Intermetallics 12 (2004) 893–897 www.elsevier.com/locate/intermet
Improvement of the creep resistance of FeAl-based alloys W.J. Zhang*, R.S. Sundar, S.C. Deevi Chrysalis Technologies Inc., Richmond, VA 23237, USA Available online 1 April 2004
Abstract By increasing the concentration of carbon or transition metals, the creep resistance of a Fe – 40Al – 0.2Mo– 0.05 Zr – 0.5Ti– 0.2C – 0.02B (at.%) alloy was significantly improved with the minimum creep rate decreasing by the order of three. The improvement was attributed to precipitation hardening and solid solution hardening. The influence of precipitation hardening and solid solution hardening on the creep behavior of FeAl-based alloys was analyzed. Formation of grain boundary microcracks and the microstructual instability are likely to be the concerns for FeAl-based alloy during long-term creep. q 2004 Elsevier Ltd. All rights reserved. Keywords: A. Intermetallics, miscellaneous; A. Iron aluminides (based on FeAl); B. Creep (properties and mechanisms); B. Corrosion
FeAl-based alloys are promising candidates for hightemperature applications due to their low density, excellent oxidation resistance, good sulfidation resistance, and highelectrical resistivity [1,2]. In addition, creep strength at high temperatures is an important property for many of the structural applications. The creep strength of binary FeAl is considerably low at high temperatures because of its noncompact B2 structure and high-vacancy concentration. Alloying additions such as Zr, Nb, Cr, Ti, C and B are effective in improving the creep resistance of FeAl alloys through precipitation hardening or solid solution hardening [3 – 5]. The creep strength of FeAl can also be enhanced via oxide dispersion strengthening (ODS) . Unfortunately, the creep resistance of FeAl-based alloys developed to date is still inferior to that of the 310 stainless steel. The objective of the present study is to examine the creep behavior of FeAl alloys influenced by precipitation hardening and solid solution hardening methods. In this article, we present and analyze the data of two FeAl alloys with enhanced creep strength, as well as that of a base FeAl alloy, and provide suggestions for further research.
The nominal compositions of the three studied FeAl alloys are given in Table 1. As compared to the base alloy 785, high carbon alloy 787 and high-transition metal alloy 912 are designed to improve the creep resistance through precipitation hardening and solid solution hardening by increasing the concentration of carbon or transition metal elements, respectively. Ingots of approximately 7 kg were arc-melted and then extruded with a reduction ratio of 21:1. Alloy 912 was extruded at 950 8C and alloys 787 and 785 were extruded at 1050 8C. Round creep samples with a gauge length of 25 mm and a gauge diameter of 5.7 mm were machined from the extruded rods with the tensile axis lying along the extrusion direction. The samples were heat treated at 1100 or 1300 8C for 2 hurs and air-cooled (AC) to room temperature. Creep tests were performed in air at 700 8C (^ 0.5 8C) in a constant load creep frame. The specimen elongation was measured with a Class B type extensometer with a strain resolution of 1025. A data acquisition unit continuously monitored the load, temperature and elongation of the specimen during testing.
3. Result and discussion
* Corresponding author. Current address: R&D T7, Philip Morris USA, 615 Maury Street, Richmond, VA 23234, USA. Tel.: þ 1-804-274-2320; fax: þ 1-804-274-2468. E-mail address: [email protected]
(W.J. Zhang). 0966-9795/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2004.02.020
Fig. 1 shows the optical microstructures of alloys 787 and 912 prior to creep testing. The alloys exhibit equiaxed grains with the grain size of 100 – 140 mm. Coarse second precipitates of 1– 5 mm in size can be seen in the alloys. These precipitates were determined to be carbide by EDS
W.J. Zhang et al. / Intermetallics 12 (2004) 893–897
Table 1 Nominal compositions (at.%) of the FeAl alloys Alloys
Base alloy (785) High carbon alloy (787) High TM alloy (912)
40 40 40
0.2 0.2 0.5
0.05 0.05 0.1
0.5 0.5 1.0
0.2 0.7 0.2
0.02 0.02 0.02
Bal. Bal. Bal.
analysis. The base alloy 785 has a similar microstructure with coarse precipitates distributed on the matrix. We believe that these coarse particles have little effect on strengthening. Fig. 2 provides the stress dependence on the minimum creep rates of the three alloys in the stress range of 50 –150 MPa at 700 8C. The data of an ODS-FeAl alloy (CEREM) are also presented for comparison . This alloy was a mechanically alloyed Fe – 40Al alloy with 1% volume fraction of fine Y2O3 particles (12 – 50 nm) dispersed in the matrix. As can be seen, the creep rates of alloys 787 and 912 are significantly lower than those of base alloy 785 and even lower than the ODS-FeAl alloy. At 90 MPa, the minimum creep rate of alloy 787 is three orders of magnitude lower than that of base alloy 785 which was heat treated at the same temperature.
