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ELSEVIER

Intermetallics 4 (1996) 351-315

0966-9795(95)00056-9

Published by Elsevier Science Limited Printed in Great Britain. All rights reserved

0966s9795196/$15.00

Nickel and iron aluminides: an overview on properties, processing, and applications

S. C. Deevi* & V. K. Sikka

Metals and Ceramics Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, USA (Received 19 October 1995; accepted 8 November 1995)

Intermetallics have long been recognized as potential candidates for a variety of high-temperature structural applications to operate well beyond the operating temperatures of conventional materials due to their excellent oxidation and cor- rosion resistances. In this paper, we compare and contrast the mechanical properties such as yield strength, ultimate tensile strength, and tensile elongations of Ni,Al-based alloys, Fe,Al-based alloys, and FeAl alloys with several of the com- mercially available superalloys such as Haynes 214 (NiCrAlY), MA-956 (yttria- dispersed FeCrAlY), and a FeNiCr alloy (HU steel) used in carburizing applications. Our comparisons clearly show that cast and wrought N&Al-based alloys exhibit superior mechanical properties over the commercially available alloys such as the FeNiCr HU steel and Haynes 214. Electrical resistivity of iron aluminides increases with the increase of aluminum content, and the electrical resistivities of Fe,Al- and FeAl-based alloys are 50-100% higher than those of commercially available heating-element materials. Processing problems associated with the melting and casting of intermetallics are discussed in light of their large, negative heats of formation; high-aluminum content of intermetallics; and the safe operating temperatures of crucible materials for melting them. A fumace- loading sequence enabled us to properly utilize the heat of reaction of inter- metallics resulting in the development of the Exo-MeltTM process for melting and casting of intermetallics for a variety of structural applications. The Exo- MeltTM process allowed us to cast a wide variety of structural intermetallic parts using sand, centrifugal, and investment casting techniques, and a total of 15 000 kg of intermetallic parts were cast by the Exo-MeltTM process during 1995.

Copyright 0 1996 Elsevier Science Ltd

Key words: A. nickel aluminides, based on Ni,Al, iron aluminides (based on Fe,Al), iron aluminides (based on FeAl). C. casting, G. automotive uses, corro- sion and erosion resistant applications,

1 INTRODUCTION

During the last several decades, ordered inter- metallics, based on aluminides of transition metals such as nickel, iron, titanium, niobium, and cobalt, have been under investigation for their possible use as high-temperature structural materi- als. I4 Aluminides of transition metals possess sufficiently high concentrations of aluminum (see Table 1) to form a continuous, fully adherent alu- mina layer on the surface when exposed to air or oxygen atmospheres. The amount of aluminum present in aluminides can range from 10 to 30

*Dr S. C. Deevi is on sabbatical at the Oak Ridge National Laboratory from Philip Morris USA under the Philip Morris Fellowship Program. His permanent address is: Research and Development Center, Philip Morris, USA, Richmond, Virginia 23234, USA.

357

wt% and is significantly higher than the aluminum concentrations present in conventional alloys and superalloys. In the case of nickel and iron alu- minides, the alumina layer formed on the surface of the materials is responsible for their excellent oxidation and carburization resistances even at temperatures as high as 1000°C or higher. There- fore, aluminides, unlike conventional steels and superalloys based on nickel, iron, and cobalt, do not necessarily require chromium to form an oxide layer on the surface of the material to protect against high-temperature oxidation and corrosion.

Alumina is much more thermodynamically stable at high temperatures than Cr,O,. Interestingly, the chemistry of aluminides is much simpler than superalloys; subsequently, they form long-range- ordered crystal structures. Apart from their oxidation and carburization resistances, aluminides possess

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358 S. C. Deevi, V. K. Sikka

Table 1. Weight percentages of aluminum, heats of formation, and melting points of intermetallics

Intermetallic

N&Al NiAl N&Al, NiAI,

Fe,AI FeAl FeAl, Fe,Al,

Weight percent (wt%) of aluminum

13.28 31.49 40.81 57.96

13.87 32.57 49.1 54.70

Heat of formation’ Melting point

A~~,ps(kcal/mol) (“C)

-66.6 f 1.2 1395

-28.3 + 1.2 1639

-67.5 f 4.0 1133

-36.0 f 2.0 854

-16.0 1502

-12.0 1215

-18.9 1164

-34.3 1171

“From selected values of the Thermodynamic Properties of Binary Alloys, eds R. Hultgren, P. D. Desai, D. T. Hawkins, M.

Gleiser & K. K. Kelley, American Society of Metals, Materials Park, Ohio, 1973; and 0. Kubachewski and C. B. Alcock, Metal- lurgical Thermochemistry, Student Edition, Pergamon Press, New York, 1979.

lower densities, high-melting points, and exhibit interesting mechanical properties due to their ordered crystal structures. The strength of some intermetallics increases with temperature instead of exhibiting a decrease; thus, they are ideally suited for high-temperature applications. Advantages of intermetallics based on N&Al, Fe,Al, and FeAl are given below.

The major advantages that can be derived from the use of nickel aluminides (N&Al) include:

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Resistance to oxidation and carburization in both oxidizing and reducing carburizing atmos- pheres up to 1100°C.

Good tensile and compressive yield strengths at 650-l lOO”C, as opposed to many nickel- based superalloys.

Fatigue resistance superior to that of nickel- based superalloys, resulting from the elimi- nation of second-phase particles such as carbides.

Superior creep strength, which is highly grain- size dependent. For the best creep resistance, coarse grain size is required; thus, castings need to be considered.

