Vol. 2 No. 4 Dec. 1998
JOURNAl. OF SHANGHAI UNIVERSITY
Formation of Intermetallic Compound in Iron-Aluminum Alloys Wang Xingqing
W o o d J. V .
(Shanghai University)
(University of Nottingham ,UK )
Sui Yongjiang
Lu Haibo
(Shanghai University)
(Central South University of Technology)
The mechanism of iron and aluminum intermetallics formation in the reactive sintering of iron and aluminum mixing powders has been studied by investigating iron-aluminum diffusion couples. The couples were treated at 600"C, 700C, 800'C, 900=C and 1000 C respectively. It was found that an Al-rich intermetallics FeAls has formed in iron adjacent to the interface of iron and aluminum by aluminum diffusion into iron at 600C (below the eutectic temperature),and that in the case above 700 ~(' (above the eutectie temperature) there was a liquid,an intermetallics Fe2Als has formed in both side of the interface. The diffusion of iron and aluminum atoms is companied with the Fe A1 reaction during the treatment under the both conditions. The diffusion coefficients of iron and aluminum and the activation energy were determined. The mechanism of the intermetallics formation in the couples is also discussed. Abstract
Key words
1
diffusion, diffusion couple, Fe-A1 alloy, intermetallics, Fe-A1 reactive sintering
Introduction
Fe-A1 intermetallics have great potentialities for structural applications owing to their good mechanical p r o p e r t i e s , excellent corrosion and oxidation resistance and inexpensive raw materials. H o w e v e r the use of these intermetallics is limited by their brittlensess and poor workability at low t e m p e r a t u r e s , even t h o u g h they can be improved by the addition of third elements ¢1-3]. Powder m e t a l l u r g y has become one of the i m p o r t a n t processing method for making Fe-A1 intermetallics E4-s?. In such process rapidly solidified prealloyed powders are consolidated by hot isostatic pressing or hot extrucion. N e v e r t h e l e s s , t h i s method involves an expensive processing route to prepare the prealloyed powders and does not lead itself to a wide range of application. An alternative powder processing method for making Fe-A1 intermetallics is the reactive sintering of elemental or partially prealloyed powders[9 12] utilizing simple powder m e t a l l u r g y routes (i. e. powder m i x i n g , p r e s s i n g and sintering). This approach is characterized by low processing t e m Received Dec. 24,1997 Wang Xinqing, Asso. Prof., School of Material Science and Engineering, Shanghai University, 149 Yanchang Road, Shanghai 200072
p e r a t u r e , s h o r t processing time and considerable flexibility in compositional and microstructural control. A l t h o u g h this method has been successfully applied, the mechanical properties of the final c o m p o n e n t s are still lower than that obtained from cast materials. This is caused by the pores which occur as a result of different diffusivity of the involved elements and possible gas evolution E9 1el. T h e sintering of compacts made from iron and aluminum powder mixtures is accompanied by the initial reaction between A1 and Fe to give a transient liquid phase w h i c h , i n t u r n , l e a d s to some densification of the compacts. All these result in the formation of intermetallics. T h e preferential diffu sion of aluminum into iron creates pores at sites occupied by the original aluminum particles. H o w e v e r according to Sheasby EgJ,the large expansions which occur above melting point (between 600 and 700 C ) can not be explained solely by unbalanced diffusion coefficients (Kirkendall effect). He concluded that binder rapid buring-out can generate stresses which contribute to expansion and therefore p o r o s i t y , possibly overshadowing the actual s y s t e m behaviour. In this paper we present a study of intermetallic compound f o r m a t i o n , pore formation as well as iron and a l u m i n u m diffusion in Fe-A1 diffusion couples. T h e aim of this research is to understand the mecha-
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Journal of Shanghai University
nisms of Fe-A1 intermetallics formation during the re active sintering of iron and aluminum powder mixtures.
2
Experimental Methods
Diffusion couples of Fe-AI were prepared by riveting a thin a l u m i n u m rod into an iron plate. In order to u n d e r s t a n d both the subsequent diffusion behaviour of iron and aluminuna and the formation of Fe-A1 intermetallic compounds below or above the reactive sintering t e m p e r a t u r e (655 C ) , some samples were initially heated to 600 C and held for 1 , 2 , 4 and 8 hours and some samples were treated at 7 0 0 C , 8 0 0 C , 9 0 0 C and 1000(' for 1 hour. All heat treatments were conduc*ed in a horizontal tube furnace under hydrogen a t m o s p h e r e . T h e samples were then rapidly cooled to ambient t e m p e r a t u r e in a stainless steel cooling zone attached to the furnace. Interdiffusion of iron and a l u m i n u m , i n t e r m e t a l l i c s was formation were analyzed by m e t a l l o g r a p h y , S E M , E D X .
