Microstructure and mechanical properties of Fe3Al-based alloys with strengthening boride precipitates

Intermetallics 15 (2007) 1172e1182 www.elsevier.com/locate/intermet

Microstructure and mechanical properties of Fe3Al-based alloys with strengthening boride precipitates R. Krein a,*, A. Schneider b, G. Sauthoff a, G. Frommeyer a a

Max-Planck Institut fu¨r Eisenforschung GmbH, Max-Planck-Str. 1, D-40237 Du¨sseldorf, Germany b Salzgitter Mannesmann Forschung GmbH, Ehinger Str. 200, D-47259 Duisburg, Germany

Received 21 December 2006; received in revised form 6 February 2007; accepted 7 February 2007 Available online 2 April 2007

Abstract Five quaternary FeeAleBeM (M ¼ Ti, Hf, Zr, V, W) alloys based on Fe3Al with strengthening boride precipitates were produced by vacuum induction melting. The alloys were investigated with respect to their microstructure and mechanical behaviour up to 1000  C. The mechanical properties were determined by tensile tests, 4-point-bending tests, high-temperature compression tests up to 1000  C as well as creep tests at 650 and 750  C. Microstructural and phase analysis were carried out by light optical microscopy, scanning electron microscopy, X-ray diffraction analysis and differential thermal analysis. The alloys were tested in the as-cast state, after homogenisation at 1200  C for 48 h and after annealing at 800  C for 624 h. Compared to a corresponding binary alloy the examined alloys exhibit significantly improved mechanical high-temperature properties as well as stable microstructures without considerable loss of ductility. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Iron aluminides, based on Fe3Al; B. Mechanical properties at high temperature; B. Precipitates

1. Introduction Iron aluminides based on the D03-ordered Fe3Al and the B2-ordered FeAl belong to the most extensively studied Fee Al alloys for structural applications [1e3]. They are well known for their outstanding high-temperature corrosion resistance in oxidising and sulphidising environment, their low density and their low cost of the constituents [4e6]. In view of their brittleness at ambient temperatures and the sharp yield stress drop above approximately 600  C it is necessary to improve the room temperature ductility and high-temperature strength to utilize these alloys as structural material for high-temperature applications. Over the past years studies have shown that iron aluminides are intrinsically ductile and their brittle behaviour is caused by environmental effects, i.e. moisture induced hydrogen

* Corresponding author. Tel.: þ49 211 6792 571; fax: þ49 211 6792 299. E-mail address: [email protected] (R. Krein). 0966-9795/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2007.02.005

embrittlement [7e13]. However, environmental embrittlement can be minimized, e.g. through careful control of grain size, composition or alloying especially with chromium [7,9,14e17]. There are various possibilities for strengthening of iron aluminides like solid solution hardening or strengthening by carbides and Laves phase which have been addressed in particular [1,18e25]. Fe3Al-based alloys with TiB2 precipitates produced through a rapid solidification process or powder metallurgy [1,26e29] as well as by conventional arc melting and casting procedures [1,32] had already been examined with respect to their mechanical properties. The influence of ZrB2 dispersoids on mechanical properties of rapidly solidified FeAl has been studied by Morris and Morris [30,31]. They showed that the addition of ZrB2 particles leads to an increase of strength up to about 600  C as well as to an increase of ductility. Doucakis and Kumar [33] investigated the thermal stability of metal diboride precipitates (TiB2, ZrB2, TaB2 and NbB2) in Fe3Al-based material. They showed that thermodynamically stable boride precipitates can be dispersed in a Fe3Al matrix through conventional casting methods.

R. Krein et al. / Intermetallics 15 (2007) 1172e1182 Table 1 Nominal composition (at.%) of the studied Fe3Al-based alloys Alloy no.