Excellent wear resistance at high temperatures (2 600°C). In fact, wear resistance increases with temperature (similar to yield strength). The increase can be as much as a factor of 1000.

The formation of Al,O, on the surface by preoxidation provides good chemical com- patibility for many environments.

The major advantages that can be derived from the use of iron aluminides include:

(1) Their density is lower than that of many stainless steels and, therefore, they offer better strength-to-weight ratio.

(2) Their resistance to sulfidation in H,S and

SO2 gases is much better than that of any other iron- or nickel-based alloys.

(3) They have a high electrical resistivity which increases with temperature.

(4) They have good corrosion resistance in many aqueous environments.

In spite of the above attractive features, low ductility, brittle fracture, and processing problems have hampered the commercial utilization of these intermetallics for several decades. Several reviews, symposium proceedings, and books describe the recent advances in alloy design, processing, and mechanical properties of intermetallics.‘-6 Excellent accounts of historical perspectives on the develop- ment of intermetallics are provided by Cahn7 and Westbrook. In this article, we primarily focus on the mechanical properties, processing, and appli- cations of nickel and iron aluminides developed at the Oak Ridge National Laboratory (ORNL) over a period of a decade or longer.

2 N&Al-BASED ALLOYS

A concerted effort by several researchers to enhance the ductility and alter the mode of fracture from brittle intergranular to ductile transgranular led to the understanding that moisture-induced hydrogen embrittlement is primarily responsible for low ductility and brittle fracture. Alloying of aluminides was considered to alleviate the above issues, and a breakthrough occurred when N&AI was micro- alloyed with boron by Aoki and Izumi’ in Japan.

Subsequent intense studies of boron and alloy stoichiometric effects by Liu and his co-workers’@‘2 at ORNL has shown that 40-50% tensile ductility can be achieved in N&Al with as little boron as 0.04 wt%. Interestingly, the addition of 0.5 at% B

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Nickel and iron aluminides: an overview 359

resulted in a room-temperature tensile ductility of 50% at 24.0 at% Al, and the ductility dropped sig- nificantly to about 5% when the aluminum content was increased to 25 at% or higher. Although several other elements, such as rhenium, zirconium, and hafnium, were observed to enhance the room-tem- perature ductility of Ni,A1,‘0~‘2 none were as effec- tive as boron. Even though the role of alloying elements in enhancing the room-temperature duc- tility is not clearly understood, the addition of boron even in small quantities is believed toincrease the cohesive energy of the grain boundaries along with suppression of environmental embrittlement.‘3~‘5 Recent observations by Deng et all6 suggest that the boron doping is beneficial because of the strengthening of bonding cohesion in grain boun- daries and suppression of hydrogen embrittlement.

Their observations were based on results obtained by the positron annihilation technique, which is sensi- tive to defects and their interaction with hydrogen.

Improved ductility and transgranular fracture of boron-doped Ni,Al allowed design of several Ni,Al- based alloys of compositions shown in Table 2 for high-temperature applications.‘3,17 It is interesting to note that the compositions of most of the alloys contain several different elements other than nickel, aluminum, and boron, The addition of alloying elements such as zirconium and chromium were found to be necessary to reduce solidification shrinkage, increase high-temperature strength and ductility, adhere the oxide film on oxidation, and confer the hot-corrosion resistance to the Ni,Al phase. An extensive evaluation of the mechanical properties of cast and wrought alloys was carried

out at ORNL to determine their suitability for commercial applications. Commercial use of Ni,Al- based alloys for high-temperature applications will be based on their overall superior mechanical, ther- mal, and physical properties over the existing materials, and the actual use of Ni,Al-based alloys in commercial applications will be a replacement of existing materials with N&Al-based alloys. Since industries generally operate by replacing existing materials with advanced materials, a comparison of mechanical properties of cast, wrought, and powder metallurgically processed alloys is of interest for engineering applications. Mechanical properties of the alloys developed at ORNL were used for com- parison with the commercially available materials.

To illustrate the strength advantages of IC-221M and IC-396 Ni,Al-based alloys over a commer- cially available FeNiCr alloy (known as HU) for high-temperature furnace applications,‘* yield and ultimate tensile strengths of IC-221M, IC-396, and HU are plotted against temperature as shown in Figs l(a) and (b). The yield strengths of IC-221M and IC-396 alloys are at least twice that of HU at room temperature and increase with rising tem- perature as shown in Fig. l(a). The strength of HU decreases with rising temperature while the strength of IC-221 M and IC-396 N&Al-based alloys at 800°C is over 200% higher than that of HU, which is commonly used in a variety of manu- facturing operations as heat-treating trays. The strength advantage of cast Ni,Al-based IC-221M and IC-396 alloys can be seen even at a temperature of 1100°C. The ultimate tensile strengths of Ni,Al- based IC-221M and IC-396 alloys are significantly

Table 2. Compositions of N&Al-based alloys selected for commercial applications and compositions of some commercial alloys

Element Alloy (wt%)

IC-221M” IC-396’ HU’ IC-50” IC-2 18LZr’ IC-221w Haynes 214’ Alloy 8Od

AI Cr MO Zr B C Fe Ti Ni Si Y

8.0 7.98

7.7 7.72 18.0

1.43 3.02

1.7 0.85

0.008 0.005

0.55 42.45

81.1 80.42 39.0

11.3 8.7 8.00 4.5 0.4

8.1 7.70 16.0 21.0

1.50

0.6 0.2 3.00 -

0.02 0.02 0.003 -

- 0.03 0.05

3.0 45.5

- 0.4

88.08 83.1 79.80 76.35 32.5

0.1

- 0.02

“Castable alloy for dynamic applications (minimum microporosity).

