0 0 7 0 5 7 4

the:csf i l O T any of their employees. irtdkes any warran!)' assumcs z n y I-gal Iishhty or responsibility for iha accur ussf-iness of any information, apparntils. prodiii:i, or

r:p:esen!s that Its u s e x c ! d not l i i f i mge pl ivately o w w d _rlghis Zc+c%?ncs herc:n to a n y specific c o i i i r i ~ c , ~ ~ a : pr0diJCt process, or service ad e name, trade ^ii-I a ik , srzr, or othewiss, does not necessarily coiris;,;Litr: e: imply its endation, or favoring by the United States Governnien! 3r any agency thereof 1 ^he views and opliiiuns of ahikors expressed herein do not necessarily state or reflect thoseof thcUnitnd StatesGoveiaEent or any sgency ths:nrJf

METALS ^AND CERAMICS DIVISION

EVALUATION OF MECHANICAL AND METALLURGICAL PROPERTIES OF Fe,Al-BASED ALUMINIDES

C. G. McKamey, C. T. Liu, J. V. Cathcart, S . A. David, and E. H. Lee

Uate Published: September 1986

NOTICE: This document contains information of a preliminary nature. It is subject to revision or correction and therefore does not represent a final report.

R e s e a r c h s u p p o r t e d by t h e Morgantown E n e r g y Technology C e n t e r , U.S, Department (of ^Energy

Prepared by the

OAK RIDGE NATIONAL LABORATORY Oak Ridge, Tennessee 37831

operated 5y

MARTIN MARIETTA ^{ENERGY SYSTEMS, INC.}

for the

U.S. DEPARTMENT ^OF ENERGY

under Contract No. DE-AC05-840R21400

3 4 4 5 6 0 0 7 0 5 7 4 3

ABSTRACT

. . .

¹

INTRODUCTION

. . .

²

BACKGROUND

. . .

³

EXPERIMENTAL PROCEDURES

. . .

⁵

ALLOY PREPARATION

. . .

⁵

CHARACTERIZATION OF ALLOYS

. . .

6

TESTING PROCEDURES

. . .

⁷

Tensile Testing

. . .

7

Oxidation Studies

. . .

⁸

Sulfidation Studies

. . . . . .

⁸

Welding Studies

. . .

⁸

Microstructural Studies

. . .

⁹

Microprobe Studies

.

Fractography

. . .

RESULTS AND DISCUSSION

. . .

MICROSTRUCTURAL STUDIES

.

TENSILE PROPERTY STUDIES FRACTOGRAPHY

. . .

OXIDATION STUDIES

. . . .

SULFIDATION STUDIES

. . .

WLDABILITY

. . .

CONCLUSIONS AND FUTURE WORK

.

REFERENCES

. . . . . .

⁹

. . .

⁹

. . .

⁹

. . .

⁹

. . .

¹¹

. . .

¹⁴

. . .

¹⁸

. . .

¹⁹

. . .

²⁰

. . .

²¹

. . .

²²

iii

C . G. McKamey, C. T. Liu, J. V. Catlicart, ^{S .} A . David, and E. ^{H .} Lee

ABSTRACT

A study is being conducted to develop aluminide alloys based on Fe3Al with an optimum combinatiori o f strength, ^duc- tility, and corrosion resistance for ^use a3 hot components in advanced fossil energy conversion systems. ^Three phases of the study have been planned: ^{( 1 )} Fe-A1 base compositions will be prepared ^{f o r} pra?liminary studies and evalluntican of potential ^for further alloy development; (2) two of the base alloys will be used in the development and characterization of the p o p e ~ k i e ~ of ternary and quaternary allays; and ( 3 ) based on the results of those t w o phases ^of the study, allays will ^be selected for preparation and characterization of large ^heaks. Studies will include the fabricability, micrsstructmres, tensile propexties, oxidation and sulfidation resjstance, and weldability of

Fe3Al-based alloys to which a fine dispersion of TiBz ^particles has been added ^{f o r} g r a i n r e f i n e m e n t . T h i s r e p o r t summarizes Lhe results of the f i r s t phase of ^{t h i s} study ^and d i s c u s s e s OUT p l a n s for Eutiire work.

0.5 wt % TiR,) were chosen as base alloys. They were p-epared by arc melting under- argon and ^{d r ~ p} casting i n t o water-c,ooled copper molds, followed by a homug~nizatioo anneal They were hot rolled starting at 1 0 0 0 ~ ~ arid finishing at ~ ~ Q O C , tht?n warm ro:led at 680"~. A I I six compositions were easily fabt-icated and exhibited ~xcellent oxidatim. and corrosiim properties.