Fe

Al

B

Ti

Hf

Zr

V

W

1 2 3 4 5

Balance Balance Balance Balance Balance

30 26 26 26 26

1 1 1 1 0.75

0.5 e e e e

e 0.5 e e e

e e 0.5 e e

e e e 0.5 e

e e e e 0.75

The aim of this study was to examine the effect of various boride precipitates on the microstructure and the mechanical properties of Fe3Al-based alloys. The alloys were produced using a conventional casting method with addition of boron and other transition elements (Ti, Hf, Zr, V and W) to precipitate stoichiometric metal borides, i.e. TiB2, ZrB2, HfB2, VB2 and WB. The aluminium content was fixed at 26 at.% and 30 at.%, respectively. 2. Experimental procedure The alloys were produced by vacuum induction melting in argon inert gas and drop cast into copper moulds to obtain rods with 32 mm in diameter and 186 mm in length. The boron was added by means of a FeeB pre-alloy (18 wt.% B). The other constituents were added as highpurity elements. The nominal compositions of these alloys are listed in Table 1.

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The alloys were tested in the as-cast and as-annealed states. Alloys 1, 2 and 3 were homogenised at 1200  C for 48 h and annealed at 800  C for 624 h. Alloys 4 and 5 were only annealed at 800  C for 624 h. All alloys were furnace cooled after each heat treatment to avoid thermal stresses. For light optical microscopy (LOM) and scanning electron microscopy (SEM) the samples were ground (up to 2400 grit), polished (with 3 and 1 mm diamond spray and finally by an oxide polishing suspension) and etched in a picrid solution (80 ml ethanol þ 20 ml HNO3 þ 1 g picrid acid). SEM and energy dispersive X-ray diffraction (EDX) were carried out with a Hitachi S-530 SEM with an integrated EDAX CPU LEAD at 20 kV and a working distance of 15 mm. The phases were identified by Xray diffraction (XRD) using a Huber X-ray diffractometer with Co Ka radiation (lCo ¼ 1.8970 nm). The XRD measurements were conducted on sheet-like samples from the bulk material of about 2 mm thickness. The critical temperatures of the D03eB2 and B2eA2 transitions were determined by differential thermal analysis (DTA) using a Setaram SETSYS-18 DTA. The samples were heated in alumina crucibles under argon atmosphere at a heating rate of 10 K/min. The accuracy of the temperature calibration was 1 K. The room temperature ductility and the fracture behaviour were observed by means of tension tests with flat samples of 3.5  5 mm2 in cross-section and 15 mm in gage length using a ZWICK 1474 universal testing machine at a strain rate of 104 s1. The temperature dependency of the yield stress

Fig. 1. LOM and SEM micrographs of (a) as-cast and (b) as-annealed Fee30Ale1Be0.5Ti.