‘Castable alloy for static applications (some microporosity).

‘Cast alloy.

“Cold workable.

‘Hot and cold workable.

‘Wrought alloy.

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360 S. C. Deevi, V. K. Sikka

- IC-396

1000 1200 Temperature (“C)

(4 g. 1250

I

.= ii 0 200 400 600 800 1000 1200

Temperature (“C) (b)

0 200 400 600 800 1000 1200

Temperature (“C) cc)

Fig. 1. A comparison of (a) yield strengths, (b) ultimate tensile strengths, and (c) tensile elongations of cast N&Al-based

alloys with HU alloy.

higher than the yield strengths shown in Fig. 1 (b), and the ultimate tensile strengths of N&Al-based IC-221M and IC-396 alloys are higher than the strength exhibited by the HU alloy. Tensile elon- gations of N&Al-based IC-221M and IC-396 alloys shown in Fig. l(c) indicate excellent room-temper- ature tensile ductility exceeding that of HU. The excellent oxidation resistance and the superior me- chanical properties of Ni,Al-based alloys suggest that IC-221 M or IC-396 can replace the FeNiCr (HU) alloy for furnace applications.

A commercially available nickel-based superal-

loy sold as Haynes 21419 by Haynes International, Inc. (Kokomo, Indiana) is compared with the tensile properties of wrought N&Al-based alloys up to 1100°C (see Fig. 2). Yield strengths of Ni,Al- based IC-50 and IC-218LZr match the yield strength of Haynes 214. On the other hand, powder- metallurgically processed IC-221 W exhibits a significantly higher strength than those of wrought alloys, even at 700°C (see Fig. 2(a)). Ultimate tensile strengths of N&Al-based alloys are significantly higher at room temperature as in the case of cast alloys, and the strength advantage of Ni,Al-based alloys over Haynes 214 is lost at or above 850°C

1500 E I 1250

5 P

1000 iz e 750

3 e 500

ifi 250

d 0

+ Haynes-214 - IC-so -)- ICI-2 1 BLZr + K-221 W (PM)

0 200 400 600 800 1000 1200

Temperature (“C)

1000 1200 Temperature (“C)

(b)

0 200 400 600 800 1000 1200

Temperature (“C) (c)

Fig. 2. A comparison of (a) yield strengths, (b) ultimate tensile strengths, and (c) tensile elongations of wrought Ni,Al-based

alloys with Haynes 214.

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Nickel and iron aluminides: an overview 361

(see Fig. 2(b)). During the intermediate temperatures (i.e. 4008OO”C), IC-50 alloy with a simple compo- sition consisting of nickel, 11.3 wt% Al, 0.2 wt%

B, and 0.6 wt% Zr is weaker than Haynes 214, and its ultimate tensile strength drops almost linearly with temperature. In contrast, IC-221 W processed by powder metallurgy exhibits significantly higher strength than Haynes 214 and other N&Al-based alloys and retains its room-temperature strength until 650°C. The strength drops above 650°C and the strength of powder-metallurgically processed IC-22 1 W is similar to the other alloys at or above 850°C (see Fig. 2(b)). Tensile elongations of 40%

or higher suggest that room-temperature process- ing of Ni,Al-based wrought alloys can be carried out similarly to Haynes 214 (see Fig. 2(c)). As expected, the tensile elongation of powder-metal- lurgically processed IC-22 1 W is much lower than that of wrought alloys at room temperature, and the strength remains constant over a wide temper- ature range. Low ductility of powder-metallurgically processed IC-221W requires a temperature of 950°C for thermomechanical processing. It is important to note that the composition of IC-50 is much simpler and is close to the composition of Ni,Al

loo0

d

i_;-;,

IO a. .‘.. .‘.. a.. .a.

30 35 40 45 50 55 60

P=(T+460)*(20+logtr)*1 O-3 (a)

000

100

IO

1 1

35 40 45 50 55

P=(T+460)*(20+logtr)*1 O-’

(b)

Fig. 3. A comparison of creep resistances of (a) cast N&Al- based alloys with commercial HU alloy and (b) wrought

N&Al-based alloys with Haynes 214.

intermetallic. Our results also indicate that the fatigue properties of IC-221 M are superior by an order of magnitude to the commercially available IN-713C (sold by Into Alloys in Huntington, West Virginia)*’ when tested in air at 650°C.

In the operating temperature range of 800- lOOO”C, creep resistance of cast and wrought alloys is of importance since creep is the primary defor- mation mechanism at high temperatures. Creep data of cast alloys are compared with HU in Fig. 3(a) while the creep data of wrought alloys are compared with Haynes 214 in Fig. 3(b). The Larsen-Miller plots shown in Fig. 3(a) indicate the superior creep strength of IC-221M by a factor of 2-4 even at high temperatures. The creep strengths of wrought alloys are also significantly better than the creep strength of Haynes 214. The high-aluminum content of nickel aluminides allows the formation of pro- tective alumina layer against oxidation and car- burization at high temperatures. The oxidation resistances of nickel aluminides in air with 5%

water vapor at 1100°C are far superior to that of Alloy 800 with no significant weight change even after 250 d exposure (see Fig. 4(a)). On the other

50

“E 4 OL 0

z _ -50 :..-.. .

+ K-50

--f K-21 8

- K-22 1

..;. . . . .._

Air-S% Water Vapor at 1100°C

-200 “’ 8” I.‘..‘....1 .‘I ..

