Effects of alloying additions on the microstructures, mechanical properties and weldability of Fe3Al-based alloys

Effects of alloying additions on the microstructures, mechanical properties and weldability of Fe3Al-based alloys

Materials Science and Engineering, A 174 (1994) 59-70 59 Effects of alloying additions on the microstructures, mechanical properties and weldability...

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Materials Science and Engineering, A 174 (1994) 59-70


Effects of alloying additions on the microstructures, mechanical properties and weldability of Fe3Al-based alloys C. G. McKamey, P. J. Maziasz, G. M. G o o d w i n a n d T. Z a c h a r i a Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6114 (USA) (Received February 2, 1993; in revised form May 27, 1993)

Abstract Several alloys based on the Fe-28AI-5Cr (atomic per cent) composition were produced to study the effects of alloying with Mo, Nb, Zr, B and C on microstructures, mechanical properties, and weldability. Optical microstructures were examined before and after heat treatments of 1 h at 750 °C, as well as after selected higher temperature anneals. Tensile properties at room temperature and 600 °C and creep-rupture properties at 593 °C and 207 MPa were determined and correlated with alloying additions. Judgments as to weldability of selected alloy compositions were made by determining the hot-crack susceptibility. The results indicate that all of these properties of iron aluminides are very sensitive to alloying additions. Some combinations of the above elements resulted in refined grain sizes, increased recrystallization temperatures, and strengthening of the base alloy (through solid solution effects, as well as the formation of precipitates) in both tension and creep-rupture tests. Increased strength, however, was generally produced at the expense of room temperature ductility and weldability. The results suggest that the design of useful iron aluminide compositions will depend on the application, with the composition being modified to provide either room temperature ductility and weldability or strength, as prescribed by the intended application.

1. Introduction Iron aluminides based on Fe3Al have excellent oxidation and corrosion resistance [1]. However, until recently their potential use as a heat-resistant structural material has been limited by low room temperature ductility (less than 5% in binary Fe3AI) and a drop in strength above 600 °C [2]. Recent studies indicate that the poor ambient temperature ductility observed in iron aluminides is caused (at least in part) by dynamic hydrogen embrittlement resulting from the dissociation of water molecules in the environment by aluminum atoms at the surface of the alloy [3]. This environmental embrittlement can be minimized by using alloying and thermomechanical processing techniques to control alloy chemistry, microstructure, and surface condition [3]. Fe3Al-based alloys can exist in one of three different structures: disordered solid solution, an imperfect B2 ordered structure, or an ordered D03 structure. According to the currently accepted binary phase diagram [4], the temperature of the disordered-B2 transition is approximately 900 °C, while the B2-D03 transition occurs at approximately 550 °C. It is also possible to produce a mixed non-equilibrium B2-D03 structure in room temperature specimens by fast quenching from above 550 °C [5]. 0921-5093/94/$7.00

The results of past studies of the effect of alloying on the properties of iron aluminides have been reviewed previously [2, 6-8]. High temperature tensile and creep strengths have been shown to be improved with additions of elements such as Ti, Mo, Zr, Hf, or Nb, but at the expense of room temperature ductility [2, 9, 10]. With B or C these elements also form precipitates which are potent strengtheners and grain size refiners [2, 11]. With the addition of several elements together, synergistic effects on metallurgical properties become very important, while at the same time becoming more difficult to determine. Fe3Al-based alloys have been developed which can achieve ambient temperature tensile ductilities of 10%-20% and tensile yield strengths as high as 500 MPa [12]. For example, alloy FA-129 (Fe-28A15Cr-0.5Nb-0.2C, atomic per cent), which was developed at Oak Ridge National Laboratory and is currently being scaled up to commercial-size castings for extensive study, has good tensile strength (400 MPa or more) at temperatures to approximately 650 °C, with a room temperature ductility of 15%-20% [12]. Its weldability is better than that of nickel aluminides and has been greatly improved by optimizing the welding process and parameters (e.g. using pre- and post-weld heat treatments to prevent cold cracking)[13]. However, for many applications its low creep-rupture strength © 1994 - Elsevier Sequoia. All rights reserved


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Alloying additions in Fe~Al-based alloys

(20-30 h life at 593 °C and 207 MPa) is not adequate. Additions of Mo or Nb in combination with a small amount of Zr result in a dramatic improvement in creep-rupture strength of binary Fe3A1 ( 100-300 h life at 593 °C and 207 MPa), but reduce the room temperature tensile ductility and weldability [10, 14]. Further efforts are therefore necessary to produce an Fe3Al-based composition with an acceptable combination of these properties. In this study we present the results of alloy development efforts to improve the mechanical properties and weldability of alloy FA-129. In particular, the effects of additions of Nb, Mo, Zr, B, and C on the microstructure, tensile and creep-rupture properties, and weldability of Fe3A1 alloys based on the FA-129 composition will be presented.