3
Results
3.1
is evidence of iron diffusing along the a l u m i n u m grain boundaries. Fig. 3 shows the distribution of Fe (white s p o t s ) along these grain boundaries. T h e analysis indicates that the composition in Fe-A1 intermetallics area does not vary with treating t i m e , although intermetallics thickness increases with heat t r e a t m e n t time, as shown in Fig. 4. Fc
FeA1FeAI3
10 20
40
30 peaks
FeAI~
Fe 60
Fe
t
Fe
80
1oo 20
Fig. 2
X ray diffraction analysis of the sample treated at 600C for 8h
Heat treatment at 600C
T h e SEM photo of the samples treated at 600~C for 8 hours is shown in Fig. l , f r o m which it can be observed that a new phase formed in iron zone near the Fe-A1 boundary. EDX analyses indicated that the new phase contained about 75°/oo A1 ( a t % ) approaching to the stoichiometric composition relation of FeAla,
?
which was proved by X - r a y diffration analysis as s h o w n in Fig. 2. It also d e m o n s t r a t e d that there was negligible AI at Fe area although there was considerable A1 diffusion into the original Fe near the interface where an A1 - rich intermetallics was f o r m e d . T h e r e Fig. 3
Fig. 1
The microstructure of the sample treated at 600C for 8h 74160 (SEM)
Fe distribution in AI area of the sample treated at 600C for 8h X100 (SEM) Fig. 5 is a higher resolution m i c r o g r a p h of the reacted region near the Fe-Fe/A1 interface. It shows the reacted intermetallics and iron. It clearly demonstrates that the intermetallics is initially forming along the grain boundary before growing into the Fe grains. No aluminum is detected in the unreacted Fe region. Fig. 6 shows that the grain size of the intermetallics (in the middle) is m u c h finer than the initial iron grain size. This implies that the intermetallics nucleation is independent of the original Fe microstructure. Fig. 6 also d e m o n s t r a t e s that there exists a n u m b e r of pores in the intermetallies.
Vol. 2
No. 4 Dec. 1998
Wang X. :
Formation of Intermetallic Compound i n . . .
307
0.45
termetallics (Fig. 8). The EDX analysis shows that
0.35
chiometric composition
c o m p o s i t i o n of t h e i n t e r m e t a l l i c s a p p r o a c h e s t h e s t o i of Fe2Als ( 2 6 - 2 8 % F e - A 1 ) ,
which contains a higher iron concentration than that
.~ 0.~5
o b s e r v e d in t h e s a m p l e t r e a t e d at 600 C , b u t t h e r e is still no e v i d e n c e of a l u m i n u m in t h e u n r e a c t e d iron.
0.I5 ~
X - r a y d i f f r a c t i o n a n a l y s i s of t h e s a m p l e t r e a t e d at 7 0 0 C s h o w s t h a t the n e w p h a s e is Fe~Als ( s e e F i g .
0.05
9). T h e m i c r o h a r d n e s s of up to 1340 H V have b e e n
0
2
Fig. 4
3
4
5 time (h)
6
7
8
k
The relation between intermetallics thickness and treating time at 600°C
k
!,.