0 100 200 300

Exposure Time (Days) (a)

Hz-5.5% CO, at 1000°C

~ ,c_50 -t-K-21 8

- K-22 1

10 20 30

Exposure Time (Days) (W

Fig. 4. A comparison of (a) oxidation resistance (in air with 5% water vapor at 1lOO’C) and (b) carburization resistance (in H,-5.5% CO,) of N&Al-based alloys with commercially available Alloy 800. (Data were developed at The Interna-

tional Nickel Company, New York).

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362 S. C, Deevi, V. K. Sikka

hand, Alloy 800 consisting of 21 wt% Cr and only 0.4 wt% Al lost over 150 mg/cm2 within 40 d exposure in air with 5% water vapor at 1100°C.

Nickel aluminides also exhibit excellent carburiza- tion resistance as compared to Alloy 800 when tested in a carburizing atmosphere consisting of hydrogen with 5.5% CO, at 1000°C (see Fig. 4(b)).

While nickel aluminides were resistant to carburiza- tion attack, Alloy 800 exhibited a significant weight gain because of carburization of the material.

It is interesting to note that IC-50 with a higher aluminum content was much more resistant to carburization than IC-221 M and IC-218.

Gas-atomized powder of IC-50 was also used to obtain composites of WC-N&Al for use as tool bits.

In this particular application, ORNL researchers2’

exploited the unique properties of Ni,Al to replace cobalt as a binder phase in the tungsten-carbide- cobalt tool bits. Figure 5 shows the flexural strengths of several composites of WC-Ni,Al at room temperature and at 800°C. The strength of the Ni,Al composites changes very little up to 800°C and the strength of WC-17 ~01% Ni,Al is

2400 k

E

3 800 1

I , IVI~AI 17co

_L

q 25°C

1 m 800°C

68 Ni3AI 80 Ni3AI 95 Ni3AI Fig. 5. Elevated-temperature flexural strength for tungsten- carbide composites containing various N&Al contents, (courtesy of T. N. Tiegs, Oak Ridge National Laboratory, Oak Ridge,

Tennessee).

Table 3. Mechanical properties of WC and TiC composites with varying N&AI contents (IC-50)

Sample no.

IBC-56 32% WC 25.7 1751

IBC-57 20% WC 18.5 1092

IBC-58 5% WC 18.4 1123

IBC-59 60%TiC 2.1 131

IBC-60 43%TiC 3.2 95

IBC-6 1 14%TiC 14.1 782

IBC-62 0 b b

Ceramic phase content (vol”?,)

Fracture Flexural toughness strength (MPa Jm) (MPa)

Hardness (GPa)

7.0 4.6 4.0 9.9 6.0 4.2 2.3

“P. Chantikul, G. R. Anstis, B. R. Lawn & D. B. Marshall, J.

Am. Ceram. Sot., 64(9) 53943.

‘Specimens bent during flexural test.

indeed higher due to the increase of yield strength of N&Al intermetallic with temperature. On the other hand, the yield strength of cobalt decreases with temperature, and therefore, the flexural strength of WC-17 ~01% Co exhibited a significant decrease at 800°C. Flexural strengths, fracture toughnesses, and hardnesses of tungsten-carbide and titanium- carbide composites with IC-50 binder are shown in Table 3. Composites of tungsten-carbide and tita- nium-carbide with IC-50 as a binder phase pro- vide the required toughness and strength with the added advantage of corrosion resistance.

3 Fe,Al- AND FeAl-BASED ALLOYS

Iron aluminides were recognized to possess excellent oxidation and sulfidation resistance even during the early 1930s but the lack of room-temperature ductility, low high-temperature strength, and inad- equate high-temperature creep resistance left them unexploited for commercial applications.22m26 Iron aluminides of interest are Fe,Al and FeAl with ordered body-centered-cubic structures correspond- ing to DO3 and B2 crystal structures, respectively.

In the case of Fe,Al, a DO, structure is stable in the 23-36 at% Al range and from room temperature to 550°C. Above 550°C an ordered Fe,AI with DO3 structure transforms to an imperfectly ordered B2 structure, which ultimately changes to a dis- ordered solid solution, (Y. On the other hand, FeAl exists with B2 structure and is stable from about 36-48 at% Al, and the transition from B2 to (Y occurs well above 1100°C

During the last decade, efforts were focused to enhance room-temperature ductility, high-temper- ature strength, and high-temperature creep resistance by alloying of iron aluminides. Two approaches, namely, solid-solution strengthening and precipita- tion strengthening, were considered for strengthening of iron aluminides. Elements such as Nb, Cu, Ta, Zr, B, and C were considered for precipitation strengthening; Cr, Ti, Mn, Si, MO, V, and Ni were added for solid-solution strengthening. In general, the addition of elements either for precipitation strengthening or solid-solution strengthening to improve high-temperature tensile strength and creep resistance resulted in low room-temperature tensile elongations. 27 In some cases, solid-solution strengthening also increased the DO,-B2 transi- tion temperature. The addition of molybdenum when added along with zirconium resulted in improved tensile and creep strengths, and only chromium was observed to increase the room-

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Nickel and iron aluminides: an overview 363 temperature ductility of Fe,Al. The addition of 2

at% Cr to Fe,Al resulted in some intergranular fracture along with transgranular fracture. Poor ductility of iron aluminides, Fe,Al, and FeAl is attributed to the environmental embrittlement due to the interaction of aluminum with water vapor present in the atmosphere as shown below:28

2 Al + 3 H,O + Al2O3 + 6 H (1) The atomic hydrogen generated in the above reaction is believed to cause hydrogen embrittle- ment in atmospheres containing water vapor, and tensile elongations of Fe,Al and FeAl improved significantly when tested either in vacuum or oxy- gen.28.29 However, practical applications exploiting the oxidation, carburization, and sulfidation resis- tances of the iron aluminides generally involve water vapor (large quantities in some cases) and may even have atomic hydrogen. Design engineers are, therefore, restricted to the strengths and tensile elongations of iron aluminides measured in air.