2. Experimental procedures Table 1 lists the alloy compositions used in this study. Except for one alloy (FA-97 with only 2% Cr), the base composition is Fe-28A1-5Cr (unless otherwise stated, all compositions are given in atomic per cent). Alloys FA-170 and -170R are different sized ingots of the same nominal composition. All alloy compositions were prepared as 500 g ingots by arc melting and drop casting into chilled copper molds. Alloy FA170R was also produced as a 7 kg ingot to provide enough material for additional tests. Fabrication to 0.8

TABLE 1. Nominalcompositions(atomic per cent) of iron aluminides under study Alloy



Nb Mo


FA-129 28 FA-97 28

5.0 0.5 2.0 0.5 2.0


FA-130 FA-131 FA-133 FA-134

28 28 28 28

5.0 5.0 5.0 5.0

0.5 0.5 0.5 0.5

0.5 0.5 0.5 0.5

0.1 0 0.1 0

FA-162 FA-163 FA-164 FA-165 FA-166 FA-167 FA-168 FA-169 FA-170 FA-174 FA-175 FA-176 FA-177 FA-178

28 28 28 28 28 28 28 28 28 28 28 28 28 28

5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0 5.0

0.5 0.5 0.5 0.5 0.5 0.5 0.5 0.5 0.5 0.2 0 0.5 0.2 0

0.5 0.5 0.5 0.5 1.0 0.25 0.8 0.25 0.4 0.4 0.4 0.4 0.4 0.4

0.05 0.05 0.05 0.05 0.02 0.025 0.05 0.025 0.025 0.025 0.025 0.025 0.025 0.025





Balance Balance

0 0 0 0

0.05 0.05 0.2 0

Balance Balance Balance Balance

0 0 0.1 0.2 0 0.1 0.03 0.05 0.05 0.05 0.05 0.05 0.05 0.05

0 0.01 0 0 0 0.005 0.005 0 0 0 0 0.005 0.005 0.005

Balance Balance Balance Balance Balance Balance Balance Balance Balance Balance Balance Balance Balance Balance


mm thick sheet was accomplished by hot rolling, beginning at 1000 °C and finishing at 600-650 °C, with the final sheet material containing approximately 70% warm work. Selected alloys were also hot rolled all the way to 0.8 mm at 850 °C. Flat tensile specimens (0.8 m m x 3 . 1 8 m m x l 2 . 7 mm) were mechanically punched from the rolled sheet, with some compositions being given a stress relief heat treatment of 1 h at 700 °C to prevent cracking during the punching operation. The same specimen size was used for both tensile and creep-rupture tests. For comparison between alloy compositions, all alloys were further annealed for 1 h at 750 °C before testing in tension and creep-rupture. This heat treatment was expected to produce a partially recrystallized microstructure [15] with a mixture of B2 and D03 order [16]. Other heat treatments were used to change the microstructure and to study its effects on properties of selected compositions. Tensile tests were performed at room temperature (RT) and at 600 °C in air at a strain rate of 3.3x 10 -3 s -1 Creep-rupture tests were performed in air at 593 °C and 207 MPa. Tensile and creep ductilities were determined from measurements of the lengths of the specimens before and after testing. Optical metallography and scanning electron microscopy (SEM) were used to study the microstructures and fracture modes respectively. Analytical electron microscopy was performed on samples cut and electropolished from the gauge portion of selected creep-rupture tested specimens using Philips CM30 (300 kV) and JEOL 2000FX (200 kV) electron microscopes. The sensitivity to fusion zone cracking during welding was established using the Sigmajig weldability test [17]. This test uses a pre-applied transverse stress during the autogenous gas tungsten arc welding of a 50 mm by 50 mm sheet specimen. The applied stress is sequentially increased until the specimen cracks. This test ranks materials based on the threshold cracking stress (TCS), above which cracking occurs in the fusion zone. Sigmajig tests were performed using argon with 70 A d.c., at 12.5 mm s -1 travel speed and 1.14 mm arc length. The welds were characterized using optical microscopy.

3. Results and discussion Table 1 shows the nominal compositions of alloy FA-129 and several alloys produced to study the effect of alloying additions on the properties of FA-129. Because initial tests of the FA-97 alloy indicated significant strengthening and grain refinement relative to FA129 [18] this alloy was also included for comparison purposes. Tests of alloys FA-130 through -134 show the effects of adding B in combination with Zr; alloys

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Alloying additions in Fe~Al-based alloys

FA-162 through -165 show the effects of separately adding B or C; alloys FA-166 through -170 indicate the effects of keeping the Nb level constant and varying Mo, Zr, B and/or C; alloys FA-174 through -178 were used to investigate the effect of keeping Mo, Zr, and C constant while varying the amount of Nb in the presence or absence of B. T h e final rolling t e m p e r a t u r e for all alloys except FA-170R through -178 was 600 °C; that group of alloys was rolled at 650 °C. All alloys except FA-129 and -97 were stress relieved for 1 h at 700 °C before test specimens were punched. Only alloys FA-170 and FA-170R have been quantitatively analyzed for actual chemical composition. T h e results on these two alloys indicate that the measured compositions are very close to the nominal compositions, with very low levels of S, N, and O. Only the C levels were determined to be higher than expected: the nominal C level was 0.05%, while the measured level was approximately 0.08%. 3.1. Optical microstructural characterization Table 2 lists the macroscopic microstructural characteristics (grain size and morphology) of the alloys used in this study in the as-punched condition


(sheet was stress relieved before specimens were punched), as well as in two heat-treated conditions. Although grain sizes vary, the microstructures shown in Fig. 1 for FA-170R are representative of all the alloys. Most of the alloys contained precipitates, some of which were in the range of 2 - 3 ,um in size [11, 19]. T h e FA-129 and FA-97 alloys were each fabricated separately, while the other alloys were fabricated to sheet in groups of FA-130 through -134, FA-162 t h r o u g h - 1 6 6 , FA-167 and -168, FA- 169 and -170, and FA-170R through -178. T h e steps in the fabrication schedule were essentially the same for each group, but slight variations in actual processing details could cause slight differences in microstructure f r o m group to group. In the as-punched condition, all the compositions through FA-170 exhibited a predominantly striated, unrecrystallized microstructure characteristic of material flow during final hot rolling, although the thickness of the bands varied (see Table 2). Alloys FA131 and -162 also displayed a small fraction of tiny recrystallized grains near the outer surfaces which were assumed to have formed during the stress-relief heat treatment of the sheet at 700 °C. T h e thicker bands can be attributed to a lack of precipitates for grain refine-

