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Intermetallics 18 (2010) 1322–1326

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Intermetallics

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Solidification and high-temperature phase equilibria in the Fe–Al-richpart of the Fe–Al–Nb system

O. Prymak, F. Stein*

Max-Planck-Institut fur Eisenforschung GmbH, Max-Planck-Str. 1, D–40237 Dusseldorf, Germany

a r t i c l e i n f o

Article history:Received 22 October 2009Received in revised form11 December 2009Accepted 14 December 2009Available online 13 January 2010

Keywords:A. Ternary alloy systemsA. Iron aluminides (based on FeAl)A. Intermetallics, miscellaneousB. SolidificationB. Phase diagrams

* Corresponding author. Tel.: þ49 211 6792 557; faE-mail address: stein@mpie.de (F. Stein).

0966-9795/$ – see front matter � 2009 Elsevier Ltd.doi:10.1016/j.intermet.2009.12.022

a b s t r a c t

Partial isothermal sections of the Fe–Al–Nb phase diagram at 1000, 1170, and 1300 �C as well as thecorresponding part of the liquidus surface were established experimentally for Nb contents in the rangefrom 0 to 33 at%. This part of the phase diagram is characterized by an extended two-phase field of the(aFe) solid solution (disordered A2) or FeAl (ordered B2) with the Nb(Fe,Al)2 Laves phase. The solubilityof Nb in all Fe–Al phases is small and reaches a maximum in the phase FeAl. A deep eutectic valley runsparallel to the Fe–Al axis across the large two-phase field and continues to the Al corner. The compositiondependence of the eutectic temperature and the exact position of the eutectic valley on the liquidussurface including the Al-rich part were determined. The high-temperature phase equilibria are stronglyaffected by the liquid phase coming from the Al corner. With increasing temperature its phase fieldgrows quickly into the ternary system resulting in a strongly reduced extension of the (aFe)þNb(Fe,Al)2

two-phase field at 1300 �C.� 2009 Elsevier Ltd. All rights reserved.

1. Introduction

A concept to improve the insufficient strength and creep resis-tance of Fe–Al alloys at high temperatures is the addition ofsolution- or precipitation-strengthening alloying elements; seee.g. [1–5]. One of the candidates frequently discussed as alloyingaddition is Nb [6–10]. A prerequisite for understanding andcontrolling the microstructure and the materials properties isa sound knowledge of the solidification behaviour and the phaseequilibria in the ternary system. However, the available informationon the Fe–Al–Nb phase diagram is quite limited. Isothermalsections at 800 �C [11] and 1000 �C [12–14] show that an extendedtwo-phase field exists between the disordered (aFe) solid solution(or its B2-ordered variant FeAl) and the hexagonal C14 Laves phaseNb(Fe,Al)2 at these temperatures. However, no information isavailable on the solid solubility of Nb in these phases, and phaseequilibria at higher temperatures are unknown.

Recently, Palm [15] established the phase equilibria in fiveFe-rich Fe–Al–Nb alloys with 7.5–40 at% Al and a constant Nbcontent of 5 at% for temperatures of 800, 1000, and 1150 �C. Hepresents partial isothermal sections of the Fe corner of the systemfor these temperatures and reports some values for the solubilityof Nb. The present investigations, which are part of

x: þ49 211 6792 299.

All rights reserved.

a comprehensive study of the whole ternary Fe–Al–Nb system,focus on the Fe–Al-rich part of the ternary phase diagramextending the work of Palm to higher Al contents and highertemperatures. Here we report isothermal sections up to 1300 �C aswell as results on the melting and solidification behaviour ofternary alloys with compositions up to 60 at% Al as obtained byelectron-probe microanalysis (EPMA), X-ray diffraction (XRD), anddifferential thermal analysis (DTA).

2. Experimental

Eleven ternary Fe–Al–Nb alloys with maximum nominal Al andNb contents of 60 and 20 at%, respectively, were prepared bylevitation melting, induction melting, or arc melting from highpurity Fe (99.9 wt%), Al (99.999 wt%) and Nb (99.9 wt%) underargon and casting into cold copper moulds. Table 1 gives compo-sitions and melting methods for all investigated alloys. A compar-ison of the nominal compositions with chemical compositions ofselected alloys as analysed by EPMA shows good agreement within�0.5 at%.