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was investigated by means of compression tests. For this 5  5  7 mm3 samples were prepared by electrical discharge machining (EDM). The compression tests were conducted up to 1000  C at a constant strain rate of 104 s1. To examine the brittle-to-ductile transition temperature (BDTT) 4-pointbending-tests were carried out. For this purpose samples with dimensions of 3  6  18 mm3 were bent at a constant strain rate of 104 s1 referring to the outer fibre. The BDTT was defined as the transition from brittle behaviour, i.e. fracture elongations of less than 1% to ductile behaviour with elongations of more than 3% in the outer fibre. Creep properties were evaluated by compression creep tests with stepwise increasing load at 650 and 750  C. The 6  6  10 mm3 samples were prepared by EDM. 3. Results and discussion 3.1. Microstructure and phase identification All investigated alloys exhibit primary Fe3Al dendrites with interdendritic precipitates. The resulting grain size of about 100  50 mm2 seems to be insensitive to thermal treatments, indicating that the particles block the motion of the grain boundaries. Fig. 1 illustrates the microstructure of alloy 1. The as-cast microstructure shows a dendritically solidified matrix with interdendritic Fe3AleFe2B eutectic seams between the grains (Fig. 1a). Fe3Al and Fe2B are clearly identified as the constituent phases of the as-cast alloy by XRD (Fig. 2(a)). Additionally the XRD pattern of the as-cast alloy contains weak lines of the TiB2 phase, which means that some small amounts of this phase are already present in the as-cast alloy. The eutectic dissolves during the heat treatments. In a first step after the homogenisation treatment the Fe2BeFe3Al eutectic changed into spheroidised Fe2B particles. Additionally, TiB2 precipitates of up to 4 mm length and 1 mm thickness were formed. After the annealing treatment, the FeB2 particles are completely dissolved, whereas the TiB2 precipitates did not coarsen considerably. Indeed the XRD pattern of the as-annealed material shows only peaks of Fe3Al and TiB2 (Fig. 2b), indicating that the Fe2B particles, which are apparently not stable, have dissolved and additional stable TiB2 particles are formed during the heat treatments. This fits well with observations by EDX, according to which the Al content of the matrix in the as-cast material is about 1 at.% higher than the global Al content. During homogenisation and annealing the Al content of the matrix decreased and finally reached the global Al content, due to the dissolution of the Fe2B particles with enrichment of Fe in the matrix. An almost identical microstructure is shown by alloys 2 and 3. These alloys exhibit primary Fe3Al dendrites along with acicular HfB2 and ZrB2 particles of approximately 5 mm in length and 0.4 mm in diameter which are located at the grain boundaries (Fig. 3a). The particle size and particle shape seem to be insensitive to thermal treatments (Fig. 3b). The as-cast alloy 4 consists of primary Fe3Al dendrites with Fe3AleFe2B eutectic in-between. During annealing two significant changes occurred. First the morphology of the precipitates changed from a lamellar eutectic structure to spherical

Fig. 2. XRD pattern of (a) as-cast and (b) as-annealed Fee30Ale1Be0.5Ti.

coarsened Fe2B particles and second all V which was primarily dissolved within the matrix diffused into the Fe2B particles as observed by EDX analysis (Fig. 4). Similar to the other alloys, alloy 5 exhibits the Fe3AleFe2B eutectic between the matrix grains (Fig. 5a). W is in solid solution within the matrix both in the as-cast and as-annealed state. Annealing led to precipitation of tiny acicular particles of about 800 nm in length and 500 nm in thickness within the grains (Fig. 5b), whereas the former eutectic disintegrated. Within this study it was not possible to determine the identity of these particles definitely. The XRD results (Fig. 6) reveal that these particles could be some different kind of iron borides or tungsten carbides as they were detected in tungsten containing D03-ordered FeeAl alloys by Sun et al. [34]. However, changes of the tungsten composition within the matrix could not be detected by EDX. 3.2. DTA results The results of the DTA tests are summarized in Table 2. It can be assumed that the critical temperatures wCrit for neither the

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Fig. 3. LOM and SEM micrographs of Fee26Ale1Be0.5Zr (alloy 3) (a) in the as-cast and (b) as-annealed state.

D03eB2 nor the B2eA2 phase transition are affected by the precipitates, i.e. these are only controlled by the matrix composition. Regarding alloys 1e3, the matrix of which consists only of Fe and Al, the examined transition temperatures fit well with the binary phase diagram of Kubaschewski [35]. Only the solidus (wSol.) and liquidus (wLiq.) temperatures of the alloys are slightly decreased, obviously due to the influence of the alloying elements. The determined critical temperatures of alloys 4 and 5 do not fit with the binary phase diagram. This may be explained by the solid solution of third elements like V or W within the matrix, respectively. It is known that these elements raise the critical temperatures of the D03eB2 as well as the B2eA2 transitions [36,37]. Both alloys show the lowest solidus temperatures wSol. resulting from the solidification of the Fe3AleFe2B eutectic. 3.3. High-temperature compression tests Fig. 7 shows the observed temperature dependence of the compressive 0.2% yield strength of the as-cast materials. Alloys 1 and 5 which exhibit relatively low room temperature strength possess the best high-temperature strength (z100 MPa at 800  C). Compared to corresponding binary alloys [22,37] the investigated materials show significantly increased yield strengths above 600  C and smaller strength values at room temperature. All materials exhibit the so called ‘‘yield strength anomaly’’, i.e. a positive temperature dependence of the yield strength, that is usually observed in