TABLE 2. Optical microstructures of F%Al-based alloys a Alloy~'

As punched Recrystallization (%)

1 h-750 °C Grain size (pm)

Recrystallization (%)

Grain size (/~m)

100 0

20-40 < 30

FA-129 d FA-97 d

(t 0

FA- 130 FA- 131 FA- 133 FA- 134

0 20 0 0

< 40 < 20 < 30 3(1-100

0 95 25 5(1

20-50 20-50 < 20 20-100

FA-162 FA-163 FA- 164 FA- 165

5 0 0 0

< 20 20-50 < 50 20-100

10 10 60 70

1(I-20 10-20 20-40 20-40

FA- 166 FA- 167 FA- 168 FA- 169 FA- 170

0 0 0 0 0

10-40 < 100 < 100 100-200 150-2(10

30 40 0 10 100

10-30 10-50 50-100 2(/-40 110

FA- 170R FA- 174 FA- 175 FA- 176 FA- 177 FA-178

0 50 50 0 25 50

< 40 < 20

1 h-850 °CO: grain size (/~m)

< 50 < 50 < 30 < 50

60 100 100 70 100 80

< 40 50 20-40 < 30 20-40 20-40

~'Grain size for 0% recrystallization represents the width of the as-rolled bands. bExcept where noted, alloy sheets were stress relieved for 1 h at 700 °C before punching tensile specimens. cPlus 3-5 d at 500 °Cto form D03 order. dNo stress relief given.

20-40 20 35 d 35 d 25 d 70~


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Alloying additions in Fe c41-based alloys

ment, as observed for FA-134 in comparison with FA170R in Table 2. Especially thick bands were observed in as-punched FA-169 and -170, although the reason could not be attributed totally to composition and may be partly due to unknown fabrication variables. Alloys FA-170R through -178 were fabricated together as a group, with a final 70% reduction at 650 °C (compared with a final rolling temperature of 600 °C for all the other alloys in this study). Of this set, only FA-170R (Fig. l(a)) and -176, which contained the higher Nb levels, have a fully as-rolled microstructure, suggesting a higher recrystallization temperature for these two alloys as compared with the other alloys in this group. The other four alloys, which contain either 0.2% or 0% Nb, exhibit microstructures which consist of 25%-50% recrystallized grains with diameters of 50/tm or less. After a heat treatment of 1 h at 750 °C, most of the alloys had a partially recrystallized microstructure (see Fig. l(b)). Alloys FA-97, -130, and -168 still exhibited a completely banded (rolled) microstructure indicating that the recrystallization temperature is higher for these compositions. Alloys FA- 129, - 170, - 174, - 175, and -177 were fully recrystallized by the 750 °C anneal. In general, alloys containing larger amounts of Mo and Zr, accompanied by B or C, exhibited less recrystallization and more grain refinement. Compositions FA-129, -97, and FA-130 through -134 were studied in the recrystallized and D0 3ordered condition (1 h at 850 °C plus 3-5 d at 500 °C), as well as in the stress-relieved but only partially recrystallized condition ( 1 h at 750 °C, which may have left some surviving B2-ordered phase). A heat treatment of 1 h at 850 °C generally was enough to recrystallize fully this series of alloys, resulting in average grain sizes of 20-40 ~tm for all compositions except FA-134. The FA-134 alloy, which had an average grain size of 70 /~m, contained no Zr, B, or C to produce precipitates and to promote grain refinement. Several general conclusions can be drawn from the optical study of the microstructures. The formation of precipitates resulted in grain refinement and increased recrystallization temperatures (see microstructures of FA-133 vs. FA-134, Table 2). The precipitates in the alloy compositions of this study were produced by the addition of Zr, Nb, and/or Mo, in combination with B and/or C. In general, the use of composition modification to control grain sizes involves an understanding of the solubility of elements in Fe3AI and their effects when in solution. Elements which have low solubility in Fe3AI and react with impurities such as C to form precipitates (e.g. Zr [8, 18, 20]) will be the most effective grain refiners. Mo and Nb, with solubilities in binary Fe3A1 of approximately 5-10 at.% [21] and less than 0.5% [9] respectively, can be effective grain refiners in the presence of B and/or C. Warm rolling at 600 °C

Fig. 1. Optical microstructures of alloy FA-170R (a) hot rolled at 650 °C and annealed for (b) 1 h at 750 °C and (c) 1 h at 1000 °C.

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Alloying additions in Fe/41-based alloys

resulted in highly deformed, unrecovered microstructures for most of the alloys. However, w a r m rolling at 650 °C of the FA-170R through -178 series of alloys resulted in partial recrystallization in all but alloys FA-170R and -176, which were the most complex alloy compositions. It is also quite evident f r o m the microstructures of alloys FA-170 and - 1 7 0 R (which were the same composition) that the fabrication schedule needs to be fully controlled to p r o d u c e the desired grain structures. 3.2. M e c h a n i c a l properties Table 3 shows the tensile properties at r o o m temperature and 600 °C. All the alloys were tensile tested after a heat treatment of 1 h at 750 °C (a heat treatment which for most compositions p r o d u c e d partial recrystallization), except FA-97 which was tested in the recrystallized and D03-ordered condition. Earlier studies have indicated that partially recrystallized or


unrecrystallized microstructures, which help minimize the environmental effect p r o d u c e d by the presence of water v a p o r in the test environment, result in higher R T yield strengths and also better R T ductilities c o m p a r e d with a totally recrystallized microstructure [15]. O n the basis of this observation, and on the results of tensile tests of several Fe3Al-based compositions in b o t h the partially and totally recrystallized conditions [12, 18], it is estimated that the R T yield strength of FA-97 after a 750 °C anneal would be over 600 MPa, while the ductility would be 3 % - 5 % . With the heat treatment used in this study, the FA129 alloy had average yield strengths of 384 M P a and 419 M P a at R T and 600 °C respectively, while R T tensile ductility was about 17%. This alloy has been p r o d u c e d commercially in larger ingots, and with controlled heat treatments has p r o d u c e d yield strengths of 4 0 0 - 5 0 0 M P a at temperatures up to 600 °C, with R T ductilities of 1 0 % - 2 0 % [12].