Heat treatments were performed for 1000 h at 1000 �C, 168 h at1170 �C, and 100 h at 1300 �C in an argon atmosphere. For the heattreatments at 1000 and 1170 �C, the samples were encapsulated inquartz ampoules which were evacuated and back-filled with argon.After heat treatment, the samples were quenched in iced brine (10%NaCl solution). For the heat treatments at 1300 �C, the sampleswere wrapped in tantalum foil and placed into closed alumina

Table 1Casting method and nominal composition of the investigated alloys as well assolidus and liquidus temperatures Tsol and Tliq as obtained by DTA (three alloys withcompositions given in italics were provided by M. Palm).

Fe [at%] Al [at%] Nb [at%] Tsol [�C] Tliq [�C]

lev.m. 87,5 7,5 5 1380 (1) 1445 (10)lev.m. 80 15 5 1381 (2) 1462 (3)ind.m. 73,5 26 0,5 1471 (1) 1496 (5)ind.m. 72 26 2 1413 (2) 1471 (5)ind.m. 70 26 4 1359 (1) 1446 (2)lev.m. 63 32 5 1327 (1) 1391 (1)lev.m. 57,5 37,5 5 1296 (1) 1358 (4)ind.m. 56 40 4 1275 (2) 1349 (2)lev.m. 55 40 5 1273 (2) 1332 (3)arc.m. 35 60 5 1153 (1) 1164 (6)lev.m. 20 60 20 1156 (1) 1460 (10)

Fig. 1. Solubility of Nb in Fe–Al (0–66 at% Al) in the temperature range 800–1300 �Cincluding literature data for pure Fe [16] and for the range 7.5–40 at% Al and 800–1150 �C [15]. The value for 60 at% Al was obtained after heat treatment at 1170 �C(instead of 1150 �C). The maximum standard deviation for the given values is 0.3 at%.

O. Prymak, F. Stein / Intermetallics 18 (2010) 1322–1326 1323

crucibles. The crucibles were filled up with titanium filings actingas oxygen getter and heated in an argon atmosphere. Rapid coolingwas performed in a jet of argon gas. In case of the two alloys with60 at% Al, where melting starts just above 1150 �C (see Table 1), theheat treatment time at 1170 �C was reduced to 24 h. The alloyFe-60 at% Al-20 at% Nb was also heat-treated at 1300 �C in thepartially liquid state. The time of the heat treatment at thistemperature was reduced to 19 h, and instead of wrapping thesample into tantalum foil, a cylindrical piece of the alloy was cutwhich exactly fitted into an alumina crucible. The crucible wascovered with tantalum foil and placed into a larger alumina cruciblefilled with titanium filings. The complete heat treatment furnacewas evacuated and back-filled with argon four times.

EPMA was performed with a Jeol JXA-8100 instrument workingat 15 kV and 20 nA. Pure Fe, Al, and Nb were used as standards. Theerror of the resulting compositions is �1% relative. For the deter-mination of the phase compositions, at least twelve points weremeasured per phase. Very fine-scaled eutectic/eutectoid regionswere analysed with a defocused electron beam covering areas of upto 20 mm in diameter with one analysis. Microstructures wereobserved on the same instrument in scanning electron microscope(SEM) mode.

DTA measurements were carried out with a Setaram SETSYS-18DTA under an argon atmosphere with samples placed in aluminacrucibles applying standard heating rates of 5 and 10 K/min. Theaccuracy of the temperature calibration is �1 K. The reversibility ofobserved effects was checked by repeated measurements on thesame specimen and by measuring various specimens of the sameas-cast or heat-treated alloy. Calibration measurements wereperformed using certified standards of pure Al, Au, Ni, and Pd.

XRD investigations on crushed and powdered samples wereperformed on a Philips X’PERT system using Cu–Ka radiation ina 2Q range from 10 to 110�.

Fig. 2. Effect of Nb addition on the liquidus temperatures of Fe, Fe-26 at% Al andFe-40 at% Al (the data for binary Fe–Nb and for the two binary Fe–Al alloys are takenfrom [16] and [18], respectively).

3. Results and discussion

3.1. Solubility of Nb in Fe–Al

According to Palm [15], the solubility of Nb in (aFe)/FeAlincreases with Al content at constant temperature (800, 1000 and1150 �C) in the investigated range up to 40 at% Al. This is showntogether with literature data for pure Fe [16] and the results fromthe present investigation in Fig. 1. The present data for the Nbsolubility at 1300 �C confirm the trend of an increase in solubilitywith increasing Al content in the composition range of the (aFe)solid solution. No value could be determined for the alloyscontaining 40 at% Al as they start melting already at about 1275 �C(see Table 1). At higher Al contents, the solubility of Nb strongly

decreases and is below 1 at% in the intermetallic phases Fe5Al8and FeAl2.