Negligible weight changes were obic.,rved after ⁵⁰⁰ b in ^{a i r -} oxidation t e s t s at tsmperatures 4-3 I O O O ~ C ^*

s

itxi l a r results occurred in sulfidat.iozi t e s t s conducted at $ 7 1 " ~ for ¹⁶⁸ b.

Resistance to oxidizing ^and sulfi3iziwg environments ^{is R} result of the formation of a self-protecting oxide layer at low oxygen pressures. The tensile strengths of ^{t h e} s i x base alloys were shown to be higher than those of type 316 stainless steel. at temperatures below 7 6 0 ~ 6 and ^those of modified W K - ~ M O steel at temperatures above 55oQC. However, room-temperature ductility is slightly higher for alloys containing more than 27 at. X A I . Preliminary weldability studies indicate that the alloys with higher aluminum content develop Z e w e r cracks in ^the fusion and heat-affected zones.

Alloys of ^{FEY-24, 25, 26,} 27, 28, and ³⁰ at. % ^{A l} (a91 with

Wesearch supported by the U.S. Department of Energy, Surface

Gasification Materials Program, Morgantown Energy Technology Center, under contract No. DE-AC05-840R21400 with ^Ma-ctin Marietta Energy Systems, Inc.

1

Based on considerations of strength, corrosion resistance, fabricability, and weldability, alloys ^of: Fe-28 and ³⁰ at. ^% A1 have been selected for further developmenl. Future work will include alloying with molybdcnum, titanium, and zirconium, singly and in combinations, to improve the high-temperature strength and room-temperature ductility.

INTRODUCTION

Iron alumini des based on Fe,Al are ordered intermetallic alloys that generally have good oxidation and corrosion resistance and low material cost. They also conserve strategic materials such as chiorciaam, ^and they have a lower density than stainless steels arid therefore a better

strength-tu-wcight ratio. However, limited ductility at ainbient tempera- tures and a sharp drop in strength above 6QO°C are major disadvantages to their use as structural m a t e r i a l s . The goal of the present study is to develop alurninide a l l o y s based ^on Fe3A1 with an optimutii combination of strength, ductility, and corrosion resistance for use as hot components in advanced fossil energy conversion systems (such as heat exchangers).

Studies invulve the fabricability, microslructure, tensile proper- ties, oxidation and sulfidatioii resistance, and welding characteristics of Fe,Al-based alloys with ternary and quaternary additions ^of molybdenum, titanium, zirconium, or other elements. Because previous studies showed that dispersions of TiB, particles served to improve the mechanical prop- erties of F~,AI via refinement of grain structitre,',2 ^0.5 wt % ("1 at. %)

of Till2 was added to alloys of Fe-24, 25, 26, 27, 28, and ³⁰ at. ^{% Al} to formulate ^{t h e} base alloys for study. This report summarizes the results to date on these base alloys, presents justification for the selection of base alloy compositions €or f u r t h e r development, and discusses ^{o u r} plans f o r future work.

BACKGROUND

.

Currently, heat-resistant alloys are either nickel-based or high- nickel-content steels containing a delicate balance ^of one or more strate- gic elements such ^as Cr, Co, Nb, Ta, and W to obtain oxidation resistance, high strength, adequate ductility, fsbricability, and thermal stability.

In spite of their high degree of dev$lopment, these state-of-the-art alloys do not meet the desired eharazteristics ^of the hot components for advanced fossil energy conversion systems because ^{o f} their susceptibility to catastrophic hot COSKOS~OII by environments containing SO2 and ^{S O 3}

(i.e., combustion gases), M2S (coal gasification plants), and alkali sulfates (gas turbines). These alloys arc also expensive to produce and suffer from aging embrittlement and chromium evaporation at high

t e m p e r a t U K @ s .

Recause of their ability to forin protective aluminum oxide scales,3 ^{q 4} iron aluminides near the Fe3A1 composition are expected to meet one of the major requirements of high-temperature components for use in environments that ^cause oxidation, s i i l f i d a t i o n , a7d carburization. Since A1203 films form _at very low oxygen partial pressuresg tlw potential applications i n c l n d e components for coal ^gas; f icatian

,

fluidizcd-bed combustion, gas- cooled reactors, ^{g a s} tiirbiries, curre-et collectors ^{f o r} fuel cells, and the

seed recovery sections of magnetohydrodynamics systems. Their rcsistance to hot corrosion by SO2, ^SO3, sulfat?~, and ^H7S is ^expected to be high because the impervious A 1 2 0 3 scale is not susceptible to fermation of a low-melting phase such ^as the Ni-Ni 3S2 eutectic (melting point ^{=: 650'6)} observed in sulfidizcd nickel-based .~lloys.~

therefore is not expected to ^{occur i b r} iron aluminides.