%
The microstructure of the sample treated at t000 C for lh (SEM)
Fig. 7
,,), ~.~ ,. i
Fig. 5
Intermetallics growing morphology in the sample treated at 600C for 8h X 500 (SEM)
t ÷
Fig. 8 e 4.-t, "-~"
Pore morphology in the intermetallics treated at 1000'C for lh X1000(SEM)
-"
Fe
t;,
-e
Ilpeaks
,
O
Fig. 6
3.2
Microstrueture of the intermetallics in the sample treated at 600~C for 8h X400
T h e S E M m i c r o g r a p h s of t h e s a m p l e s t r e a t e d at 1000°C for 1 h o u r is s h o w n in F i g . 7. A n e w p h a s e has f o r m e d in t h e i n t i a l A1 a r e a a n d t h e p a r t of F e a r e a ,A1 has d i s a p p e a r e d a n d a l a r g e p o r e r e m a i n s in t h e c e n t r e of t h e i n t e r m e t a l l i c s . P o r o s i t y is also p r e s e n t in t h e in-
Fe JI
Fe2AI~
H e a t t r e a t m e n t at 7 0 0 ' C - 1 0 0 0 ' C 10
20
J: t 40
Fo Fe2AI~ 1
Fe
LA Fe't",LI 60
80
Ioo 2O
Fig. 9
X-ray diffraction analysis of the sample treated at 1000 C for lh
Journal of Shanghai University
308 measured for the intermetalIics,which is also higher than that of the samples treated at 6 0 0 C . Measure-
gram [~4]. The discontinuous interface between Fe-At and Fe demonstrates that A1 diffusion is controlled by the intermetallics formation. Fig. 12 is a schematic drawing of the process based on these observation:A1
ments of the thickness and diameter of the intermetallies zone show that it grows into Fe and AI and the overall thickness increases with temperature as
d i f f u s i o n ~ r e a c t s with Fe--~FeA13. This demonstrates
expected. Fig. 10 is a general s u m m a r y of these mea-
that the reaction of intermetallics formation is con-
s u r e m e n t s , from which it can be seen that the inter-
trolled by A1 diffusion. As a result A1 does not exist
melallics grows mainly into the original aluminum
in Fe region. There are some evidences that a number
area after initial reaction with iron.
of Fe atoms diffuse toward A1 region even at 600~C.
0.8 0.7
Fe does not react with A1 and distributes along A1
0.6
and Fe concentration in the A1 increases with treat-
0.5
ment time at 600°C.
grain boundaries,so that Fe distribution is continuous
,~ 0.4 : 0.3~
Al
0.21
m Fe
0.1 0 700
800 temperature
Fig.
10
4
Discussion
900 ('C)
1000
AI
The relation between intermetallics thickness and treating temperature (for lh)
D i f f u s i o n of AI and Fe at 600C
4. 1
0
From Fe A1 binary phase diagram E14],the melting temperature of pure aluminum and the lowest eutectic temperature is 660C and 655 C respectively. Therefore all these phases should be solid at 600C. The results show that the intermetallics FeAI~ has been formed by A1 diffusion into Fe. Fig. 11 shows the distribution of iron and aluminium across the reacted region after 8 hours. There is littel evidence of matrix diffusion into the unreacted Fe and AI zone. The discomposition in the intermetallics region varies about 1%--2%, "-'100
which is consistent with the phase dia-
--
-- -- -
-
A ~
~
Fig.
4. 2
12
>X
Schematic illustration of A1 diffusion and intermetallics formation in Fe zone
Diffusivity of Fe and AI below reactive sintering temperature
In order to measure the diffusivity of iron and aluminum through the intermetallics zone, the distance from the reference plane across the intermetallics and the concentration at the interface have been measured. Consequently, the diffusion coefficients of iron and aluminum can be calculated by Fiek's second law:
8 CA/3 t : D A ?~CA/O .2:2,
.
where CA is the concentration of A atoms (mole fract i o n ) , t is diffusion time ( s ) , x is diffusion distance ( m ) and DA is the diffusion coefficient of A atoms.
0J
"~ oo
o
60
Fe
4O
20
The diffusion coefficient D usually varies with concentration in a given diffusion couple and markedly !
AI
.l~ interface -0.8 -0.6-0.4 0.0 0.4 0.6 The distance from the i n t e r f a c e o f F e a n d Al(mm) Fig.
11
Distribution of Fe and A1 in the diffusion couple treated at 600 C for 8h
increases with temperature rising. Here the source of Fe and A1 diffusion is pure Fe and pure A1 and this remains constant for diffrent treatment t i m e , s o that DA can be expected to be constant for a given temperature. Therefore it is possible to use the Grube solutionE~S]of Fick's second law as follows:
Vol. 2 No. 4 Dec. 1998
Wang X. :
Formation of Intermetallic Compound in.,.
309
zone. T h u s intermetallics zone grew with increasing treatment temperature.
C , - - C a / C ~ - - C o = e r f ( x / 2 Dv/DTtat), where C, is the surface concentration or the concen-
4. 4
tration at the Matano interface,CA is the concentra-
If the intermetallics thickness on each side of the prior Fe and A1 is selected as tile largest diffusion dis tance to the Matano interface,lh., and DA, can be cal-
tion of A atoms at some distance x from the surface and Co is the initial concentration of A atoms in the diffusion couple B zone. so t h a t , C , = 1 and Co=O. So, CA = 1 - - e r f ( x / 2
~/DAt).