TABLE 3. Tensile properties of Fe3Al-based compositions Alloy a

Nominal composition (at.%)

600 °C b

RT b Strength c (MPa)

Elongation (%)



Strength (MPa)

Elongation (%)



FA-129 FA-97 d

Fe-28A1-5Cr-0.5Nb-0.2C Fe-28AI-2Cr-0.5Nb-2Mo-0.1Zr-0.2B

384 454

930 520

16.9 1.2

417 436

520 529

38.3 25.9

FA-166 FA-167 FA-168 FA-169 FA-170

Fe-28AI-5Cr-0.5Nb -1Mo-0.02Zr -0.25Mo-0.025Zr-0.1C-0.005B -0.8Mo-0.05Zr-0.03C-0.005B -0.25Mo-0.025Zr-0.05C -0.4Mo-0.025Zr-0.05C

505 509 686 502 474

888 822 773 803 680

10.7 8.6 4.9 8.7 5.5

525 506 586

609 565 653

27.9 21.1 24.8

FA-134 FA-131 FA-130 FA-133

Fe-28AI-5Cr-0.5Nb-0.5Mo -0.05B -0.1Zr-0.05B -0.1Zr-0.2B

404 419 681 639

568 559 783 890

4.0 3.2 3.1 6.3

448 433 561 523

512 514 596 552

40.3 43.9 32.9 25.2

FA-162 FA-163 FA-164 FA-165

Fe-28AI-5Cr-0.5Mo-0.5Nb-0.05Zr -0.01B -0.1C -0.2C

523 572 441 416

716 758 614 782

4.8 4.8 3.3 8.7




FA-175 FA-174 FA-170R e FA-178 FA-177 FA-176

Fe-28AI-5Cr-0.4Mo-0.025Zr-0.05C -0.2Nb -0.5Nb -0.005B -0.2Nb-0.005B -0.5Nb-0.005B

354 336 421 357 364 538

672 578 722 686 632 804

7.6 4.5 6.7 7.8 6.4 7.6

383 390 455 385 403 494

422 430 486 417 437 504

49.1 45.7 50.8 50.2 46.5 39.6

~Except where noted, all alloys were heat treated for 1 h at 700 °C before punching specimens, then specimens were annealed for I h at 750 °C. bTested in air at a strain rate of 3.3 × 10 -3 s ~. cYS, 0.2% yield strength; FS, fracture strength. dHeat treated for 1 h at 850 °C plus 7 d at 500 °C. eSame composition as FA- 170, but melted in 7 kg ingot.


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Alloying additions in Fe ~Al-based alloys

Because previous studies have indicated that a combination of Nb, Mo, and Zr additions produces the best high temperature creep and tensile strengths [10], the alloy compositions in this study were selected to investigate the effect of holding the concentration of one or two of these three elements constant while varying the other one or two, in combination with the presence or absence of B or C. The highest tensile strengths (at both RT and 600 °C) were exhibited by alloys FA-168, -130 and -133, all of which contain the largest amounts of Mo, Nb, and Zr in combination with B and/or C. Alloys without B or C, even though they contained signifiant levels of Mo, Nb, and Zr (e.g. FA-162), in general had less strength. In other words, without B or C present, more Mo, Nb, and Zr went into solution (to produce some solid solution strengthening [8]). B addition appears to be especially important for strength. Recently, Liu and George have found that B strengthens grain boundaries in iron aluminides just as it does in nickel aluminides [22]. Since it segregates to grain boundaries, it could also segregate to dislocations and affect their behavior. The addition of B with Zr to increase room temperature strength was especially noticeable in alloys FA-130 and -133 compared with FA-134, and in FA-163 compared with FA-162. Increased strength, accompanied by a transition from an intergranular to a transgranular fracture mode, has also been noted by others for the addition of B to Fe3Al-based alloys [20, 23-26]. It appears, therefore, that B produces strengthening through at least two different mechanisms: (1) solid solution or precipitation strengthening of the matrix and (2) grain boundary strengthening as shown by Liu and George [22]. The role of C is less clear. Tests of alloys FA-163 and -164 indicated reduced strengths with the addition of C in comparison with FA-162, while tests of FA-164 and FA-129 suggested no adverse effect on RT ductility. This is contrary to the work of Kerr [27] which suggested that C had an adverse effect on ductility and fracture mode. However, such contrasting results also suggest that the synergistic interactions between minor elements such as B and C added to these complex alloy systems are still relatively unknown, and their effects on properties are important. The results presented in Table 3 generally indicate that Mo, Nb, and Zr added to this alloy system produced increased tensile strengths at both RT and 600 °C. However, the increased strength came at the expense of ductility. Embrittlement was especially significant at RT where two factors interacted to produce low ductilities. One, the large number of coarse precipitates which contributed to strengthening also provided crack-initiation sites. Two, since these alloys deform at higher stresses, more stress likely exacer-

bates the environmental embrittlement effects of hydrogen. The RT fracture modes were determined using SEM. All the fractures occurred predominantly by transgranular cleavage, as shown in Fig. 2. The scale of

Fig. 2. Scanning electron micrographs of RT tensile fracture surfaces of Fe3Al-based alloys: (a) FA-133 hot rolled at 600 °C and annealed for 1 h at 750 °C; (b) FA-133 hot rolled at 600 °C and annealed for 1 h at 850 °C plus 5 d at 500 °C; (c) FA-169 hot rolled at 850 °C and annealed for 1 h at 750 °C.

