Tensile properties of a DO3-containing FeMnAlSiC alloy at elevated temperatures

Scripta METALLURGICA et MATERIALIA

Vol.

26, pp. 1417-1421, 1992 Printed in the U.S.A.

Pergamon Press Ltd. All rights reserved

TENSILE PROPERTIES OF A D03--CONTAINING FE--MN-AL-SI--C ALLOY AT ELEVATED TEMPERATURES

J.S.CHOUt and C.G.CHAO~ tInstitute of Mechanical Engineering Hnstitute of Material Science and Engineering National Chiac>-Tung University, Hsinchu, Taiwan, R.O.C.

(Received December 27, 1991) (Revised March 3, 1992) Introduction

A series of results concerning the effects of silicon addition on the mechanical properties, corrosion resistance and high-temperature oxidation resistance of F e - M n - A 1 - C alloys have recently been presented by many authors (1-9). They suggested that the addition of silicon effectively improved the strength, corrosion resistance and high-temperature oxidation resistance of the alloys. A few papers have reported mechanical properties of F e - M n - A 1 - S i - C alloys at elevated temperatures (1,10). They chiefly focused on a fully austenitic matrix: An F e - 3 0 M n - 8 A I - I . 5 S i - I C alloy was reported to possess a yield strength of about 530 MPa and an elongation of 5 pct at -650" C by Schmatz (1). Silicon in F e - M n - A l - S i - C alloys was suggested to possibly act as a solid solution strengthener. The addition of silicon enhancing the formation of ferrite has also been indicated by a limited number of papers (1,6). D03 (Fe3A1 type, fcc cubic: Fm3m) anti-phase domains were found to be formed by an ordering transition during quenching and exist up to 850"C in our previous study of an Fe-29.0Mn---8.0Al-l.SSi-0.gc alloy (11). Investigating the elevated-temperature mechanical properties of a duplex (7 + DOs) F e - M n - A 1 - S i - C alloy is the purpose of this study. Another fully austenitic F e - M n - A 1 - C alloy has also been prepared for comparison. Experimental Procedure A duplex Fe-26.6Mn-7.SAl-l.65Si-0.73C Mloy (alloy A) containing D03 and austenite and an austenitic Fe-28.1Mn-7.5AI-0.75C alloy (alloy B) were melted in an induction furnace. The ingots were homogenized at 1200" C for two hours under an argon atmosphere; they were then hot forged and hot rolled to a final thickness of 3.0 mm. The chemical compositions were analyzed by atomic emission spectrometry. After solution treatment at 1050"C and quenching, tensile specimens were machined to a gauge length of 25 mm with a cross section of 6.25 mm x 2.0 mm. Tensile tests were carried out in an Instron tensile testing machine at temperatures ranging from 25 ° C to 650" C, with an initial strain rate of lxl0 -4 s -~. Each specimen was kept at a test temperature for 10 minutes before the test. This was done to ensure a uniform test temperature of the specimen. Thin film specimens for transmission electron microscopy (TEM) were prepared by means of a double jet electropolisher with an electrolyte of 30 pct ethanol, 60 pct acetic acid and 10 pct perchloric acid. Polishing temperature was kept in a range from -5°C to 5"C, with current density in a range from 1.5x104 to 2.5x104 A/mS. TEM was performed on a JEOL-2000FX scanning transmission electron microscope operating at 160 kV. The fracture surfaces were examined by Hitachi S-2500 scanning electron microscope (SEM) operated at 20 kV. Results and Discussion Figure 1 (a) and (b) are, respectively, optical micrographs (OM) showing solution-treated microstructures of alloy A and alloy B. The microstructure of alloy A consists of equiaxed austenite grains

1417 0036-9748/92 $5.00 + .00 Copyright (c) 1992 Pergamon Press Ltd.