Fig. 1. Optical microstructures of (a) alloy 787 heat treated at 1300 8C/2 h/AC and (b) alloy 912 heat treated at 1100 8C/2 h/AC.
Fig. 2. Stress dependence of minimum or steady-state (Min/SS) creep rates of the base alloy 785, high carbon alloy 787 and high transition metal alloy 912 tested at 700 8C. For comparison, creep results of an ODS-FeAl alloy (CEREM) are included .
The effect of heat treatment on the creep resistance of these alloys was reported elsewhere . Note that the 787 alloy exhibits a typical creep rate vs. stress curve of ODS alloys with the stress exponent n changing from 3 to 5 at high stresses to , 20 at low stresses . This behavior is related to the back-stress effect of fine oxide or precipitates on dislocation movement during creep. With high carbon concentration and high-temperature annealing at 1300 8C, a high density of fine precipitates in the size of 10 – 20 nm were observed in alloy 787 as shown in Fig. 3. These precipitates appear to be titanium carbides as determined by EDS analysis. The precipitates blocked the dislocation sources in the heat-treated sample prior to creep (Fig. 3a), and the density of dislocations in the crept samples is thus low. These dislocations were also pinned by in situ formed fine precipitates (Fig. 3b). Therefore, the low creep rate of alloy 787 is likely a result of low-dislocation density and low-dislocation mobility due to the pinning effect of fine precipitates. In addition, a deep valley was observed on the creep rate curves of alloy 787 at 120 –150 MPa before the onset of steady-state creep, suggesting that the dislocations had to overcome the pinning of precipitates at the initial stage of creep. Similarly, Maziasz et al.  reported that the presence of fine FeTiP and Ti-rich carbides are responsible for the good creep resistance of a complex Fe – 36Al – 0.4C alloy. Fig. 4 compares the creep rates and creep strains of base alloy 785 and high carbon alloy 787 as a function of time at 700 8C and 90 MPa. Alloy 787 exhibits a well-developed steady-state creep regime in addition to its low creep rate.
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Fig. 4. Creep rates and creep strains of (a) base alloy 785 and (b) high carbon alloy 787 as a function of time at 700 8C and 90 MPa after heat treatment at 1300 8C/2 h/AC.
Fig. 3. TEM micrographs showing fine precipitates in the size of 10– 20 nm (marked by arrows) pinned (a) the dislocation source in the 787 sample prior to creep and (b) the dislocations in the crept 787 sample.
In contrast, no steady-state creep was observed in the base alloy 785. The creep rate of this alloy increases by approximately 10 times after creep for 3 h. Dislocation bands, grain boundary bulging and recrystallization were observed in the crept 785 sample. The steep increase in creep rate of this alloy may likely have resulted from the localized deformation, which may have led to high dislocation density and finally recrystallization. In high carbon alloy 787, development of localized deformation bands and recrystallization was suppressed due to the pinning of dislocations by fine precipitates. Therefore, a long
steady-state creep regime was observed. Because of its low creep rate and well-developed steady-state creep, alloy 787 has a rupture life (1058 h) over 100 times longer than that of base alloy 785 (9 h) at 700 8C and 90 MPa. As can be seen in Fig. 2, the creep rates of high-transition metal alloy 912 are three orders of magnitude lower than those of base alloy 785 in the whole stress range studied. In the high stress range, the creep rates of alloy 912 are even lower than that of high carbon alloy 787 and ODS-FeAl alloys by two orders of magnitude. We believe that alloy 912 is strengthened by solid solution due to its high concentration of Ti, Mo and Zr. In contrast to precipitation hardening, solid solution hardening will lead to a lower creep rate in the whole stress range. The stress exponent of n ¼ 3 – 4 for 912 is also typical for solid solution alloys . Furthermore, we observed the initial decrease of creep strain at the early stage of creep in 912 tested at 700 8C and 90 MPa as shown in Fig. 5. A similar phenomenon has been reported in a TiAl-based alloy with W in solution . It is speculated that an impact stress larger than the nominal stress of 90 MPa was applied on the sample upon loading, leading to a high instantaneous primary strain. However,
W.J. Zhang et al. / Intermetallics 12 (2004) 893–897 Table 2 The minimum creep rates and rupture life of the FeAl-based alloys and 310 stainless steel tested at 700 8C and 90 MPa
Fig. 5. Creep curve of high-transition metal alloy 912 tested at 700 8C and 90 MPa. Note the initial decrease in creep strain at the early stage of creep.