Extensive alloying efforts to optimize properties such as creep resistance; room- and high-temperature strengths; and oxidation, carburization, and sulfi- dation resistances, while maintaining reasonably good room-temperature ductility, resulted in several Fe,Al alloys containing 2-5 at% Cr.27.30,3’ Compo- sitions of Fe,Al alloys, along with a commercially available mechanically alloyed FeCrAlY with yttria dispersion (sold by Into Alloys (Huntington, West Virginia) as MA-956),*’ are shown in Table 4. It is clear from Fig. 6 that yield and ultimate tensile strengths of wrought Fe,Al-based alloys are signifi- cantly higher at room temperature and retain their

strengths up to 600°C as compared to the MA-956 superalloy. The strengths are similar at 600°C and are very low at 800°C for their use as structural materials. Tensile elongations of wrought Fe,Al alloys increase with temperature while the temper- ature has a minimal effect on the tensile elonga- tion of MA-956 (see Fig. 6(c)). Creep resistance of MA-956 is superior to that of iron aluminides because of the dispersion of yttria by mechanical alloying (see Fig. 7). It is to be recognized that the cost of mechanically alloyed MA-956 is higher by a factor of 30 as opposed to the Fe,Al-based alloys discussed in this paper. Our preliminary results suggest that yttria dispersion can improve the creep properties of Fe,Al alloys, and the cost of mechanical alloying will also increase the price of Fe,Al alloys.

The high-aluminum content of iron aluminides allows the formation of a protective alumina layer that imparts sulfidation resistance to iron alu- minides. Therefore, iron aluminides are particu- larly suitable in harsh, aggressive, and corrosive environments up to 800°C as compared to FeCrAl and FeCrNi alloys (see Fig. 8). Also, Fe,Al alloys could easily be hot worked at or above 800°C and further processing such as rolling can be done in the temperature range of 600&8OO”C.

Aluminides based on FeAl exhibit better oxida- tion and corrosion resistance than Fe,Al alloys and are also lower in density by as much as 30-40”/;1 as compared to steels and other commer- cial iron-based alloys. Mechanical properties of binary FeAl are dependent on the aluminum con- tent as can be seen in Fig. 9, which shows the

Table 4. Compositions of Fe,Al-based alloys developed for commercial applications and compositions of some commercial alloys

Element

FAS” FALh FA- 129’

Alloy (wt’%)

FAH Fecralloy MA-956 Type 310

Al Cr I3 Zr Nb C MO Si Y y703 T;

Ni Fe

15.9 15.9 15.9

2.20 5.5 5.5

0.01 0.01

0.15

- I.0

- 0.05

-

d

- -

d d

15.9 4.5 -

5.5 16.0 20.0 25.0

0.04 0.15 I.04

- 0.02 0.15

1.00 -

- - 0.50

0.3 -

0.5

d

- d

0.5

d 20.0

d

“Sulfidation-resistant alloy.

hHigh room-temperature (RT) tensile ductility.

‘High-temperature strength with good RT ductility.

‘Balance.

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364 s. c. Deevi, V. K. Sikka

3 800

s

= 600 F x e 400

%

‘F 200 ZR (u

d 0

c 2

1200

r, 1000 c”

?! 800 x

Q 600

= I-” : 400

200

g 150 6

‘G zl : 100

5 z z 50 I-

O

0 200 400 600 800 1000 1200

Temperature (“C)

(a)

Temperature (“C)

@I

0 200 400 600 800 1000 1200

Temperature (“C) (c)

Fig. 6. A comparison of (a) yield strengths, (b) ultimate tensile strengths, and (c) tensile elongations of Fe,Al-based alloys

with MA-956 superalloy.

variation of yield strengths and tensile elongations of binary FeAl alloys. While the yield strengths increase with the increase of aluminum content up to 40 at0/o Al, room-temperature tensile elongations of aluminides decrease with the increase of aluminum content. Also, the yield strengths of FeAl are insensitive to the aluminum content at 700°C and are too low for structural applications at or above 700°C. Low ductility of iron aluminides necessitates hot working above 700°C and in practice, a tem- perature of 900°C or higher is employed. While

m--t FA-129

25 30 35 40 45 50 55 60

P=(T+460)*(20+logtr)*1 O-3

Fig. 7. A comparison of creep resistance of Fe,Al-based alloys with MA-956 superalloy.

00

0 50 100 150 200 250 300 Exposure Time (h)

Fig. 8. Weight gain as a function of time for FeAl alloys exposed to a gas mixture at 800°C with P,, = 10m6 atm and PO, = 10m2’ atm (courtesy of J. H. DeVan, Oak Ridge

National Laboratory, Oak Ridge, Tennessee).

the addition of chromium up to 5 at% does not improve the ductility significantly, 0.12 at% B (300 ppm) enhances the ductility of Fe-35.8 at% Al from 2.2-5.6%. On the other hand, even as little as 0.06 at% B resulted in severe edge and end cracks during hot rolling. Such edge cracking during hot rolling was not observed with the addition of zirconium, and the ductility enhancement was observed in the range of 0.1-0.2 at% Zr.32 Exten- sive evaluation of FeAl containing different levels of boron and zirconium additions led to the under- standing that the boron-to-zirconium atom ratio plays an important role in enhancing the room- temperature ductility, and the optimum ratio was observed to be about two.