For a given composition, the solubility of Nb in the (aFe) solidsolution increases with temperature, whereas in case of the B2-ordered alloy with 40 at% Al, the data of Palm [15] indicate theopposite behaviour. The reason for this trend reversal is not yetclear.

3.2. Effect of Nb on the melting temperatures of Fe–Al

Alloying pure Fe with Nb results in a linear decrease of the liq-uidus temperature as a function of the Nb content until the eutecticcomposition of about 10 at% Nb is reached, see e.g. [17]. Fig. 2 showsthat the same is true for two series of Fe–Al–Nb alloys with fixed Alcontent and Nb replacing Fe. The liquidus temperatures of Fe–Alwith 26 and 40 at% Al (temperatures for the binary alloys are takenfrom [18]) are reduced linearly with about the same slope as for thebinary Fe–Nb case. Nb also lowers the melting temperatures of Al-rich Fe–Al alloys. Whereas the liquidus temperature of binary Fe-60 at% Al is 1231 �C (see [18]), it is reduced to 1164 �C for a ternaryalloy with nominal composition Fe-60 at% Al-5 at% Nb.

Fig. 4. Composition dependence of the eutectic temperature along the eutectic valley(squares: [24], circles: present results).

O. Prymak, F. Stein / Intermetallics 18 (2010) 1322–13261324

3.3. The eutectic valley

It is known from the literature (e.g. [15]) that as-cast Fe–Al–Nballoys in the two-phase region (aFe) solid solutionþNb(Fe,Al)2

ternary Laves phase have microstructures consisting of primarysolidified (aFe) or Nb(Fe,Al)2 precipitates surrounded by a eutecticmixture of these two phases. The (aFe) and Nb(Fe,Al)2 phase fieldsof primary solidification are separated by a eutectic valley which istentatively indicated by Bejarano [13,14] as a straight line runningparallel to the Fe–Al axis at about 11 at% Nb. From a combination ofdata available from the literature [10,14,19,20] and results from thepresent investigations, the exact position of the eutectic valley onthe ternary liquidus surface was determined, see Fig. 3. In Fig. 4, therespective temperatures of the eutectic reaction along the eutecticvalley as measured by DTA are shown. As expected the eutectictemperatures decrease with increasing Al content. However, at lowAl contents the values run through a shallow maximum at about10 at% Al. A similar observation has recently been made for themelting temperatures of the (aFe) solid solution in binary Fe–Alwhich also show a maximum at about the same Al content [18]. Asthe eutectic valley is the boundary line between the primary phasefields of (aFe) and Nb(Fe,Al)2, the application of the Alkemadetheorem [21,22] on the two maxima indicates that there should beanother melting temperature maximum lying within the Nb(Fe,Al)2

primary phase field, the position of which is determined by theAlkemade line through the two maxima. However, as an exami-nation of the Nb(Fe,Al)2 Laves phase itself was beyond the scope ofthe present investigations, this point has experimentally not beenexamined.

Fig. 3 shows that the eutectic valley runs parallel to the Fe–Alaxis at a Nb concentration of 10 at% up to about 40 at% Al. Theanalyses of the alloys with higher Al contents indicate that theeutectic valley continues at least up to 65 at% Al. The exact positionof the invariant reaction line at higher Al contents between theprimary solidification fields of Al3Nb on the one side and Fe2Al5 orFe4Al13 on the other side has not been studied and the line shown inFig. 3 is only tentative. Before finally reaching the Al corner, theinvariant reaction line splits into a eutectic line ending at the binaryFe–Al eutectic L 4 Fe4Al13þAl and a peritectic line ending at thebinary Al–Nb peritectic LþAl3Nb 4 Al.

From examination of the microstructures of the two alloys with60 at% Al containing 5 and 20 at% Nb, respectively, and phaseanalyses by EPMA and XRD, the primary solidifying phases and thephases in the eutectic mixture were determined. Addingthe information from DTA experiments on the same alloys, thecomplete solidification sequence could be established. As anexample, Fig. 5 shows the as-cast microstructure of a piece of theFe-60 at% Al-5 at% Nb alloy. The grains of the primary solidified

Fig. 3. Position of the eutectic valley and fields of primary solidification on the Fe–Al-rich part of the liquidus surface of the Fe–Al–Nb system (squares: [14,19], triangles:[10], star: [20], circles: present results). The question mark on the line between theprimary solidification fields of the Laves phase and Al3Nb indicates that the exactposition of this line for higher Nb contents has not yet been investigated.