Catastrophic hot ^corrosion

Iron aluminides are projected to be much cheaper than conventional alloys by virtue of the lower cost ^{o f} iron ^and aluminum- Furthermore, the successful substitution of these alloys ^{f o r} the current heat-resistant alloys would reduce our nation's dependence on elements such as chromium and cobalt. Together, those two factors offer alloys based ^on an abundant supply of raw materials and without the sharp price excursions associated with import disruptions of critical strategic elements.

A practical advantage of these alloys ^{n e a r} the Fe3A1 composition for- fixed components ( e . g . , heat exchangers) and r o t a t i n g components ( e . g . , steam turbine components) ^{i s} their lower density ^{( p =: 6 . 6 gecrn-’)} as compared with the steels and nickel-base alloys ^{( p z} 7 . 8 to ^{8 . 5} g-cm-’).

The usefulness of iron aluminides for structural applications is, however, expected to be limited ^by their low room-temperature ductility (-1-2%) and their poor hot strength above ^600OC.

characteristic of cast and fabricated aluminides tested at room tem- perature, ceases to be a problem at high temperatures as the ductility increases to about ^50% or more at 60OoC.*

hot extrusion of powders have been shown to have a room-temperature ^duc- tility of ^about The creep strength of these alloys ^{i s} comparable to t h a t of a 0.15% carbon steal at ^550”C.7

Low ductility, which is

Tron aluminides consolidated by

Recent revi sions of the iron-alumi ^nuni phase diagram (Fig. 1) confirm thc existence of three bcc phases [a disordered solid solution (a),

ordered Fe,A1 ( D O , ) , and ordered FeAl ^{( R 2 ) ]} and the two-phase regions ^{01 +} DO’ and a ^f B2 €or alloys containing ^{24 t o 30} at. ^% . 4 1 . * 9 ’ Transmission

ORM L-DWG 84- 3 6726R

700

-

0 650

W rx

3

600

2

⁵⁵⁰

5 a

W t- 500

450 400

-

/ ^_I OKAMO-1.0 AND BECK

21 22 23 24 25 26 27 28 29 30 31 ³²

ALUMINUM (at. %)

F i g . 1. The iron-aluminum phase diagram showing the phases ^{o f} interest to this study.

electron microscopy, Mossbauer spectroscopy, and X-ray diffraction studies show that the phase transformations resulting from thermal heat treatments generate large coherency stresses and that these transformations proceed by nucleation and growth, spinodal decomposition, continuous ordering, or combinations of those effects.","

results indicate that Fe3A1 alloys with less than 26 at. % A 1 can be age- hardenable through control of precipitation processes.

For the purposes of this study, these

Grain structures in Fe3A1 alloys, prepared by powder metallurgy methods or by conventional melting and casting techniques, have been refined through additions of 1 to ^2% titanium diboride (TiB,)

inclusions,',* produced by adding boron and titanium to the melt during casting. These inclusions effectively strengthen Fe,Al at temperatures to 60OoC.

Most of the data presently available on the iron aluminides was gathered during the 1950s and 1960s on binary alloys a€ approximately 16 to 28 at. % Al. Some studies in alloy development were conducted at that time, and discussions ^of these studies can be found in refs. 12 through 15.

EXPERIMENTAL PROCEDURES

&Lay PREPARATION

Six alloys with compositions shown in Table 1, each containing 0.5 wt X TFB, added for grain refinement, were prepared by arc melting under argon and drop casting into watex-cooled copper molds of size 12.7 by ^{2 5 . 4} by 127 mm. One half of each ^500-g ingot was homogenized for ⁵ h at 1000°C.

approximately ^{0 . 9} mm, starting at lOOO'C and finishing at 650OC.

warm rolling to approximately ^{0 . 7 6 mtn} was done at ^600'C.

schedule produced sheet of uniform thickness, with minor edge ^cracks on all six alloys and a slight curling of one alloy.

The homogenized alloys were then hot rolled to a thickness of Final This .rolling

CHARACTERIZATION OF ^ALLLOYS

The density of each rolled alloy was measured by the Archimedes"

method using toluene as the liquid medium. Results are given in Table 1.

Wet chemistry techniques were used to confirm the accuracy o f the nominal

Table 1. Iron-aluminum alloys presently being studied A1 loy Compositiona Density Grain diameterb designation (at. % Al) (g/cm3

1 (w)

FA-36 ²⁴ 6.62 38

FA-40 25 6.60 52

FA-38 26 6.56 44

FA-41 27 6.52 55

FA-37 28 6.49 46

FA-39 30 6.42 47

"A11 alloys contain 0.5 wt % TiB2 (-1 at. %) added bGrain ^{s h e} was measured after a standard anneal o f for grain refinement.