the normal stress of 90 MPa may not be large enough to activate further dislocation multiplication due to highfriction stress as a result of solid solution, and thus anelastic recovery such as back motion of dislocations occurs resulting in the initial decrease of creep strain. Note that the creep strain kept nearly constant up to 40 h after the initial decrease, followed by a steady-state creep regime. Then the creep strain increases rapidly at around 80 h, and the sample finally ruptured at 190 h with a creep elongation of 42%. The reason for the rapid increase of creep rate after 80 h is not clear yet. No rapid increase of creep rate was observed in 912 tested at 600 8C and 207 MPa for 1200 h (see Fig. 6). It may suggest that the rapid increase of strain observed at 700 8C resulted from microstructural instability which is highly temperature dependent. Table 2 compares the minimum creep rates and rupture life of the three FeAl alloys and the 310 stainless steel. It can
Min. creep rate (1/s)
Rupture life (h)
785 787 912 310SS
2.5 £ 1027 2 £ 10210 1.0 £ 1029 ,5 £ 1028
9 1058 190 ,1150
be seen that the minimum creep rates of FeAl alloys were significantly reduced via precipitation hardening and solid solution hardening. The minimum creep rates of 787 and 912 are more than an order of magnitude lower than that of the high-grade stainless steel 310. Comparison with other commercial alloys was reported elsewhere . However, the improvement in rupture life is not as significant as in creep rate; while the rupture life of 787 alloy is comparable to that of 310 stainless steel but the rupture life of alloy 912 is inferior to that of 310 stainless steel. In consideration of its low-creep rate, the alloy 787 should have a rupture life longer than that of the 310 stainless steel. The shorter than expected rupture life of alloy 787 is due to the early fracture of the sample. As shown in Fig. 6, the sample fractured at a low strain of around 0.3% after creep for 1050 h at 700 8C and 90 MPa. A similar low-rupture strain of 0.34% was observed in the 912 sample tested at 600 8C and 207 MPa with the rupture life of 1000 h (Fig. 6). Examination of the crept sample showed microcracks along grain boundaries even in the head of crept samples. Therefore, formation of grain boundary microcracks appears to be responsible for the early rupture of the samples.
Fig. 6. Creep curves showing the 787 and 912 alloys ruptured at a strain of around 0.3% after creep for over 1000 h.
The creep resistance of a FeAl alloy Fe –40Al – 0.2Mo – 0.05Zr – 0.5Ti – 0.2C –0.02B was significantly improved by increasing the carbon content or by increasing the concentration of transition metals. The minimum creep rates decrease by three orders of magnitude at 700 8C and 90 MPa. The improvement was attributed to precipitation hardening and solid solution hardening, respectively. The high carbon alloy 787 has a creep rate vs. stress curve typical of oxide dispersion strengthened alloys with the stress exponent n in the range of 3– 5 at high stresses and , 20 at low stresses. Dislocations were found to be pinned by fine precipitates in the size of 10– 20 nm in the crept 787 sample. Precipitation hardening also promotes the development of a steady-state creep regime in addition to reducing the minimum creep rate. Solid solution hardening greatly decreases the minimum creep rate at both low and high stresses. However, the strengthening effect may be weakened after a certain period of creep due to microstructual instability. After long-term creep for over 1000 h, samples
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of both alloys fractured at a relatively low creep strain of around 0.3%, probably due to the formation of grain boundary cracks.
References  Stoloff NS, Liu CT, Deevi SC. Intermetallics 2001;8:1313.  Lilly AC, Deevi SC, Gibbs ZP. Mater Sci Engng A 1998;A258:42.  Munroe PR, Kong CH. Intermetallics 1996;4:403.
 Titran RH, Vedula K, Anderson GG, Mater Res Symp Proc, vol. 39. MRS; 1984.  Maziasz PJ, Goodwin GM, Alexander DJ, Viswanathan S. In: Deevi SC, et al., editors. Proceedings of International Symposium on Nickel and Iron Aluminides: Processing, Properties and Applications. Cleveland, OH: ASM International; 1996. p. 157.  Morris-Munoz MA. Intermetallics 1999;7:653.  Sundar RS, Deevi SC. Metall Trans A 2003; 34A:2233.  Evan RW, Wilshire B. Creep of Metals and Alloys. The Institute of Metals; 1985. p. 114.  Zhang WJ, Deevi SC. Intermetallics 2002;10:603.  Sundar RS, Deevi SC. Mater Sci Engng A 2003; 357A:124.