Maziasz et a1.33 have recently shown that cast FeAl alloys designated as FA-385, consisting of Fe-21.1 -41-0.42 MO-O. 1 Zr-0.03 C (all wt%), can possess reasonable room-temperature ductilities (on the order of 68%) and high-temperature strengths, particularly with microadditions of boron (see Figs 10(a) and (b)). Tensile elongations and strengths of FeAl alloys in air improved

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Nickel and iron alumhides: an overview 365

420 1 s ; 340 sl 5 8 260 0 .9

* 160

A 200T . 4OO’C

l 6OO’C

l 700%

100 1 I I I I I

30 34 30 42 46 50

60

Aluminum (at. X) (4

1

48 A 200*c

v 4oo*c

2 l 600.C

36 ._ 5 s P o 24 z

12

0 -I

30 34 38 42 46 50

Aluminum (at. %)

@I

Fig. 9. (a) Plot of yield strength as a function of aluminum concentration for FeAl alloys tested at temperatures to 7OO”C, and (b) plot of tensile elongation as a function of aluminum concentration for FeAl alloys tested at temperatures to 700°C.

significantly when they were processed by extrusion as opposed to rolled or cast properties; extrusion of nitrogen-atomized powder of FA-385 showed further ductility enhancement (see Fig. 11). Creep resistance of iron aluminides is significantly lower than that of many of the high-temperature alloys, and lack of creep strength limits their present use as structural materials above 800°C. Low density, high-specific strength, and high-specific stiffness of FeAl over steels and other nickel-based alloys are attractive for a variety of applications, and the tensile elongations of FeAl are significantly better than TiAl and Ni,Al intermetallics. The high-tem- perature strengths of FeAl alloys are significantly better than a variety of polymeric composites and aluminum alloys, and also have the added advantage in that FeAl alloys can be obtained at a relatively low cost as compared to many of the intermediate temperature range materials.

FA-385 FA-385 FA-385

w/25 ppm w/so ppm

Boron BWOil

(4

1250 # I

El Y&-Tested I” Aor

‘;;; 1000 - n YeWTested I” Oxygen

._. .

q UTS-Tested in AN

% 750 1 ,E+ UTS-Tested en Ox ZY9

Jz +a

E E! 500

G 250

0 i

FA-385

L

FA-385 FA-385

w/25 pprn w/50 ppm

BCXOIl BOrOn

(b)

Fig. 10. Effects of boron microalloying additions on as-cast FeAliFA-385 alloys tensile tested at room temperature, including (a) total elongation and (b) yield and ultimate tensile strengths (courtesy of P. J. Maziasz and C. T. Liu, Oak Ridge

National Laboratory, Oak Ridge, Tennessee).

.- w 6 g 10 5 iii

ii 5

5 I-

O

Rolled at Rolled at Extruded As-Cast Powder

900°C 900°C at 900°C Extruded

HTl HT2 at 11oo’c

Fig. 11. Tensile elongation data at room temperature for FeAl (FA-385) processed by different methods in both air and oxygen environments. HTl is a heat treatment of 1 h at 9OO”C, and HT2 is 1 h at 1000°C. All other specimens were

only stress relieved for 1 h at 750°C in air.

Apart from the above advantages, iron aluminides (unlike many commercial heating element materials such as Ni-20 wt% Cr, iron alloys, nickel alloys, and superalloys based on nickel, cobalt, and iron) possess high electrical resistivities. Figure 12(a) shows a linear increase in electrical resistivity of binary FeAl with an increase in aluminum content, the

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366 S. C. Deevi, V. K. Sikka

I

0 4 8 12 16 200

WtXotAl (4

---t Haynes 556

ii 2

b 0 ~~~‘~~~‘~~~I~~~I~~~‘~~,

iii 0 200 400 600 800 1000 1200

- w

Temperature (“C) (b)

Fig. 12. (a) Variation of electrical resistivity of Fe-Al alloys with increase of aluminum content and (b) a comparison of electrical resistivities of iron aluminides with those of com- mercially available Haynes 214, 230, and 556 alloys. The composition of FAH is Fe-28 Al-5 Cr-0.05 B (at%), and the composition of FA-129 is Fee28 Al-5 Cr-0.5 Nb-0.5 MO-0.1

Zr-O-2 B (at%).

electrical resistivities of FeAl with 35 at0/o of alu- minum being on the order of 150 pohm-cm or higher. The high electrical resistivity of iron alumin- ides is in stark contrast to the electrical resistivities of most conventional electrical-resistance alloys currently in use, and the substitution of transition elements to binary FeAl may actually increase the resistivity values beyond the 150 pohm-cm. It is important to note that the electrical resistivities of iron-aluminide alloys are typically above 150 pohm-cm and can be as high as 200 pohm-cm (see Fig. 12(b)). Electrical resistivities of superalloys such as Haynes 214, Haynes 230, and Haynes 556 are 2040% lower than the electrical resistivities of Fe,Al-based alloys such as FA-124 and FAH. It is interesting to note that both Fe,Al and FeAl possess a positive temperature coefficient of resistance up to at least 600°C.

Consistent with their high electrical resistivities, the thermal conductivities of iron aluminides are

also lower than for nickel-based alloys. The thermal expansion coefficients of iron aluminides are similar to those of stainless steels. The similarity of thermal expansion coefficients of steels and iron aluminides suggests that iron aluminides can be used as bond- coat materials for thermal spray coatings, or iron aluminides can be coated or weld overlayed on a variety of chemical and petrochemical reactors against high-temperature oxidation, carburization, and sulfidation.