1 h at 850°C plus 7 d at 500OC.

compositions. fie chemical analyses (Table ^{2 )} generally agrce well with the nominal compositions, except for PA-33 ^(Fe--30 at. % Al) where the analyzed aluminum level is lower than the nominal one. However, a plot of the densities of all six alloys versus their nominal compositions (Fig. 2) results in a straight line, indicating that alloy ^FA-39 should not be off i t s nominal composition as far as indicated by the chemical analysls. ^The alloys were also analyzed for carbon, oxygen, hydrogen, and nitrogen. ^The

Table 2. Chemical analysis of iron aluminides

Chemical ana 1 ^ys is Nominal cOmpOSitiOn

A1 loy [wt % (-at.

% ) I

_(wt

%I

number

A1 Ti

Fe A 1 Ti Is

...

FA-40 85.69 (75) 13.81 (25) 0.38 0.12 14 0.38 FA-41 84.43 (73) 15.07 (27) 0.38 0.12 15 0.37 FA-39 82.45 (70) 17.05 (30) 0.38 0.12 16 0.36

7

6. 9- 6.

6.

ORNL-DWG 86-1925

-.

8-

7-

h E ^{6 . 6 -}

\ 2

6.5-

+ cn w

H

z 6 . 4 - n

6.

6.

6.

3 - 2 - 1 -

I l - - I I

27 2a 29 ³⁰ 31 ³²

6 _{3 1}

22 23 ²⁴ 2s 26

ALUMINUM (AT. %)

Fig. 2. Composition and density of the iron-aluminum alloys under study. All alloys contain 0.5 wt % TiB2.

analytical results for three of the alloys, presented in Table ^{3 ,} indicate that the concentration of these interstitial elements are very low and should cause no complications in the analysis of our results.

TESTING PROCEDURES Tensile Testing

Tensile samples with a gage section of 0.76 by 3.18 by 12.70 ^mrn were punched from the rolled sheet.

samples were first given a standard heat treatment of 1 h at 850°C plus 7 d at 500°C.

a strain rate of ^{3 . 3} x

room temperature and 800OC.

400'C were cleaned and deburred by either electropolishing or ^vapor blasting. Three of the alloys in the as-rolled condition were selected

for testing at room temperature in air following 30-min vacuum anneals at various temperatures between 550 and 1000°C.

For tensile testing at various temperatures,

All tests were conducted in an Instron testing machine at s - l . Temperatures o f the tests varied between

Samples to be tested at temperatures below

Table 3. Chemical analysis for interstitial elements in iron aluminides

Nominal composition Chemical analysis

A1 loy [wt X _(-at.

%)I

(wt PPm)

number

Fe Al Ti B C O B N

FA-36 86.33 (76) 13.17 (24) 0.38 0.12 50 <1 <1 21 FA-38 85.07 (74) 14.43 (26) 0.38 0.12 46 <1 <1 17 FA-37 83.77 (72) 15.73 (28) 0.38 0.12 56 <1 <1 7

Oxidation Studies

Oxidation studies were performed ^on rectangular samples measuring approximately 10 by 15 mm, cut from the 0.76--mm-thick rolled sheet. Tlie samples were prepared ^{f a r} testing by mechanically polishing with ^4-0 emery paper, followed by annealing in vacuum for 1 h at ^80OOC.

tests were then performed at 600, 800, and 1000°C, each test totaling about 500 h in duration. Measurements of the weight gain as a function of time indicated the degree of oxidation.

Air-oxidation

Sulfidation Studies

Samples of alloys containing 24 to ²⁷ at. % A1 were cut from the rolled sheet in 1- by 1-crn squares ^{f o r} sulfidation studies. Surfaces were sanded with 4-0 emery paper, cleaned, then annealed for 1 h at 80OoC in dry hydrogen.

in platinum foil, sealed in evacuated capsules, and heated at 700OC ^{f o r} 168 h.

gas pressure.

The cleaned samples were embedded in CaSO, powder, wrapped Additional tests were done at 871OC to produce a higher sulfur

Welding Studies

Samples for preliminary welding studies were prepared by mechanically polishing as-rolled material as above, then heat treating in vacuum for 1 h at 8OO0C plus 7 d at ^500OC. Autogeneous electron-beam welds were made on four alloys varying in aluminum content from 24 to 27 at. %. The welds

were made with a 15-kW Leybold-Heraews electron beam-welder.

speed ranged from 25 to 102 c m / m i n at a power level of 75 kV and 5 mA.