phase, which obviously has decomposed during cooling, are clearlyvisible. The analysis of the grain composition indicates that thesegrains have been the Fe5Al8 phase (frequently called 3 phase),which is known from the binary Fe–Al system to be only stable athigh temperatures and decomposes during cooling into FeAl andFeAl2 in a very quick eutectoid reaction which can not besuppressed [23]. EPMA of the two phases in the grains confirmsthat they are FeAl and FeAl2 both containing about 0.3 at% Nb. Thematrix surrounding the former Fe5Al8 grains is a mixture of Al3Nb,FeAl, and FeAl2 where fine-scaled areas of FeAlþ FeAl2 again resultfrom the decomposition of Fe5Al8. The overall composition of theAl3Nbþ (former) Fe5Al8 eutectic-like matrix area is Fe-62.4 at%Al-5.2 at% Nb marking the position of the eutectic valley betweenthe primary solidification fields of Al3Nb and Fe5Al8 in Fig. 3. TheDTA measurements of this sample show that primary solidificationof the Fe5Al8 phase occurs at 1164 �C and that solidification isalready completed at 1153 �C resulting in the three-phase equilib-rium Al3Nbþ FeAlþ Fe5Al8. At 1110 �C, Fe5Al8 starts to decomposeand at 1080 �C, the final eutectoid reaction with the formation ofFeAl2 occurs resulting in the three-phase equilibriumAl3Nbþ FeAlþ FeAl2.

Fig. 5. Microstructure of the as-cast alloy Fe-60 at% Al-5 at% Nb (SEM back-scatteredelectron micrograph) showing the primary 3 grains, which have decomposed into FeAl(grey)þ FeAl2 (dark) during cooling, and the surrounding FeAlþ FeAl2þ Al3Nb (light)matrix.

O. Prymak, F. Stein / Intermetallics 18 (2010) 1322–1326 1325

3.4. Isothermal sections

Fig. 6 shows partial isothermal sections of the Fe-Al-rich part ofthe ternary system at 1000, 1170, and 1300 �C as obtained fromEPMA, XRD, and DTA. The phase boundaries of the (aFe)/FeAl andthe Nb(Fe,Al)2 phase fields in the 1000 and 1170 �C sections weredrawn according to the results reported by Palm [15] on samplesannealed at 1000 and 1150 �C. As the present investigations wereperformed at 1170 instead of 1150 �C, the two phase boundaries aregiven as broken lines. However, the deviation from the true posi-tion can be assumed to be only small as is also indicated bya comparison with the sections for 1000 and 1300 �C. The whitecircles in the 1170 and 1300 �C sections mark data obtained frompartially or completely molten specimens from the two alloys withnominal compositions Fe-60 at% Al-5 at% Nb and Fe-60 at%Al-20 at% Nb. The partially molten samples were found to beseparated into different, macroscopically homogeneous parts ofdifferent composition after the heat treatment where then eachpart could be dealt with as a new sample.

The extended (aFe)/FeAlþNb(Fe,Al)2 two-phase field knownfrom the existing isothermal sections up to 1150 �C becomes muchsmaller at 1300 �C because the liquid phase field originating fromthe Al corner extends already far into the ternary system. TheB2-ordered FeAl phase does no longer form equilibria with the

Fig. 6. Partial isothermal sections of the Fe–Al-rich part of the Fe–Al–Nb system at 1000,1170, and 1300 �C. The squares in the 1000 section indicate data taken from [15], blackand white circles as well as triangles are from the present investigations. The black andwhite circles represent the alloy and phase compositions of the equilibrated alloys, whitecircles indicate data from partially or completely molten samples, and the black trianglesmark the intersection with the eutectic valley (as is explained in the text).

Laves phase but instead coexists with the liquid phase at 1300 �C.The Fe-rich end point of the liquid phase field, which forms a three-phase field with the (aFe) solid solution and the ternary Lavesphase, follows the eutectic valley which has been discussed in thepreceding Section 3.3. The position of this point, which is markedwith black triangles in the isothermal sections at 1170 and 1300 �C,was determined first by interpolating the data in Fig. 4 to find the Alcontent and then reading the exact position from Fig. 3. Withincreasing temperature the liquid phase grows more and more intothe ternary system and with that cuts the three-phase fields con-necting the Fe–Al boundary with the Al–Nb phase Al3Nb one afterthe other, until at 1170 �C no more equilibria between phases of theFe–Al and Al–Nb system exist.