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Alloying additions in Fe ¢Al-based alloys

the cleavage was determined by the general microstructure; finer grains or narrower as-rolled bands resulted in a finer-scale cleavage ( c o m p a r e Fig. 2(a) with Fig. 2(c)). A small fraction of intergranular failure was noted in some specimens containing a finer recrystallized grain structure (e.g. Fig. 2(b)), but its presence could not be correlated simply with alloying additions. T h e c r e e p - r u p t u r e properties at 593 °C and 207 MPa are shown in Table 4. For specimens heat treated for 1 h at 750 °C, the longest creep lives and lowest M C R s were observed for alloys FA-168 and -130, followed by -133, -97, and -163. T h e worst compositions for creep were FA-175 and -178, followed by FA129 and -177. T h e s e are the same alloy compositions which produced the best and worst tensile strengths respectively, indicating again the importance of specific strengthening mechanisms for the p r o p e r t y of interest in these Fe3Al-based alloys. Especially obvious from the data in Table 4 is the effect of C on c r e e p - r u p t u r e life. C additions of 0.1% and higher were extremely detrimental to creep life as shown by the data for the


FA-162 through -165 group of alloys. This effect has also been observed for other Fe3Al-based alloys with similar compositions [18]. T h e reduction of N b below the 0.5% level also appears to lower the creep resistance as indicated by the creep lives of the FA-170R through - 178 group of alloys. Although the best tensile properties were generally produced by a partially recrystallized microstructure, the grain size in fully recrystallized material played an important role in producing c r e e p - r u p t u r e resistance. Table 5 shows the effect of heat treatment on the c r e e p - r u p t u r e lives of several Fe3A1 compositions. Significant i m p r o v e m e n t s in life were noted for alloys FA-97 and -176 after heat treatments to recrystallize totally the microstructure. O t h e r compositions (e.g. FA-129 and -170R) exhibited only slight improvements on recrystallization. H o t rolling to sheet at a higher temperature also p r o d u c e d uniformly larger grain sizes and dramatic i m p r o v e m e n t s in creep resistance for alloys FA-97 and -170, but did not appear to improve FA-169. T h e r e certainly a p p e a r to be several complicated variables involved in the production of

TABLE 4. Creep-rupture properties of Fe3Al-based iron aluminides at 593 °C and 207 MPa Alloy ~

Microstructure Recrystallization (%)

FA-129 FA-97

100 0

FA-166 FA-167 FA- 168 FA- 169 FA- 170

30 40 0 10 100

FA- 134 d FA- 131J FA-130 d FA1-133 d

100 100 100 100

FA-162 FA-163 FA- 164 FA- 165 FA-175 FA- 174 FA-170R" FA- 178 FA- 177 FA-176

10 10 60 70 100 100 80 80 100 70

Life (h)

Elongation (%)

MCR b(% h i)

16.8 115.5

68.9 67.3

1.5 0.4

68.6 84.4 286.7 30.8 41.1

60.1 47.4 65.0 52.1 50.6

0.6 0.3 0.08 0.9 0.6

49.8 38.4 201.6 102.7

58.5 59.0 60.9 70.7

0.7 0.8 0.1 0.2

78.1 116.2 27.1 26.3

62.6 55.0 71.0 62.2

0.5 0.2 1.2 1.3

4.0 35.5 33.4 4.8 13.9 50.4

68.6 76.2 68.4 68.8 68.2 61.9

7.3 0.8 0.8 6.2 2.0 0.5

Grain size (/~m) 20-40 <30 c < 30 < 50 < 100 c < 40 110 60-80 25-45 25-45 20-30 <20 < 20 < 40 < 40 20-40 40-60 < 40 < 40 20-40 < 30

~'Unless otherwise noted, all alloy sheets were stress relieved for 1 h at 700 C before punching tensile specimens: then specimens were annealed for 1 h at 750 °C. bMinimum creep rate. ~Number represents the width of the as-rolled bands. ~No stress relief of sheet. Specimens were heat treated for 1 h at 850 °C plus 3 d at 500 °C. ~Cast in 7 kg ingot and hot rolled at 650 °C.


C. G. McKamey et al.


Alloying additions in Fec4l-based alloys

TABLE 5. Effect of heat treatment on creep-rupture properties of Fe3Al-based alloys Alloy

Processing, heat treatment"

Microstructure Recrystallization

Grain size







MCR (%h -1 )


lh-750°C 1 h-850 °C+ 5 d-500 °C HR850,1h-750°C

0 100 100

<30 b 10-20 150

115.5 463.6 790.4

67.3 47.2 31.0

0.4 0.04 0.02


lh-750°C lh-850°C + 5 d-500°C

100 100

20-40 20-40

16.8 21.8

68.9 75.4

1.5 1.0


1 h-750 °C HR850,1 h-750 °C

10 100

< 40 200

30.8 38.2

52.1 35.4

0.9 0.6


lh-750°C HR850,1 h-750 °C HR650,1h-750°C HR650,1h-1000°C HR650,1h-750°C HR650,1 h-750 °C + 70min-950 °C