The welding

Microstructural Studies

yicrostructural studies were performed on the samples after polishing with 0.5-pm diamond powder (Linde A ) , followed by a chemical etch ^{w i t h 50%}

CK,CQBH, ^33% HNQ;,, ^{and 97% HC1.}

Microprobe Studies

The TiBz precipitates, which were added fox grain refinement and dispersion strengthening, were identified by ^using back-scattered

electrons in a E O L Superprobe. Bath l i n e scxnsrfng element mapping were used to determine compositions of second-phase particles,

Fsactography

The tensile fracture surfaces were examined ^by using secondary

electrons in a Super ^111-A scanning electron microscope from International Scientific Instrument, Inc,

RESULTS AND DISCUSSION

Figure ^{3 shows t h e} microstructures ^{of the} a s - m l l e d and annealed alloy w i t h a. nominal compasiB.ion of Fe-27 at. ?$ A I , Such microstructures are ",pica1

or

the a l l u y s used i n this study. The as-rolled structure

S?'~(ZWS elongated g r a i n s less than 50 plm wide?. A f e w recrystallized grains were observed i n the 25 and 27 at, 41 alfoys a f t e r the last rolling ad;

a temperature of 600°C, taut mPg f i b r a u s grains appeared in the other alloys. The recrystallized microstruztaares, studied after a standard anneal of 1 fi at B ~ O * C followed by ⁷ d at 50 ^{O G ,} sIiowed that g r a i n diame- t e r s f o r all alloys varied between 38 and ^5.5 pm ( s e e Table 1 ) .

Second-phase ^partic1 es presumably t i t a n i u m d i b s r i d e s a were m i - formly distributed ^over the grains an3 grain ZPsundaries

-

Hicroprobe ^and scanning electron mlcroscopa studies indicated ^{t h a t} the p a r t i c l e s were no larger than 1 to ³ pm in diameter, although an in-depth study of the size and ^exact composition of t h e particles remains to be done.

Y

Fig. 3. Optical microstructures of FA-41 alloy (Fe-27 at. X _A1

+

0.5 ut % TiB2).

at 50OoC. ( a ) As rolled. (b) After annealing 1 h at 85OoC plus 7 d

The Fe-27 at. % A1 alloy was studied to determine the recrystalliza- tion temperature (Fig. 4). By annealing €or 30 rnin at various temperatures, it was determined that the alloy started to recrystallize at about 6 5 0 ° C , as shown in Fig. 4 ( b ) .

recrystallized with a grain diameter of approximately ⁴² pm. Grain ^{s i z e} increased very little with increasing temperature above 7 0 0 ° C , remaining at ⁴⁰ to 50 vm in diameter to 1000°C.

After ³⁰ min ^{E a t} 700°C the alloy is totally

These observations indicate that precipitates are effective l".n retarding grain growth in these a1 lays,

F i g u r e 5 ^shows the tensile properties of iron alurainide as functions of aluminum eonc.entration and test temperature.

0 . Z y i e l d stress was highest for the ²⁴ to 26 at. % A 1 alloys (-7.50 MPa) and ~IIPP, decreased rapidly to about 3'30 MPa for the ³⁰ at. X A I alloy,

~ i ~ e same trend w a s ssxen for samples ^:InneaPed at ²⁰⁰ and 400'6, ^altInough

s t r e s s levc?ls were ^lower. f l i t ? s u b s t a n t i a l decrease in yield stress with a11:~iinrirn content from ²⁴ to 27% is p o s s i b l y related to a ^cliange in dis- location s t r ~ ~ c t i i x e from unit ^Ji slocations to superlattice dislocations;

however., detailed ^{transmiss ion} electron microscope studies are needed to v e r i f y the dCsloc.atian structiares i.n these alloys ^{a A t} test temperatures cf dOC, ^700% and 80Q°C, ^{~ h e} opposite trend was s e e n : the lower alumiraim content alloys e x h i h i t e d a slightly lower y i e l d strength, e . g . , 43 MPa for 1~a-25 a t .

x

AI at ~ O O ~ C compared w i t h $8 M P ~ for Fe-30 at. % ^{A B .}

ultimate tensile slxength (?UTS) showed similar trends ^~

A t room temperature the

~fai:

The ductility of iron abumnsnide also depends on aluminurn content and test t@mpeKatUre, ^{The 24} at. % A I a l l o y exhibited a tensile elongation

of abrrraat 1% at roan temperature. The room-temperature el ^oragation

increased steadily w i t h increasing aluminum content and reached about 5%

for the ^3e3 at. % A 1 a l l o y , For samples annealed at ^{200 and} ~ O O ~ C , the durtilil;y increased with aluminum addition up to 26 to ²⁸ at. %.

increase in aluminum to ³⁰ at.

