus it limited local concentration of plastic deformation and suppressed
Fig. 9. Optical micrograph of an FL specimen which has been fatigued (25 000 cycles with an applied bending stress of 442.6 MPa). A microcrack has penetrated through a neighbouring ferrite phase and linked with another microcrack in the austenite region.
o:FL specimen o:FT specimen
0.9 0.8 0.7 0.~ 0.5 0.3 O.2 03 O.C cycles ( xlO 4 )
Fig. 8. Total crack lengths of FL and FT specimens vs. fatigue cycles. The crack propagation rates of both FL and FT specimens are comparable.
Fig. 10. Optical micrograph of an FT specimen which has been fatigued (25 000 cycles with an applied bending stress of 442.6 MPa). A crack penetrated into the ferrite phase then turned back to the austenite region, meanwhile, the crack blunting phenomenon occurred within the ferrite region at the lower right area.
30
fatigue damage. Wan et al. [19, 25, 26] observed that Fe-AI-Mn steels with an austenite structure are better than those with a ferrite structure in many properties, such as strength, ductility and corrosion resistance. However, their steels contain certain amounts of elemental carbon ranging from 0.1% to 1.5%. The carbon mostly dissolves in the austenite phase, or precipitates to form carbides in the ferrite phase. This effect creates a significant solid solution strengthening in austenite, and raises both the strength and the ductility which greatly contributed to the resistance of fatigue crack initiation and propagation. However, the hard and brittle carbides in the ferrite region are detrimental to the resistance to fatigue cracking by both the stress concentration on the interfaces and the breakage of carbide particles. The present study proves that once the carbon effect is removed, the different crystal structure rather than the stacking fault energy provides the important fatigue resistance, and the ferrite phase proves to be a better crack arrester than austenite by the crack blunting effect. The better resistance character in ferrite also explains the comparable crack propagation rates between FL and FT specimens as shown in Fig. 8. Since the frequencies of crack grown into the ferrite area are comparable between those two specimens, similar impediments to crack propagation hold for each sample even though the string arrangements of constituents are completely different. Finally, there is no dislocation cell substructure observed in the ferrite region after tensile testing, and ferrite cannot play the role as a crack arrester, meanwhile, ferrite is softer than austenite; therefore, both the strength (YS and UTS) and the elongation of ST specimens are low.
4. Conclusions The extent of the string arrangement in microstructure and associated tensile and bending fatigue properties in an austenite-ferrite dualphase Fe-8.2wt.%AI-30.1wt.%Mn steel have been investigated. The results are summarized as follows. ( 1 ) The carbon plays an important role in both strengthening and fatigue resistance in the austenite phase. Once the carbon effect is removed, the ferrite phase provides better resistance to fatigue crack than austenite. (2) A well-defined dislocation cell substructure developed in the ferrite region by high ampli-
rude of cyclic stress which corresponded to the observed crack blunting phenomenon and made the character as fatigue crack arrester of the ferrite phase. (3) The low stacking fault energy in the austenite phase raised a planar deformation mode which generated less resistance to fatigue crack by dislocation pile-up and stress concentration. (4) No dislocation cell substructure was observed in the ferrite region after tensile testing; therefore, both the strength and the elongation ot ST specimens were lower than those of SL.
Acknowledgment This work was supported by the National Science Council, Taiwan.
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