100 100 50 100 70 100

110 260 < 30 90-100 < 30 30-40

41.1 276.0 33.3 61.3 50.8 107.6

50.6 36.2 78.9 56.9 56.6 49.1

0.6 0.04 0.8 0.35 0.4 0.1

FA-170R FA-176

aAll alloys were warm rolled at 600 °C except where noted. HR650 and HR850 designate hot rolling at 650°C and 850°C respectively. All alloys were given 1 h at 700 °C before punching tensile specimens. bNumber represents the width of the as-rolled bands.

improved creep resistance in these alloys, including grain size, compositions of matrix and precipitates, ordered structure (B2 vs. D03) , and possibly subtle aspects of the dislocation-subgrain structure. Each of these factors is affected by both the specific details of fabrication conditions and the post-fabrication heat treatments. More systematic studies are needed to sort out the individual roles of each of these factors. 3.3. Analytical electron microscopy microstructural characterization Specimens of several alloys that were creep-rupture tested at 593 °C and 207 MPa were selected for further microstructural characterization using transmission electron microscopy (TEM) to observe the dislocation structure, and X-ray energy-dispersive spectroscopy (XEDS) for microcompositional identification of precipitates. T E M specimens were cut from the gauge portions of tested specimens. The FA-129 and FA-97 alloys were heat treated for 1 h at 850 °C (recrystaUization) followed by 5 d at 500 °C (D03 ordering) prior to creep testing, and had rupture lives of about 22 h and 464 h respectively. The FA-170 and -170R alloys were similarly creep tested after different processing-heat treatment conditions, including (a) hot rolling at 850 °C followed by 1 h at 750 °C (FA-170, 281 h), (b) hot rolling at 600 °C followed by 1 h at 750 °C (FA-170, 43 h), and (c) hot rolling at 650 °C followed by 1 h at 750 °C (FA-170R, 31 h). Since the B 2 ~ D 0 3 transition temperatures for Fe3Al-based alloys are near or below

550 °C [2, 4], it was assumed that creep deformation at 593 °C took place in the ordered B2 phase regime. The FA-129 alloy specimen had very coarse grains and very few dislocations (Fig. 3(a)). B2-type antiphase boundaries (APBs) were associated with the 2-fold superdislocations and defined the coarse domain structure; within these domains was mottled strain contrast that suggested very fine (less than 10 nm) D03 phase domains. Although this particular specimen exhibited no coarse or fine precipitates, coarse ( 0 . 3 - 2 / t m ) nonuniformly distributed clusters of NbC (80Nb-20Fe, atomic per cent) have been identified via XEDS in other specimens of this same material [10, 14, 16]. By contrast, the FA-97 alloy specimen had a finer grain size and a much higher concentration of 2-fold superdislocation networks within the matrix. A m o n g the network dislocations was also a very high concentration of fine defects (1-38 nm in size) tentatively identified via their contrast behavior as dislocation loops rather than precipitates (Fig. 3(b)). This specimen also contained some very coarse (above 1 pm) clusters of NbC (Nb rich with some Zr) particles, with a somewhat more uniform distribution of smaller precipitate particles (not yet identified) in the matrix. This specimen also exhibited no evidence of ordered B2 or D03 APBs, although electron diffraction was consistent with a B2 phase matrix. The several FA-170 and -170R alloy specimens examined exhibited microstructural behavior that was somewhat intermediate between those of the two alloys

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Alloying additions in Fec4l-based alloys

described above. The effects of variations in processing-heat treatment conditions manifest themselves mainly in differences in the matrix dislocation content and in subtle aspects of the creep-induced subgrain boundary structure. The FA-170 specimen hot rolled


at 850 °C had coarser subgrain structure (1-13 pm in size) with a fairly low matrix dislocation network content; many of the subgrain boundaries were lowangle grain boundaries while others were planar, dense honeycombed arrays of dislocations. The FA-170

Fig. 3. Transmissionelectron micrographs showing(a) grain and dislocationstructure in FA-129, (b) dislocationloops in FA-97, and (c) formationof a subgrain boundary in FA-170 (hot rolled at 650 °C),all creep-ruptured to failureat 593 °C and 207 MPa.


C. G. McKamey et al. / Alloying additions in Feg4l-based alloys

specimen hot rolled at 650 °C had a somewhat coarser subgrain structure and a higher concentration of matrix network dislocations (roughly 2-3 times more); the subgrains in this specimen were defined almost exclusively by dense dislocation arrays, which can be seen forming from matrix dislocations in Fig. 3(c). Finally, the FA-170R specimen hot rolled at 600 °C had a finer subgrain structure ( 1-4/~m) with a matrix network dislocation structure similar to that of the specimen hot rolled at 850°C. Some coarse clusters of NbC ( 6 5 N b - 2 7 Z r - 6 F e - 2 M o , atomic per cent) precipitates were found in the specimen hot rolled at 850 °C, but no fine precipitates or fine dislocation loops were detected in any of the FA-170 or -170R alloy specimens. All the FA-170 and -170R specimens also had a predominantly D03 phase matrix. These microstructures correlate reasonably well with the creep-rupture properties and give some insight into underlying mechanisms. The FA-129 alloy fails rapidly with a high secondary creep rate (Table 5), and the microstructure, with very few dislocation networks and ordered B2 APBs, suggests a high degree of recovery during creep. Similarly, the FA-97 alloy exhibits a much longer rupture life and a much lower creep rate, and the microstructure, with a high concentration of network dislocations and loops, is consistent with significantly more deformation hardening and less recovery. The fact that the hardening is due to the unusual occurrence of dislocation loops rather than fine precipitates in this particular alloy is somewhat surprising, and will require further study. The properties behavior of the FA-170 and -170R alloy series of specimens falls between these two cases. Although the two specimens hot rolled at 600 and 650 °C have higher matrix dislocation contents than the specimen hot rolled at 850 °C, they have shorter rupture lives and higher creep rates. This suggests that the microstructural evidence for rearrangement of the matrix dislocation networks into planar honeycomb arrays reveals instability of the dislocation structure, the mechanism by which dislocation recovery occurs. The specimen hot rolled at 850 °C most likely had fewer dislocations to begin with, as well as a processinginduced subgrain boundary structure that was different from that caused by creep deformation at lower temperatures. Both factors would contribute to the longer creep life and lower creep rate observed in this specimen by allowing some deformation to occur initially and by allowing less dominance of the instabilityrecovery processes during creep. In summary, T E M analysis of these creep-rupture tested specimens reveals complex interactions between alloying element effects and processing conditions on the microstructural evolution during creep. The microstructural analysis further indicates that the high