X

^caused no significant improvement.

P b u c t i l i t y increased s h a r p l y above ^400°C; the elongation increased to over 50% at temperatures at and above 60Q°C. Because of the excellent hot duc- tilities, there is no difficulty in fabricating iron aluminide above 600°C.

Further

Fig. 4. Optical microscrucxxues or F A 4 1 alloy (Fe-27 at. X A1

+

(b) Annealed 30 min at 0.5 w t % TiB2).

65OoC. (a) Annealed 30 min at 60OoC.

(c) Annealed 30 min at 70OoC.

13

m I

(D P

s

n 0

1

(VdW) WL3N3b'lS 31ISN3J. 31VWIlln

... ...

I-. --

PI u

SI ?a. '3 D,

LJ LI c E

.. L .... 1 ... l... .... I- -1 rnmr-u3mPm~_,

(X) NOIlV3N[173 1 ... 1 I__.-L ... 0 G 0 0 0 0 0 0 0 t

The alloys with higher ^aluminem content (27, 28, and 30 at. % Al) were further tested by annealing the as.--.rolled alloys ^{f o r 30} min at various temperatures between ^{5 0 0} and 1000°C and then tensile testing at room temperature. The results are presented in Fig. 6. The y i e l d stress decreased continuously with increased annealing temperature until about 7 0 O o C .

remained constant, ^{w i t h} Fe-27 at. % A1 exhibiting the highest strength (477 MFa) and Fe---3Q at. % AI the ^lowest strength ^{(416 MPa). The} trend i.n room--iemperature ductility was for the high@r-alumii1uin alloys (28 and 30 at. % A l ) to exhibit the ^higher elongations. Maximum room-temperature elongations ^{o f 7} to ^97, were reached for the Fe-30 at. % A1 alloy at ^tean- peratures ^{of 625} to ^713OOC.

From ^{7 0 0} to 1000°C the strength of each alloy in this series

The tensile properties ^{o f} Fe-28 and 30 a;. 7; _{A1 alloys as a Eunct:ion} of test temperature are compared with ^the proprtties ^{o f} type ³¹⁶ stainless steel and modified ^91Zr-11Mo s L c e 1 in ^{F i g .} 7. It can be seen that the yield strength ^of the iron aluiiminides i s better than that of type ³¹⁶ stainless s t e e l up to 7 6 0 ° C and better thain that ^of modified 9C;r-lMo steel ^above 5.SO°C. However, the ducti~liiy at temperatures below ^400°C needs to be improved.

E’RACTOGH APHY

Scanning electron h,iicroscopy was used to examine the fracture ^{s u r -} faces of samples that had been tested at room temperature. ^A previous study showed ⁴ hat thp fracture mode in the ^as-i o l l e d matPLial is mainly transgranulsr while the annealed ^samples fail intergranularly.*

r c s u l t s were substantiated by r e s u l t 5 of ^tlrc present s t u d y . Tensile samples that had ^bsen heat ^tiealetl before testing ^{(1 €1} at 8 5 0 ° C plus 7 ^d a% ^500°C) exhibited a fracture node typical of ^inairily intersgrani~lar failure [Fig. 8 ( a ) ] , with approximately 10 to ^20% Lransgranular character.

At ^higher magnification, [Fig. B ( b ) ] , cracking along grain boundaries is apparelit, and TiB2 particles can be seen ^{O N} the g r a i n surfaces, as well ^as the few transgranular fracture regions.

Those

1000 900- 800-

2 ^’700- 5 Q I 600 c 0

5

sm-

F

u) 0 J W

400-

;I ^300-

..._..

... .-I_____.

-

ORNL-DWG a6-1930

45 40 35-

fs ^{A 30-}

z 0

<

u z

-1 W

r-( t- 25-

a 20- 15-

AT.% AL

* ²⁷

+ 28

X 30

- -

“OI

¹⁰⁰

500 575 650 725 800 875 950 1025 1100 1175 1250

ANNEALING TEMPERATURE P O

AT.% AL

* ²⁷

+ 28

x ³⁰

I I I ₁ I I I I

ANNEALING TEMPERATURE (‘C)

Fig. 6. ^Effect of 30-min anneals at various temperatures on tensile properties o f iron aluminides.

ORNL- DYJG 86-1932

1

____ ~~

I

g L - - 1- - L I - L - 1 - 1 I - L

0 10C 20C ³⁰⁰ 400 ^3UO 600 700 800

TFST TEMPEQ4TURt {‘CJ

1-

9 c o

9ci-

^/ /’

- x- -/- x i

A L

Fig.. 7 . E f f ~ ~ t of test temperature ^{Q H ~} t e n s i l e properties sf ^fe-28

A I , and comparison of tensile properties with those of ^type 316 stainless s t e e l and modified 9Cr-lMs steel.