4 PROCESSING OF INTERMETALLICS

Intermetallics can be processed by various means, and the processing methods vary depending on the product (i.e. plate, sheet, foil, bar, rod, and tube). Figure 13 provides a schematic of possible processing methods by means of melting and by the use of elemental powder-processing (reaction synthesis) techniques. Although the elemental powder approach is novel, melting remains the primary processing technique to obtain a variety of cast, wrought, and powder-metallurgical prod- ucts. Melting techniques will determine whether an intermetallic can be obtained economically with good control of the composition, and with minimal or no porosity in the cast structure. Melting methods can be used to obtain atomized pow- ders for powder-metallurgical processing of inter- metallics.

The melting techniques such as air-induction melting (AIM), vacuum-induction melting (VIM), vacuum-arc melting (VAR), and vacuum arc-double electrode melting (VADER) are applicable to melt and obtain intermetallics with different levels of purity. Along with melting techniques, process parameters such as crucibles; melting atmosphere (air, inert gas, vacuum); furnace loading of low- and high-melting-point metals; and most impor- tantly, the ability to use recycled material with close control of target composition need to be considered.

Of the above techniques discussed in the previous section, AIM is the most economical technique and can be used if a protective slag can be gener- ated easily on top of the melt. Our experience shows that nickel aluminides can be obtained by the AIM process with an argon cover while the VIM process is most suitable for iron aluminides.

The VIM process will eliminate the solidification porosity associated with the escape of dissolved hydrogen (generated from dissociation of mois- ture) in AIM castings.

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Nickel and iron aluminides: an overview 367

Fig. 13. Possible processing methods to obtain intermetallic compounds such as bar, sheet, foil, and tube.

4.1 Factors affecting melting of intermetallics

Unlike conventional steels, alloys, and superalloys, melting of intermetallics requires consideration of factors including: (1) the difference between the melting points of aluminum and transition metals such as nickel, iron, titanium, and niobium (see Table 5); (2) the large amounts of aluminum (lo-30 wt%) as opposed to < 0.5% of aluminum in most commercial alloys and < 5 wt% Al in superalloys; (3) the exothermic nature of the inter- metallic compound formation (see Table 1); (4) the

Table 5. Densities, crystal structures, melting points, and boiling points of aluminum, nickel, and iron

Element Density (g/cc)

Crystal structure

Melting Boiling point (“C) point (“C)

Al 2.702 fee 660 2520

Ni 8.90 fee 1453 2910

Fe 7.86 bee (a). fee (a) 1536 2860

Ti 4.5 hcp(a), bee(P) 1670 3290

Nb 8.57 bee 2468 4750

co 8.9 hcp(a)> fee(P) 1495 2870

higher melting points of some intermetallics com- pared to aluminum and transition metals; and (5) the reactivity and vapor pressures of alloying ele- ments in atmospheres such as air and vacuum at the melt temperatures.

In general, aluminide formation is an exothermic process (see Table 1 for heats of formation), and the exothermic nature of aluminide formation has been a great concern for melters due to the uncon- trolled nature of the reaction. The lack of expertise in melting alloys with high-aluminum contents and the safety issues associated with the melting practice limited the commercial-scale melting and casting of intermetallics. On the other hand, the exothermic reaction has been successfully implemented by a variety of investigators to obtain aluminides from elemental powders by combustion- or reaction- synthesis techniques. 3b39 It has been shown that only melting of aluminum is needed for the forma- tion of intermetallics by combustion- and reaction- synthesis techniques while melting of aluminum and transition metal is needed for most industrial melting and casting techniques. Although reaction-

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368 S. C. Dee\ li, V. K. Sikka

synthesis principles have been extended by the authors to obtain rods of nickel and iron alu- minides,40,4’ reaction synthesis and densification approach cannot provide large parts with intricate shapes and configurations. Also, it cannot be applied to every intermetallic system in an economical way. Therefore, conventional melting processes need to be modified to effectively utilize the exothermic reactions during the melting of inter- metallics. Proper utilization of the exothermic reaction, as in the case of combustion-synthesis technique, can result in substantial energy savings and may also reduce the time needed for melting.

The proper use of the heat released can also result in control of the chemical composition and forma- tion of oxidation product (slag).

4.2 Conventional melting process

In a conventional melting process, nickel aluminide is obtained by melting nickel and loading alu- minum into the molten nickel melt stock. This process requires heating the melt stock to a tempera- ture of 1600°C prior to the addition of aluminum.

Melting of aluminum is generally not preferred since molten aluminum can seep through the crucible as the temperature is raised to 800°C or higher and damage the induction coil with serious safety issues to the furnace and the operators. The addition of aluminum to the molten nickel melt stock can raise the temperature of the crucible by several hundred degrees and will also result in oxidation of aluminum and some other alloying elements.

To validate the issues associated with the conven- tional melting process, a batch of alloying elements of composition IC-221M (see Table 2) was loaded into the crucible. This composition exhibits high strength as compared to the other N&Al-based alloy compositions developed at ORNL and also requires about 6 at% less aluminum than pure N&Al (IC-50). Alloying elements such as Cr, MO, Zr, and B were sandwiched within the nickel melt stock to prevent oxidation. The entire melt stock was inductively heated in a zirconia crucible.