temperature creep behavior of these alloys is being dominated by the solid solution effects of the various alloying elements on dislocation generation-recovery behavior, rather than precipitation effects, which might be expected when C, Nb, and Zr are added. 3.4. Weldability of micro-alloyed iron aluminides Studies of the weldability of iron aluminides have identified a general propensity for cold cracking and a specific dependence of hot cracking on minor changes in alloying composition [13, 28]. Initial work on the weldability of alloy FA-129 showed that the alloy could be welded by controlling the welding process and parameters [13]. This composition had a good resistance to hot cracking, withstanding a stress of 172 MPa in the Sigmajig test without cracking (Table 6). The results also indicated that cold cracking, cracks which develop several hours or days after welding (Fig. 4),

TABLE 6. Stress to cause hot cracking in selected iron aluminide alloysa Alloy

Threshold cracking stress (MPa (klbf in- 2)

FA-97b FA-129c FA-167c FA-168c FA-169b FA-170 b

,~ 138 (20) 172(25) 138 (20) ~ 138 (20) < 103 (15) 103-138 (15-20)

FA-174c FA-175c FA-176 c FA-177c FA-178c

-> 103 (15) -> 138 (20) 103-138 (15-20) 103-138 (15-20) 103-138 (15-20)

aDetermined by the Sigmajigtest. bFinal processing at 850 °C. CFinal processing at 600-650 °C.

Fig. 4. Optical micrograph showing cold cracking in a weld of alloy FA-129.

C. G. McKamey et al.



Alloying additions in Fec4l-based alloys

could be eliminated on laboratory specimens by using a minimum preheat of 200 °C and a post-weld heat treatment of 400 °C for 1 h. Pre- and post-weld heat treating, in general, may lower residual stress as a result of differential thermal expansion and most certainly drives off residual hydrogen. The preliminary results on the Sigmajig threshold cracking stresses measured for the various alloys used in the present study are presented in Table 6. For most alloys there was not enough material to identify the exact TCS. However, for some alloys it was possible to bracket the TCS and for others to show that it was greater or less than some selected starting stress. The results clearly showed that the weldability of iron aluminides is sensitive to changes in composition. Alloy FA-97 which contained larger amounts of Mo, Zr and B exhibited very poor weldability compared with alloy FA-129. Previous welding studies on Fe3Al-based compositions have suggested that B is deleterious to weldability above some nominal level [13, 28]. B has also been found to be deleterious for the welding of FeAIbased alloy compositions [29]. Likewise too much Zr appears to promote hot cracking during welding. Among the alloys tested, all alloys except FA- 168, - 169 and -97 appeared to have acceptable hot crack resistance. As noted earlier, FA-168 and -97 had the best high temperature tensile and creep strengths, while FA129, which had good hot crack resistance, had the best RT ductility. These results would seem to indicate that compositional changes which promote good hot crack resistance also promote good tensile ductility, but are detrimental to strength. Thus hot cracking may not be a problem for all of these alloys and cold cracking may be prevented by a proper combination of pre- and post-weld heat treatments, as is commonly done with many commercial ferritic alloys [13]. These preliminary weldability results are encouraging and suggest that some of the advanced Fe3Al-based alloys may have comparable or better weldability than FA-129. Figure 5 shows that the TCS for the Fe3Al-based alloys falls within the range of values for austenitic stainless steels and for some other aluminides.

4. Summary The mechanical properties and weldability data presented in this study indicate that, in addition to a dependence on microstructure, the properties of iron aluminide intermetallic alloys are very sensitive to alloying additions. The properties are affected through both solid solution and precipitation mechanisms, and, with increasing complexity of composition, synergistic effects between the various alloying elements become more important. For this alloy system, the results show













304 SS

316 SS

Fig. 5. Comparison of values for TCS (for initiation of hot cracking during welding) of Fe3Al-based alloys with those of other common alloys.

that additions of Mo, Nb, Zr, C, and B produce increased tensile and creep strength, but at the cost of room temperature tensile ductility and weldability. The development of one composition will require that trade-offs be made regarding the acceptable values of all three properties. For example, creep resistance of over 200 h (at 593 °C and 207 MPa) and tensile strengths of 500 MPa or more (at temperatures of RT to 600 °C) can be attained if the intended application will allow for lower RT ductility and weldability. To produce one composition with an acceptable combination of properties, the results of this study indicate that (1) levels of 0.03-0.05 at.% C are acceptable for both creep resistance and weldability, (2) B may promote creep resistance, but should be kept at 0.005% or less for weldability, (3) for weldability, the Zr level should be kept below 0.05%, (4) the Mo level needs to be at least 0.4% for creep resistance, but not more than 1% for good RT ductility, and (5) Nb levels of near 0.5% are desirable for strength. It is expected that, with optimization of fabrication techniques, some of the better alloy compositions from this study can exhibit a better combination of tensile and creep properties and weldability than they do at present, as well as adequate corrosion resistance.