Fig. 8. Fractographs of FA-39 alloy (Fe-30 at. % ^{A 1}

+

^{0 . 5} wt 4: _TiB2) after annealing 1 h at 85OoC plus 7 d at ^5OO0C and tensile testing at room temperature.

OXIDATION STUDIES

Oxidation studies ^were conducted on the ^YE--24, 25, 26, ^27, and

30 at.

X

^AI alloys at 600, 800, and 1O0OoC, each test lasting approximately 500 h .

f o r all alloys tested, and oxide films with colors in the interference range were observed. ^Thr. resijlts at 800OC are presented in ^{F i g . 9} for the Fe-24, 2 7 > and ³⁰ at. % A1 alloys, along ^{w i t h} the data for.type ³¹⁶

stainless s t p e l al.. t h i s tcmperature. After approximately ^{5 0 0 h} at this temperatuie, all iron aluminide alloys had a dull b l u i s h gray color, ^{w i t h}

no appareni spalling, and ^weighL gains of less than 5 ^x lo-' g/cm2. The type 316 stainless ^{st-eel, on} the ^other hand, gained about 12 ^x

ak 120 h, when it began to spall. Hesults ^{o f} the tests at 1000°C are shirwn j n Table ^{4 .} Weight gains ^601- the iron aliiminides ^remained low

At 500°C, weight ^gairis of less than ^{0 . 3 x} g/cm2 were recorded

g/cm2

ORNL-DWG 86-1933

A L L O Y

* 2 4 % Al.

+ 2 7 % AL.

X 3 0 % AL 0 3 1 6 SS

TIME AT TEMPERATURE (H)

Fig. ^{9 .} Oxidation resistance of iron aluminides ^and of type ³¹⁶ stainless steel at ^80OoC.

Table 4 . O x i d a t i m r e s i s t a n c e of iron aluminidgs a t 1000°C

+

Weight change a f t e r 526-h exposure Composition

( a t . X A l )

( g h 2

1

Alloy

FA-36 24

FA-40 25

FA-38 ²⁶

FA-4 I ²⁷

28-39 ³⁴⁾

31.6

ss

4 . 4 9 x 10-4 5 . 4 0 x 10-4

4.48 x

lo-&

5.38 x IO--@

4.74 ^x

-1516.78 x

lo-"

€2

Table 5 . Corrosion o f irosr aluminides i n c a p s u l e test"

Composition Weight gain ( a t . % Al) ( ^[ll& / ^c m 2 ) Alloy

FA- 3 6 EA-40 FA-38 FA-41

FA- 3 6 FA-40 FA-38 FA-41

2 fe

2.5 26 27

24 2 5 26 27

0 . 0 5 A l l a l l o y s showed 0 . 0 3 i n t e r f e r e n c e 0.04

0.02 C O l O K S .

0.24 A11 alloys ^were 0.27 covered w i t h _a 0.22 d u l l gray c o a t i n g 0.25

"Exposure for 168 h ^{i a a} s c a l e d , evacuated q u a r t z c a p s u l e to t h e gaseous decomposition p r o d u c t s o f CaSOb.

Samples were not pre-oxidized.

WELDARILITY

P r e l i m i n a r y w e l d a b i l i t y r e s u l t s from electron-heam welding and micro- s t r u c t u r a l a n a l y s i s performed on Fe-24 t c ~ 27 a t . % A 1 a l l o y s indicated t h a t the a l l o y with highest aluminiim content (Fe-27 ^{a t .} % AI) had t h e b e s t w e l d a b i l i l y , with ^no cracks ( F i g . 10). The ^{welds of} sther a l l o y s showed t r a n s v e r s e or crater cracks and sometimes both. The r e s u l t s are ^sum- marized i n Table 6 "

Table 6 . Electron-beam w e l d a b i l i . t y of iron alutuinides

Composition ( a t . % Al)

Alloy W e l d a b i l i t y

._.._.

FA-36 ²⁴ Transverse c r a c k s FA-40 25 Transverse and c r a t e r

c r a c k s

FA-38 26 Crater c r a c k s

FA-41 ²⁷ No cracks

Y-205228

I

,

, ^Warn ,

Fig. 10. Section of electron-beam weld in annealed Fe-27 at. % Al.

The top view of the Fe-27%

AI

weldment in Fig. 10 shows the epitaxial growth of grain structure in the fusion zone, starting from the HAZ.

elongated grains grow and impinge against each other along the fusion line.

effective in pinning grain boundaries even at temperatures to the melting point.