4. Summary

By EPMA, XRD, and DTA investigations, high-temperature phaseequilibria and the solidification and melting behaviour of Fe–Al-rich Fe–Al–Nb alloys were studied. Three partial isothermalsections 1000, 1170, and 1300 �C as well as the corresponding partof the liquidus surface were established. The solubility of Nb in allFe-Al phases is small and reaches a maximum in the B2 FeAl phaseof about 3.5 at% at 800 �C. The reason for the decrease of the Nbsolubility in B2 FeAl with increasing temperature which wasalready observed by Palm [15] is not yet clear. The addition of Nb toFe–Al lowers the melting temperature for all investigated Fe–Alcompositions. In case of the (aFe) solid solution and FeAl, thisdecrease can be described as a linear function of the Nb contentwith its slope being independent of Al content. At Nb contents ofabout 10 at%, the liquidus temperatures reach a minimumbelonging to the eutectic valley which characterizes the extended(aFe)/FeAlþNb(Fe,Al)2 Laves phase two-phase field. The eutecticvalley runs parallel to the Fe-Al axis and continues to the Al cornerwhen leaving the (aFe)/FeAlþNb(Fe,Al)2 two-phase field. Thehigh-temperature phase equilibria are strongly affected bythe liquid phase which grows quickly from the Al corner into theternary system and reaches a value of 35 at% Al at 1300 �C.

More detailed investigations will be performed to answer theopen questions and to complete the isothermal sections as well asthe liquidus surface and the reaction scheme of the whole Fe–Al–Nb ternary system.

Acknowledgements

The authors would like to thank Mr. G. Bialkowski (samplecutting), Mrs. H. Bogershausen (preparation of metallographicsections), Dr. M. Palm (provision of three samples), Mr. S. Voss(preparation of alloys), and Mrs. I. Wossack (EPMA analyses).Financial support by the Max Planck Society within the frameworkof the inter-institutional research initiative ‘The Nature of LavesPhases’ is gratefully acknowledged.

References

[1] Morris DG, Munoz-Morris MA, Chao J. Intermetallics 2004;12:821.[2] Morris DG, Morris-Munoz MA, Baudin C. Mater Res Soc Symp Proc

2005;842:27.[3] Morris DG, Munoz-Morris MA, Gutierrez-Urrutia I, Requejo LM. Intermetallics

2009;17:404.[4] Palm M. Intermetallics 2005;13:1286.[5] Palm M. Int J Mater Res 2009;100:277.[6] Dimiduk DM, Mendiratta MG, Banerjee D, Lipsitt HA. Acta Metall

1988;36:2947.[7] Stein F, Schneider A, Frommeyer G. Intermetallics 2003;11:71.[8] Morris DG, Munoz-Morris MA, Baudin C. Acta Mater 2004;52:2827.[9] Morris DG, Munoz-Morris MA, Requejo LM, Baudin C. Intermetallics

2006;14:1204.[10] Park JM, Kim KB, Kim WT, Lee MH, Eckert J, Kim DH. Intermetallics 2008;16:642.[11] Burnashova VV, Ryabov VR, Markiv VYa. Dop Akad Nauk Ukr RSR/A 1970;1970:747.

O. Prymak, F. Stein / Intermetallics 18 (2010) 1322–13261326

[12] Raman A. Z Metallkd 1966;57:535.[13] Bejarano JMZ. Estudo do diagrama ternario ferro-alumınio-niobio, PhD thesis,

Universidade Estadual de Campinas, Brazil; 1993. pp. 1-148.[14] Mota MA, Coelho AA, Bejarano JMZ, Gama S, Caram R. J Crystal Growth

1999;198–199:850.[15] Palm M. J Alloys Compd 2009;475:173.[16] Paul E, Swartzendruber LJ. In: Okamoto H, editor. Phase diagrams of binary

iron alloys, monograph series on alloy phase diagrams. Materials Park, OH:ASM Int; 1993. p. 247.

[17] Okamoto H. J Phase Equilib 2002;23:112.

[18] Stein F, Palm M. Int J Mater Res 2007;98:580.[19] Mota MA, Coelho AA, Bejarano JMZ, Gama S, Caram R. J Alloys Compd

2005;399:196.[20] Milenkovic S, Palm M. Intermetallics 2008;16:1212.[21] van Rijn van Alkemade AC. Z Phys Chem 1893;11:289.[22] Malakhov DV. Calphad 2004;28:209.[23] Stein F, Vogel SC, Eumann M, Palm M. Intermetallics 2010;18:150.[24] Mota MA. Obtencao e caracterizacao de materiais compositos in situ no sis-

tema Fe–Al–Nb, Master thesis, Universidade Estadual de Campinas, Brazil;1998. pp. 1-78.