Fe3Al-based alloys. The yield strength peak is typically observed between 550 and 600  C followed by a rapid decrease in yield strength [2,15]. However, alloy 1 shows a behaviour that is slightly different from that of the other alloys of this study, i.e. the peak temperature is about 50 K shifted to a higher value. It is known that a shift of the anomaly peak can be caused e.g. by solid solution of Ti within the Fe3Al matrix [37,38]. However, this effect could be excluded by EDX measurements, where Ti was not detectable within the matrix. Therefore it is supposed that this behaviour is related to the non-equilibrium microstructure of the as-cast alloy where three phases are present (see Section 3.1). After homogenisation the yield strength peak of alloy 1 is at the expected temperature of about 550  C (Fig. 8), which supports the above conclusion. The data indicate increased room temperature yield strengths, which may be due to increased amounts of TiB2 after the homogenisation treatment, whereas the high-temperature strength is still on the same level. The room temperature yield strength values of alloys 2 and 3 are still on a very high level, but slightly lower than in the as-cast state. Both alloys show similar yield strength behaviour above 600  C. Since there are no significant changes in the microstructure of these alloys after homogenisation, the lower yield strength values are probably due to recovery effects such as reduction of internal stresses or annihilation of vacancies. Fig. 9 shows the yield strength behaviour after annealing at 800  C for 624 h. Apart from alloy 5, all materials exhibit

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Fig. 4. SEM micrograph with the EDX pattern of the matrix of (a) as-cast and (b) as-annealed Fee26Ale1Be0.5Zr as well as (c) SEM micrograph of as-annealed precipitate with the EDX pattern.

reduced room temperature strength values compared to the ascast or as-homogenised state. The significant increase in room temperature yield strength of alloy 5 is caused by the precipitation of very fine particles within the matrix (see Fig. 5). However, these particles are not effective at elevated temperatures. Just like for materials in the as-cast and as-homogenised state, alloy 1 possesses the highest yield strength at elevated temperatures. Alloy 4, the microstructure of which also changed significantly during heat treatment (Fig. 4), shows little change of yield strength. The values are slightly decreased,

which could either result from recovery processes or the mentioned changes in microstructure. It is noted that this alloy shows a surprisingly high yield strength value at 800  C, the reason of which is not yet clear. 3.4. Room temperature tension tests Table 3 lists the values of yield strength and fracture strain from room temperature tension tests. All alloys exhibit a fracture strain below 1%. The highest fracture strain was shown by

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Fig. 5. LOM and SEM micrographs of (a) as-cast and (b) as-annealed Fee26Ale0.75Be0.75W.

alloys 1 and 5 in the as-cast and as-annealed state. Alloys 2 and 3 partially fractured already during elastic elongation. It is noted that these alloys contain acicular precipitates (see Fig. 3) which may act as internal notches and hence may be the reason for this poor ductility. The values of the yield strength measured in tension are about 20e40% lower than in compression. This asymmetry was also detected in earlier studies of FeeAl alloys. Details were given in Refs. [22,39]. Fracture surfaces of two as-cast and as-annealed alloys after room temperature tension tests are shown in Fig. 10. The fractographs show transgranular cleavage, which occurs in all investigated alloys in both as-cast and as-annealed states. Transgranular cleavage is also observed in corresponding binary alloys [22,15], indicating that the boride precipitates do not weaken the grain boundaries as was observed in some carbide strengthened Fe3Al-based alloys where the fracture mode changed partially to intergranular cleavage [22]. However, there is a difference in scale of the cleavage facets in the as-cast state and after heat treatments, e.g. Fig. 10 (a) and (b), which is supposed to results from the change of the particle shape. 3.5. Creep tests Fig. 11 shows secondary creep rate data as a function of the applied stress at 650  C, i.e. in the B2-range of the as-cast