Acknowledgments The authors would like to thank Hershel Pierce for testing and compilation of data for this report, Wade Jones for preparation of TEM specimens, and Drs. E. P. George and S. Viswanathan at Oak Ridge National


C. G. McKamey et al.


Alloying additions in Fe c41-based alloys

L a b o r a t o r y for reviewing the manuscript. This research was sponsored by the Fossil Energy A R & T D Materials Program and by the Assistant Secretary for Conservation and Renewable Energy, Office for Industrial Technologies, A d v a n c e d Industrial Concepts (AIC) Division, A I C Materials Program, US Department of Energy under Contract DE-AC058 4 O R 2 1 4 0 0 with Martin Marietta Energy Systems, Inc.

References 1 J.H. DeVan, H. S. Hsu and M. Howell, Sulfidation/Oxidation Properties of Iron-Based Alloys Containing Niobium and Aluminum, ORNL/TM-11176 (May 1989) (Oak Ridge National Laboratory, Oak Ridge, TN). 2 C. G. McKamey, J. H. DeVan, P. E Tortorelli and V. K. Sikka, J. Mater. Res., 6 (8)(1991) 1779. 3 C. G. McKamey and C. T. Liu, in R. D. Kane (ed.), Proc. ADVMAT/91, 1st Int. Symp. on Environmental Effects on Advanced Materials, National Association of Corrosion Engineers, Houston, TX, 1992, paper 17-1. 4 T. B. Massalski (ed.), Binary Alloy Phase Diagrams, American Society for Metals, Metals Park, OH, 1986. 5 J.R. Knibloe, R. N. Wright and V. K. Sikka, 1990Advances in Powder Metallurgy, Vol. 2, Metal Powder Industries Federation, Princeton, NJ, 1990, p. 219. 6 D. Hardwick and G. Wallwork, Rev. High Temp. Mater., 4 (1978)47. 7 P. Tomaszewicz and G. R. Wallwork, Rev. High Temp. Mater., 4(1978) 76. 8 M. G. Mendiratta, S. K. Ehlers, D. M. Dimiduk, W. R. Kerr, S. Mazdiyasni and H. A. Lipsitt, in N. S. Stoloff, C. C. Koch, C. T. Liu and O. Izumi (eds.), High Temperature Ordered Intermetallics II, Materials Research Society, Pittsburgh, PA, 1987, p. 393. 9 D. M. Dimiduk, M. G. Mendiratta, D. Banerjee and H. A. Lipsitt, Acta Metall., 36 (1988) 2947. 10 C.G. McKamey, P. J. Maziasz and J. W. Jones, J. Mater. Res., 7(8) (1992) 2089.

11 R. G. Bordeau, Development of Iron Aluminides, AFWALTR-87-4009, May 1987 (Air Force Wright Aeronautical Laboratories, Wright-Patterson Air Force Base, OH). 12 V.K. Sikka, SAMPE Q.,22(4)(1991)2. 13 S.A. David and T. Zacharia, in K. Natesan and D. J. Tillack (eds.), Heat Resistant Materials, ASM International, Materials Park, OH, 1991, p. 169. 14 P. J. Maziasz and C. G. McKamey, Mater. Sci. Eng. A, 152 (1992) 322. 15 C. G. McKamey and D. H. Pierce, Scr. Metall. Mater., 28 (1993) 1173. 16 P.J. Maziasz, C. G. McKamey and C. R. Hubbard, in G. M. Stocks, D. P. Pope and A. E Giamei (eds.), Alloy Phase Stability and Design, Materials Research Society, Pittsburgh, PA, 1990, p. 349. 17 G.M. Goodwin, Weld. Z, 77(2)(1987) 33s. 18 C.G. McKamey, 1988, unpublished results. 19 C. G. McKamey, C. T. Liu, S. A. David, J. A. Horton, D. H. Pierce and J. J. Campbell, Development oflron Aluminides for Coal Conversion Systems, ORNL/TM-10793, July 1988 (Oak Ridge National Laboratory, Oak Ridge, TN). 20 D.J. Gaydosh, S. L. Draper and M. V. Nathal, Metall. Trans. A, 20(1989) 1701. 21 C. G. McKamey and J. A. Horton, Metall. Trans. A, 20 (1989)751. 22 C. T. Liu and E. P. George, Scr. Metall. Mater., 24 (1990) 1285. 23 I. Baker and D. J. Gaydosh, in N. S. Stoloff, C. C. Koch, C. T. Liu and O. Izumi (eds.), High Temperature Ordered Intermetallics II, Materials Research Society, Pittsburgh, PA, 1987, p. 315. 24 R. H. Titran, K. M. Vedula and G. G. Anderson, in C. C. Koch, C. T. Liu and N. S. Stoloff (eds.), High Temperature Ordered lntermetallics, Materials Research Society, Pittsburgh, PA, 1985, p. 309. 25 M.A. Crimp and K. Vedula, Mater. Sci. Eng., 78(1986) 193. 26 D. G. Morris and M. A. Morris, Acta Metall. Mater., 39 (8) (1991)1771. 27 W. R. Kerr, Metall. Trans. A, 17(1986) 2298. 28 S. A. David and T. Zacharia, Weld. Res. Suppl., 72(1993) 2015. 29 P.J. Maziasz, G. M. Goodwin, C. T. Liu and S. A. David, Scr. Metall. Mater., 27(1992) 1835.