Three W-5% Re versus W-26% Re thermocouples were located in the crucible to measure the tempera- tures. The addition of aluminum to the molten nickel melt stock at 1600°C initiated a violent exothermic reaction instantaneously. A peak tem- perature of 2300°C was noted within 1 min. The crucible was at a temperature of 2100°C for several minutes, and a high-temperature vapor cloud escaped from the crucible as a result of the oxidation of aluminum (see Fig. 14). It is important to note

Fig. 14. A photograph of the high-temperature vapor cloud formed after the addition of aluminum to molten nickel melt

stock.

that the oxidation of aluminum and zirconium is much more exothermic and releases significantly higher energies than the heats of formation of nickel aluminides.

2 Al + 3/2 O2 + A1203, AH = AOO.2 k Cal/g mole (2) Zr + O2 + ZrO,, AH = -261.8 k Cal/g mole (3) Therefore, the addition of aluminum to the melt stock is undesirable, owing to the possibility of exceeding the safe operating temperatures of a ceramic crucible (such as alumina, zirconia, and mullite). Also an uncontrolled oxidation could raise the temperature of the crucibles close to their melting points. Further, the temperature rise asso- ciated with the addition of aluminum to the molten nickel melt stock could potentially lead to disastrous results if molten metal comes into con- tact with water if the crucible cracks.

4 3 . Exo-MeltTM process

Research at ORNL addressed the above safety issues and focused on the melting, casting, and processing of N&Al- and Fe,Al-based alloys for commercial utilization of intermetallics. Melting and casting of over 100 heats of N&Al- and Fe,Al- based alloys were accomplished successfully with the Exo-MeltTM process to obtain several thousand kilograms of nickel and iron aluminides. The Exo- MeltTM process consists of dividing the melt stock into several parts in a furnace-loading sequence

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Nickel and iron aluminides: an overview 369

such that a very exothermic reaction with a high- adiabatic combustion temperature is favored initially, leading to a molten product. In the case of nickel aluminides, the formation of NiAl is highly exothermic, and the adiabatic combustion tempera- ture is the melting point of NiAl. On the other hand, N&Al formation would only give rise to a combustion temperature of 13OO”C, and a value of

= 1400°C (melting point of Ni,Al = 1395°C) can be achieved only if the melt stock is preheated to the melting point of aluminum.42A4

An optimized furnace-loading sequence, based on the Exo-MeltTM process, is shown schematically in Fig. 15 for melting N&Al- and Fe,Al-based alloys.

In this sequence, nickel was split into two parts.

First, the top part of the loading is combined with all the aluminum to initially form NiAl. Second, the remaining nickel of the melt stock is loaded at the bottom. All of the alloying elements of a par- ticular melt stock are loaded between the top and bottom layers of nickel. It is important to note that several pieces of aluminum melt stock were vertically placed inside the wall of the zirconia crucible so that the molten aluminum would react with the nickel melt stock in a controlled manner.

Melting was carried out under an argon cover by monitoring the temperature of the melt stock.

Exothermic-diffusional reactions between the molten aluminum and the solid nickel melt stock were observed at different locations in the crucible.

Temperature measurements and a video recording of the melting confirmed the occurrence of exo- thermic reaction in the top layer. Thermodynamic calculations suggest that the reaction of molten aluminum with the nickel melt stock at a tempera-

Thermocouple Leads

I

Fig. 15. A schematic representation of furnace loading employed in the Exo-MeltTM process.

ture of 700°C or above leads to molten NiAl. Alloy- ing elements are dissolved as the molten NiAl drips down onto the heated nickel melt stock.

Additional NiAl liquid is formed as the furnace temperature is raised, and the exothermic reaction continues until the reaction between molten alu- minum and heated nickel melt stock is complete.

The furnace-loading sequence shown in Fig. 15 allowed the rate of reaction to be controlled. The rate of reaction between the molten aluminum and the nickel melt stock is also controlled by the large- size nickel melt stock which is at least on the order of 0.5-2.0 cm diam. The crucible temperature never exceeded the melting point of NiAl. The crucible life is extended since the crucible experi- ences a maximum temperature of 1640°C for about 10 min at most.

4.4 Advantages of the Exo-MeltTM process over conventional melting process

In the Exo-MeltTM process, exothermic reactions are initiated just after the melting of aluminum, whereas the melting of nickel melt stock is required in the conventional process. In addition to the above, the power required to melt a batch of IC- 221M by the Exo-MeltTM process is 47% lower (2.92 kWh) than with the conventional process (5.5 kWh). Time required to melt and pour the intermetallic by the Exo-MeltTM process can be reduced by 50% as opposed to a conventional pro- cess. All of these resulted in substantial savings in energy and a reduction in processing cost. Unique attributes of the Exo-MeltTM process are com- pared with the conventional process in Table 6.

Target composition of an alloy can be reached easily and safely with the Exo-MeltTM process as compared to a conventional process. Recovery of the elements by Exo-MeltTM and conventional processes is shown in Table 7. Target composition can be attained even in melting of larger melt stocks, and recovery of the elements from a 272 kg melt stock is also shown in Table 7. The tech- nique has successfully been implemented in the melting of iron and titanium aluminides. High levels of aluminum present in the alloys allowed forma- tion of A1203 (slag) on top of the molten metal.

The low levels of oxygen and nitrogen (< 40 ppm by weight) indicate that the oxide film was highly impervious. Oxygen, nitrogen, and carbon impuri- ties are generally around 0.012 wt%. Melting of the alloys in air under an argon gas cover allowed the reduction of oxygen, nitrogen, and hydrogen levels and enhanced the recovery of elements.

References

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