materials. The creep data follow Norton’s power law as already reported earlier [40] and the stress exponents n were evaluated accordingly. It is noticeable that alloys 1 and 5 exhibit a change of slope in the double-logarithmic stress vs. creep rate plot. Alloy 5 possesses a stress exponent of n ¼ 3.2 within a stress range of 50e75 MPa and 8.3 above. Alloy 1 exhibits an n value of 8.9 up to 80 MPa and 5.6 above. The other alloys show no change of slope. Alloy 2 has an exponent of n ¼ 5.0, for alloy 3 n ¼ 5.7 and for alloy 4 n ¼ 4.8. Since there were significant changes of microstructure after heat treatments, the slope changes in the stress vs. creep rate plots of alloys 1 and 5 are believed to result from these microstructure changes. The data for the homogenised alloys in Fig. 12 indicate a constant stress exponent n ¼ 4.0 for alloy 1. This confirms the supposition that the stress exponent variation of alloy 1 in Fig. 11 is due to microstructural changes. The hafnium alloyed material (alloy 2) again gives n ¼ 4.0 (with a deviating data point at 85 MPa), whereas alloy 3 gives n ¼ 4.6. At 750  C creep tests with stepwise increased load were only performed for the as-cast titanium and tungsten alloyed materials. The other alloys were only tested with constant load of 25 MPa. The results are presented in Fig. 13. As already visible at 650  C (Fig. 11), alloys 1 and 5 exhibit a stress exponent variation. Relatively high stress exponents (n ¼ 9.1 for alloy 1 and n ¼ 7.9 for alloy 5) were observed for both

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Fig. 7. Temperature dependence of the yield strength of as-cast materials.

increased creep resistance of alloy 1 is due to the precipitation of either TiB2 or Fe2B or both, whereas after the equilibrium annealing only TiB2 could be effective. The creep resistance of alloy 5 with the highest creep resistance at 650  C is attributed to the precipitation of Fe2B as well as to solid solution strengthening of the matrix by tungsten. The positive effect of tungsten on the creep properties of Fe3Al-based alloys has already been shown by Sun et al. [34] and Zhang et al. [25]. Alloy 4 which also contains Fe2B exhibits a relatively low creep resistance. 3.6. Brittle-to-ductile transition temperature (BDTT)

Fig. 6. XRD pattern of (a) as-cast and (b) as-annealed Fee26Ale0.75Be 0.75W.

Fig. 14 illustrates the results of the 4-point bending tests for the investigated as-cast alloys. The transition from brittle to ductile behaviour was observed between 100 and 200  C.

alloys in the lower stress range between 25 and 45 MPa. At higher stresses both alloys possess the same n value of 5.4. It is again believed that these n changes result from changes of microstructure during the creep test. The observed stress exponents indicate dislocation creep in agreement with earlier work on similar alloys [40]. Compared to a binary alloy with a similar composition (Fee27.6Al) that was examined by Sundar et al. [41], the investigated alloys exhibit significantly improved creep resistances. Obviously the

Table 2 Transition temperatures wCrit for ordering and melting evaluated by DTA Alloy no.

wCrit D03eB2 [ C]

wCrit B2eA2 [ C]

wSol [ C]

wLiq [ C]

1 2 3 4 5

471 548 547 561 562

1062 841 833 905 929

1358 1380 1374 1203 1160

1439 1466 1474 1474 1481

Fig. 8. Temperature dependence of the yield strength of as-homogenised materials.

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Table 3 Yield strength and fracture strain evaluated by room temperature tension tests Alloy no.

1 2 3 4 5

Fig. 9. Temperature dependence of the yield strength of as-annealed materials.