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(i.e., marked as "'),") with an average grain size of -17 p~n and D03 grains (i.e., marked as "Dr') with an average grain size of -14 #m. The volume fraction of D03 phase is about 27 pct. Metallographic observation for alloy B reveals a uniform distribution of equiaxed recrystallized austenite grains having an average size of "50 #m. No precipitate was found within austenite matrices in either of the two solution-treated alloys, which was confirmed by TEM. TEM observations of D03 in a solution-treated specimen of alloy A and corresponding electron diffraction pattern are shown in Figures 2(a) and (b). Figure 2(a) is a g = l l l D03 dark-field (DF) electron micrograph taken from a region within a D03 grain. An is.tropic D0~ antiphase-boundary (APB) microstructure is observed. Figure 2(b) is the [0Ill---zone axis diffraction pattern showing intense {111} D0~ spots, giving clear evidence of a high degree of D03 order. The above TEM observations illustrate that D03 ordering was not suppressed during quenching. The relationships between tensile properties of the two alloys and test temperature are shown in Figures 3(a) through (c). Yield strengths (0.2 pct offset strain) of alloy A in Figure 3(a) are always higher than those of alloy B at various test temperatures. Both alloy A and alloy B show positive temperature dependence of yield strengths from 200"C to 500" C. Peaks of yield strengths appear in the vicinity of 500' C. The small increment in yield strengths of alloy B from 200" C to 500"C can be attributed to the effect of K-phase ((Fe,Mn)3A1Cx, L'12 structure) precipitation hardening (10,12). K-phase precipitation may also have a similar hardening effect on the austenite matrix of alloy A. Figure 4 is a [011] zone-axis diffraction pattern which is taken from the austenite matrix in a specimen of alloy A after fracture at 500" C. The ~100} superlattice spots indicate the formation of the K-phase during tensile testing at an elevated temperature. The causes of higher yield strengths and distinct yield strength peak in the (7 + D03) alloy are probably more complicated. Many D03 alloys possess high yield strengths and exhibit an anomalous yielding behavior in the intermediate temperature region (13-18). This fact hints at the D03 phase having a dominant effect on the yielding behavior of the (7 + D03) alloy. Clarifying the role of D03 on the anomalous yielding behavior in alloy A is difficult since the mechanism of anomalous yielding in common D03 alloys is still not very clear. Elongations of alloy B in Figure 3(c) increase from 25"C to 200'C with increasing temperature. Elongation then gradually decreases with increasing temperature, which may be due to the K-phase precipitates formed. Only a dimple fracture surface is observed in fractured specimens of alloy B. Elongations of alloy A are limited below 20 pct at test temperatures below 550"C. A significant increase in elongations of alloy A above 550'C is, however, observed. OM observation (Figure 5) illustrates the formation of transgranular cracks within D03 grains (i.e. marked as "D") in a fractured tensile specimen of alloy A, which is tested at 25" C. Figure 6 is an SEM fractograph taken from the same specimen shown in Figure 5. This figure reveals transgranular cleavage-like surfaces (i.e., marked as "C") with cracks and portions of dimple fracture surfaces (i.e., marked as "D"). The proportion of cleavage-like surfaces decreases with increasing temperature. Cleavage-like fracture surfaces can be observed up to 550" C. OM and SEM observations give an explanation on the limited elongation of alloy A below 550" C. The existence of the D03 phase can obviously significantly reduce the elongations of Fe---Mn-A1-C alloys below 550" C. Summary The tensile properties of the duplex (7 + D03) Fe-26.6Mn-7.8Al-l.65Si-0.73C alloy (alloy A) and the austenitic Fe---28.1Mn-7.5A1-0.75C alloy (alloy B) have been investigated. 1. The (7 + D03) alloy exhibited higher yield strengths than the anstenitic alloy at various temperatures from 25"C to 650" C. The (7 + D03) alloy also exhibited a more distinct increase in yield strengths with increasing temperature in the intermediate temperature region from 2000C to 500' C. The austenitic alloy, however, exhibited only a small increment in yield strengths, with increasing temperature in the same temperature region. D03 apparently played an important role in the difference of yielding behavior between the two alloys. 2. The elongations of the (7 + D03) alloy were significantly lower than those of the austenitic alloy when tested at temperatures below 550" C. The existence of the D03 phase apparently reduced the elongations of the Fe--Mn-A1-Si---C alloy in the lower temperature region.