T h u s the diffusion functions of Fe atoms and A1 a t o m s can be given as the following: CAI= 1 - - e r f ( x / 2 Cw= l--erf(x/2
D i f f u s i v i t y of Fe and A! above 700 C
culated by the G r u b e solution ~'I. T h e relation between the diffusion coefficients and t e m p e r a t u r e can be given according to the expression: D = D,,e xp ( -- E . . I / R T ) , where Do is a constant for a given diffusion couple,E~
q~Alt),
is the activation energy (J • tool L) of diffusing of A
D,/~vd).
T h e largest values of diffusion distance and concentration were used to calculate Dv~ and DAt in the present case. Here DA, is the A1 diffusion coefficient in
a t o m s , R = 8 . 313J • mol x • K k,and T is temperature in Kelvin. This equation can be expressed as l o g D = l o g D 0 - - EA/RT.
iron and DF, is the Fe diffusion coefficient in alu-
Plotting lOgDA, and logDF, versus 1 / T (Fig. 13) gives
minum. T h e results given in T a b l e 1 show relatively constant values of D with treating time.
an a p p r o x i m a t e l y liner relation, so l h a t , D O F ~ = I . 3 X 1 0 3units, E > = S 3 . @ g K J - m o l D o A I = 5 . 2 X 1 0 ~units, EAI = 6 . 6 1 K J ' m o l
Table 1 Relationship of diffusion coefficient with treating time treating time (h)
1
2
4
8
DF~(cm2sec-1 )< 10 n )
4.62
5.22
4.92
4.02
DAl(cm2see-~ )< 10 9)
4.45
3.72
3.85
4.39
4.3
T h e results confirm that iron diffuses into the liquid phase A1 more rapidly than aluminum diffuses into sold phase Fe. 0
D i f f u s i o n of Fe and AI above 700 C
T h e r e is limited solubility ( 0 . 0 2 ~ )
-1
of Fe in AI at
t e m p e r a t u r e lower than the A1 melting point E~n. H o w e v e r , u p to 44°/oo Fe can be dissolved in liquid A1. A1 and Al-rich phase with low melting point will melt above reactive t e m p e r a t u r e ( 6 5 5 C ) , so that an amount of Fe is dissolved in the liquid A1 when the samples are treated above 700~C. T h u s Fe dissolut i o n , F e and A1 atoms interdiffusion, synthesis reaction and intermetallics formation occur in the liquid A1. T h e initial process of intermetallics formation in the liquid state proceeds as follows .. eutectic reaction between FeA13(formed before A1 melting during hea'cing) and A1--~eutectic liquid-~Fe and AI interdiffusion --~ intermetallic reaction ~ FeeAl~. Fe atoms diffuse rapidly in liquid A1, so that the process of intermetallics formation is finished very quickly. Most of intermetallic compounds were formed in the initial A1 zone. This was attributed to the synthesis reaction in the liquid A1. As a result intermetallics Fe,~AI~ was formed. This intermetallics has been reported in previously reactive sintering of Fe and A1 powder mixture :a'~'~e?. Due to Fe dissolving into A1 and A1 also diffusing into F e , i n t e r m e t a l l i c s developed toward iron
~; *
-2
9-4 -5 DAI
-6 -7 8.0
8.5
9.3
10
I/T(xl0aK) Fig. 13 The relation between diffusion coefficients and temperature
4.5
Intermetallics and pore f o r m a t i o n
T h e results show that an Al-rich intermetallics FeAI~ will form below the reactive sintering tempera ture by A1 diffusion into Fe and another intermetallics FeeAls will form above the reactive sintering temperature by Fe diffusion into the liquid A1 and AI diffusion into solid Fe. T h e products formed below and above reactive sintering t e m p e r a t u r e have a constant stochiometric composition which does not change with holding time at a given temperature. These two intermeta[[ics have no fixed orientation relationship
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Journal o f Shanghai University
with the orginal Fe and A1. T h e y have finer grain
Properties, Microstructure
t h a n the original Fe and AI. During A1 diffusion into F e , the m o b i l i t y of A1 a t o m s along grain b o u n d a r i e s
3
of Fe is higher t h a n t h r o u g h the m a t r i x . As a r e s u l t the s y n t h e s i s reaction and i n t e r m e t a l l i c s f o r m a t i o n occur p r e f e r e n t i a l l y at the grain boundaries. A b o v e the
4
reactive s i n t e r i n g t e m p e r a t u r e Fe2AI5 is more s t a b l e t h a n the o t h e r i n t e r m e t a l l i c s . A large n u m b e r of pores
5
exist in the i n t e r m e t a l l i c s . The m o r p h o l o g y of the pores s h o w t h a t they are unconnected and thus p r o b a bly f o r m as a r e s u l t of gas p r o d u c t s (see Fig. 8).