The The

HAZ

is not well defined, indicating that the TiBi particles are

CONCLUSIONS ^AMD FUTURE WORK

The present work indicates that alloys of Fe-24 to 30 at. % A1 with small additions of TiBz are easily fabricated and exhibit excellent oxidation and corrosion properties.

dizing environments is conferred on these alloys by a self-protecting oxide layer that forms at low oxygen pressures.

are higher than those for type 316 stainless steel at temperatures below Resistance to oxidizing and sulfi-

The tensile strengths

T60°C and for modified 9Cr-lMo steel at temperatures above ^550OC.

Fe-A1 alloys tested, the room-temperature ductility is slightly higher ^{f o r} those containing more Lhan 27 a L . % Al. Our preliminary weldability stud- ies using electron-beam processes indicate that the alloys with higher aluminum content have fewer cracks in the fusion zone and ^IIAZ. Based on considerations of strength, corrosion resistance, fabricability, and weldability, we have selected the ^Fe-28 and 30 at. ^{A 1} alloys ^as base alloy compositions for further alloy development.

Of the

Planned f u t u r e work includes alloying with molybdenum, titanium, and zirconiiim ta improve the high-temperature strength and roam-temperature ductility. Preliminary studies involving molybdeniim indicate that

molybdenum-containing precipitates form in the matrix, causing a reduction in grain size. Grain dianieters in ^{c a s t} ingots were siibsLantially reduced (Prom 105 to 25 ~ i m ) by addition of only 2 at. Mo. Tests io determine the solubility limit of molybdenum i n Fe-28 at. % A I are in progress, along with the fabrication and preparation of specimens for mechanical testing. Future plans also involve further mechanical testing, including creep studies, and further corrosion, oxidation, and welding studies.

ElecLron microscopy s t u d i e s well be conducted to determine precipitate composit.iun and morphology. Studies ^of the dislocation structures, antiphase domain StrucLiires, and ordering processes and kinetics as a function of alloy additions will ^{a l s o} begin.

REFERENCES

1. ^{E . R .} Slaughter and S . K. Das, "Iron-Aluminum Alloys with Titanium Dibori de Dispersions ^by Rapid Solidification,

''

pp. 354-63 in Praceedii~gs of ^{t h P} Second Iii-temat i o n a l Conference on R . ~ ~ s i d Solidification Processing, ed. R. Mehrabian, ^tl. H. ^Kear, and K. ^Cohon, Claitor's Publishing Division, Baton Rouge, La., ^1980.

2. II. Znouye, C. T. L i u , and ^{3 .} A . Rorton, P h y s i c a l MeLfallurgy and Mechanical P r o p e r t i e s of Iron Aluminides, final report (unpublished) for Office of Naval Research, D e p a r t w n t of the ^Navy, Arlington, Va., ^by Oak Ridge National I,aboratory, for work from November 1982 to September 1984.

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( 1 9 7 8 ) .

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,

liigh-TemperatLrr-e Ordered l a t c r m e t a l l d c Afloys, ed.

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P i t tsburgh, 1985.

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of S ~ ~ m d a r d , s Symposium on Mapid Solida'hisat ion Prni*esh;issp: Princ ip1es and Technil4ogies I V , DPC. ^{6-8, 1382,} Gaithcrshurg, Md.

,

National Wiare,iu o f Standards, Gaithersburg, M d . 1983.

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(1960) ⁹

8 . K . Oki, M. Kasaka, ^and T. ^{EgiichF, Jap. J .} A p p l . Phys. _12(28), 1522 (1973).

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11. S . M. Allen, P h i l . Ndg. 34(1), 1 8 1 (1977).

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WADC

Technical Report ^52-298, Wright Air Development Center, Wright-Patterson AFB, Ohio, May 1957.

13. J. W, Holladay, '*Review of Development in Iron-Aluminum-Base Alloys," DMIC ^{Memo 8 2 ,} Defense Metals Information Center, Battelle Memorial Institute, Columbus, Ohio, Jan. 30, ^1961.

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15. M. ^G. Mendiratta and H. ^{A .} Lipsitt, "DOJ-Domain Structures in Fe,Al-X Alloys

,"

pp. ^155-62 in Materials Research Society Symposia

Proceedings, vol. ^{3 9 ,} High-Temperatore Ordered I n t e r m e t a l l i c Alloys, ed.

C. C. Koch, C. T. ^{L i u ,} and N. S. S t o l o f f , Materials Research Society, Pittsburgh, ^1985.

16. Charles D. Hodgman, ed., ^Kandbook of Chemistry and Physics, 40th E d i t i o n , Chemical Rubber Publishing Co., Cleveland, 1958, p. ^3072.

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