Compared to a corresponding Fee26Al alloy [42] the BDTT did not increase significantly through the precipitate hardening of the present alloys. During bending serrated yielding accompanied by frequent sound emissions were detected in all alloys between 200 and 500  C. Fig. 15 shows examples of the bending behaviour of alloy 1. It can be seen that at 250  C (Fig. 15 (b)) small

Yield strength [MPa]

Fracture elongation [%]

As-cast

As-annealed

As-cast

As-annealed

398 435 585 378 428

345 435 472 356 465

0.8 e 0.2 0.5 0.6

0.6 0.2 0.1 0.5 0.8

frequent stress serrations occur in the region of plastic deformation, which become most pronounced at 400  C (Fig. 15(c)), whereas nearly no serrations are observed at 200 and 500  C (Fig. 15 (a) and (d)). Each serration is accompanied by a sound emission. The amplitudes of the stress serrations increase with increasing temperature. The observed serrated yielding in a restricted temperature range is characteristic for the PortevineLe Chatelier (PLC) effect. Acoustic emissions have been observed and related to the PLC effect e.g. in AleMg and CueMg alloys [43,44]. Serrated yielding was also reported for binary FeeAl alloys with lower Al contents (up to 17 at.% Al) in high-temperature compression tests, which was attributed to the interaction of mobile dislocations and short-range ordering and disordering reactions [45]. In

Fig. 10. SEM fractographs after tensile testing of (a) as-cast and (b) as-annealed Fee30Ale1Be0.5Ti as well as (c) as-cast and (d) as-annealed Fee26Ale0.75Be 0.75W.

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Fig. 11. Secondary creep rate as a function of stress at 650  C of as-cast materials.

Fig. 13. Secondary creep rate as a function of stress at 750  C of as-cast materials.

D03-ordered Fee25Al single crystals which were investigated by Engelke and Neuha¨user [46] the stress serrations, which were observed between 80 and 500  C were explained by means of a pipe diffusion process of vacancies or solutes into the dislocation core. Furthermore, Briguet and Morris [47] reported about yield point effects during unloading at intermediate temperature in a mechanically alloyed B2-ordered Fee40Al material, which were ascribed to be due to dislocation locking by spreading the cores of a/2<111> dislocations. However, the serrated yielding was not observed for the compression tests of the studied alloys with similar elongations which needs further study.

between room temperature and 1000  C, fracture at room temperature, BDTT as well as creep at 650 and 750  C. It was found that Ti, Hf and Zr form stable metal diborides which lead to significantly improved strengthening at elevated temperatures as well as higher creep resistances compared to corresponding binary alloys. V and W do not form metal borides, but improve strengthening through solid solution hardening. The precipitate particles prevent grain coarsening during the annealing treatments and have only slight effect on ductility which appears promising for future alloy developments. Annealing of the alloy with tungsten leads to the appearance of finely dispersed precipitates, the character and effect of which is not yet clear. All investigated alloys showed serrated yielding with sound emissions between 200 and 500  C in bending corresponding to the PLC effect, but not in compression.

4. Conclusions Fe3Al-based alloys with strengthening boride precipitates were investigated with respect to microstructure, yield strength

Fig. 12. Secondary creep rate as a function of stress at 650  C of as-homogenised materials.

Fig. 14. Total elongation (referred to the outer fibre) in 4-point bending of ascast materials. The dashed lines indicate the BDTT range. The arrows indicate total elongations of more than 3%.

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Fig. 15. Stressestrain curves in 4-point-bending for Fee30Ale1Be0.5Ti at (a) 200, (b) 250, (c) 400 and (d) 500  C with pronounced serrated yielding at 250 and 400  C.

Acknowledgments The authors would like to thank Mr. K. Markmann for casting the alloys, Mr. G. Bialkowski for preparing the specimens and mechanical testing and Mrs. I. Wossack for carrying out the X-ray diffraction experiments. The authors are grateful to Dr. M. Palm and Dr. F. Stein (MPIE) for helpful discussions.

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