Acknowledgements The authors are pleased to acknowledge financial support of this research by the National Science Council, Republic of China, under Grant No. NSC 79--0405-E--009-11. They are also grateful to Miss P.F.Chou for typing.

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Reference 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18.

D.J. Schmatz, Trans. ASM, 52, 898 (1960) S.K. Bauerji, Met. Prog., Apr., 59 (1978) R. Wang and F.H. Beck, Met. Prog., Mar., 72 (1983) J.S. Dunning and M.L. Glenn, Met. Prog., Oct., 19 (1984) C.J. Altstetter, A.P. Bentley, J.W. Fourie and A.N. Kirkbride, Mater. Sci. Eng., 82, 13 (1986) J. Charles, A. Berghezan, A. Lutts and P.L. Dancoisne, Met. Prog., 119, 71 (1981) R. Wang, M.J. Straszheim and R.A. Rapp, Oxid. Met., 21, 71 (1984) J. Sauer, Oxid. Met., 18, 285 (1982) S.K. Banerji, The 1982 Status Report On Fe--Mn-A1 Steels, Foote Mineral Co., Extort, PA (1982) R.K. You, P.W. Kao and D. Gan, Mater. Sci. Eng., Al17, 141 (1989) T.F. Liu, J.S Chou and C.C. Wu, Metall. Trans. A, 21A, 1891 (1990) K.S. Chan, L.H. Chen, T.S. Lui and Y.W. Chao, Proc. Janpanese--Sino Symposium on Casting, Jan., 102 (1990) J.W. Park, I.G. Moon and J. Yu, J. Mater. Sci., 26, 3062 (1991) S.K. Ehlers and M.G. Mendiratta, J. Mater. Sci., 19, 2203 (1984) H. Inouye, Mat. Res. Soc. Symp. Proc., 39, 255 (1985) R.S. Diehm and D.E. Mikkola, Mat. Res. Soc. Symp. Proc., 81, 329 (1987) C.G. Mckamey, J.A. Horton and C.T. Liu, Mat. Res. Soc. Symp. Proc., 81,321 (1987) M.G. Mendiratta, S.K. Ehlers, D.M. Dimiduk, W.R. Kerr, S. Mazdiyasni and H.A. Lipsitt, Mat. Res. Soc. Symp. Proc., 81,393 (1987)

(a)

(b)

FIG.1. OM of solution-treated microstructures of (a) alloy A and (b) alloy B (austenite: marked as "7"; D03: marked as "D").

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FIG.2. TEM micrographs of D03 taken from a solution-treated specimen of alloy A: (a) g = l l l D03 DF, (b) [011] zone-axis diffraction pattern. 1000

1000

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800

N

600

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100 200 ,.ZOO 400 500 Temperature

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TemperGture

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FIG.& Temperature dependence of tensile properties: (a) yield strength (0.2 pct offset strain) (b) ultimate tensile strength (c) elongation.

0

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100 200 300 400 500 600 700 Temperoture

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alloy A alloy B

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FIG.4. [011] zone-axis diffraction pattern taken from a fractured tensile specimen of alloy A, which is tested at 500 ° C.

FIG.5. OM of a fractured tensile specimen of alloy A, which is tested at 25" C (austenite: marked as "7"; D0~: marked as "D").

FIG.6. SEM fractograph, taken from the same specimen shown in Figure 5 (transgranular cleavage-like fracture surfaces: marked as "7", dimple fracture surfaces: marked as "D").