5
6
Conclusions (1) A t 600"C the i n t e r m e t a l l i c s FeA13 is formed on
the iron side of a Fe-A1 diffusion c o u p l e , w h i l e at temp e r a t u r e s above the melting point of A1 the intermetallics Fe2A15 is formed in whole A1 area and p a r t of Fe area beside the interface. In both cases t h i c k n e s s of the i n t e r m e t a l l i c s increases with heat t r e a t m e n t time and t e m p e r a t u r e . (2) T h e i n t e r m e t a l l i c c o m p o s i t i o n s are c o n s t a n t for a given heat t r e a t m e n t t e m p e r a t u r e and do not change with heat t r e a t m e n t time. ( 3 ) T h e i n t e r m e t a l l i c reaction is c o n t r o l l e d by Fe and A1 interdiffusion. Fe and A1 d i s t r i b u t i o n in the couples is discontinuous. ( 4 ) T h e i n t e r m e t a l l i c s reacts and forms p r e f e r e n tially at the grain boundaries. A large n u m b e r of pores form d u r i n g the reaction and exist in the intermetallics.
References 1
Mckamey C. G. and Horton J. A. ,The effect of molybdenum addition on properties of iron aluminides ,Metallurgical Transaction A ,20A : 751 (1989) 2 Mckamey C. G. ,Maziasz P. J. ,Goodwin G. M. ,et al. ,Effects of alloying additions on the microstructures,mechanical properties and weldability of FezAl-based alloys ,Materials
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(1994) Wang X. and Wood J. V. ,Role of nickel in Fe3AI intermetallic compound formation, Powder Metallurgy, 38 (1) : 59 (1995) Bordeau R. G. ,Development of Iron Aluminides,AFWALTR-87-4O09(Air force Wright Aeronautical Laboratories, Wright-Patterson Air Force Base ,OH May 1987) Knibloe J. R.,Wright R.N. and Sikka V. K.,1990 Advances in Powder Metallurgy, Metal Powder Industries Federation, Princeton, N J, 2 : 219 ( 1990) Sikka V. K. ,Baldwin R. H. ,Howell C. R. ,et al. ,Powder production, processing and properties of Fe3A1,1990 Advances in Powder Metallurgy, Metal Powder Industries Federation, Princeton, N J, 2 (1990) Knibloe J. R. ,Wright R. N. ,Sikka V. K. ,et al. ,Elevatedtemperature behavior of Fe3A1 with chromium additions, Materials Science and Engineering A-Structural Materials Properties, Microstructure and Processing, 153 (1-2 ) : 382 --
386 (1992) 8 Prakash U. ,Buckley R.A. and Jones H. ,Formation of B2 antiphase domains in rapid solidified Fe-A1 alloys,Philos. Mag. ,1991:213 9 Sheasby J. S. ,Powder metallurgy of Iron-aluminum,Int. J. Powder Metall. Powder Technol. ,21(4) :301 (1979) 10 Lee D.J. and German R. M. ,Sintering behaviour of ironaluminum powder mixes, Int. J. Powder Metall. Powder Technol. ,21(1) : 9 (1985) 11 Rabin 13. H. and Wright R. N. , Synthesis of iron aluminides from elemental powders:rection mechnaisms and densification behavior, Metallurgical Transaction A , 22A (2) :277 (1991) 12 Wang X. and Wood J. ,Influence of nickel on reactive sintering of FesA1 intermetallic compound,Powder Metallurgy,36(3) :187 (1993) 13 Zhuang L. Z. ,13uekenhout L. and Duszczyk J. ,Reaction phase-forming and mechanical properties of Fe3A1 produced from elemental powders, Scripta Metallurgica Et Materialia ,30(7) : 909-- 914 (1994) 14 Massalski T. B. ,Binary Alloy Phase Diagrams,ASM,Metals Park, OH,1986:111 15 Giles F. Carter,Principles of Physical and Chemical Metallurgy, ASM, Metals Park, OH, 1979 : 282