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Oxide transformations during sintering of in situ oxidised FeAl powders
MATERIALS SCIENCE & EBGlMEERlNG
Materials Scienceand EngineeringA207
(1996)
105-114
Oxide transformations during sintering of in situ oxidised FeAl powders K. Wolski, F. Thivenot, J. Le Coze Cextre
SMS,
Ecole
des Mines
de Saint-Etienne
158, cows
Fawiel,
42023
Saint-Etienne
Cedex,
France
Received17 April 1995;in revisedfonn 10 August 1995
Abstract
Powder metallurgy has been used to prepare FeAl powders with in situ created oxidation products. The aim of this study was to analyse the types of oxide formed as a function of sintering conditions. Highly oxidised powders have been used to give evidence of these transformations by X-ray diffraction. It has been shown that low temperature sintering (T< 1200 “C) results in the formation of FeAl,O, spinel, while higher sintering temperatures lead to its transformation into n-A&O3 and coarsening. An attempt, based on thermodynamic calculations, has been made to explain the formation of FeAl,O, spinel. Keywo~&: Tntertnetallics; FeAl powders; Mechanical alloying; Internal oxidation; Oxide dispersion; Oxide transformation
1. Introduction
FeAl intermetallic has a B2 ordered structure for aluminium content between 32-52 at.%. For ductility purposes, FeAl is considered for applications in the composition range restricted to 32-40 at.% Al [l]. This level of aluminium is still high enough to provide excellent oxidation and corrosion resistance at high temperatures (T > 0.5 r,> [2] and together with relatively low density and high melting temperature (5.6 g cmU3, 1340 “C for Fe-40at.%Al) offers a potential for high temperature applications [3]. The creep strength of FeAl is, however, insufficient [4,5]. One of the most powerful methods of improving creep properties is oxide dispersion strengthening (ODS) [6]. The oxides can be introduced as a powder or formed in situ by oxidation [7]. In both cases ODS FeAl has to be prepared in two steps: (i) ball milling, which is necessary to ensure a homogeneous distribution of oxides or oxidation products and leads to a composite powder; (ii) sintering, which is necessary to ensure densification. During sintering, the oxidation products are trans0921-5093/96/$15.000 1996- ElsevierScienceS.A. All rights reserved
formed into oxides. It is worth noting that the first step, leading to the composite powder, can be performed either by milling prealloyed FeAl powders (attrition) [7] or by milling elemental Fe and Al powders (mechanical alloying, MA) [8,9]. The aim of this study is to clarify the transformations of oxidation products that occur during sintering. The starting materials for this study are composite FeAl powders with dispersion of oxidation products prepared either by attrition or by MA. Owing to the relatively short milling times, the formation of FeAl in composite powders prepared by the MA process is not complete [lo]; however, the homogeneity of oxidation products dispersion is sufficient for the purpose of the present investigation. Two aspects of this study have already been published: (i) formation of FeAl by MA process [lo] and (ii) formation of carbides resulting from the use of a process control agent [ll]; consequently these aspects will not be discussed in this paper. This investigation is a part of a wider study aimed at improving high temperature slow plastic flow properties of FeAl intermetallic, by dispersion of nanosized oxide particles [12].
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Table 1 Elemental powders used in this study Power
Fe Al FeAl
Producer
Prolabo Koch Osprey
Diameter (pm)
10 8-15 5-50
Oxygen content (wt.%)
Carbon content (ppm)
certified
measured
certified
measured
0.2 1.0 300 ppm
0.5 5.0 3500 ppm
200
140 400 30
2. Experimental procedure 2.1. Powders
Prealloyed FeAl powder containing 0.35 wt.% oxygen has been used for attrition. Very fine (therefore highly oxidised) powders of elemental Fe and Al were used to prepare the composite powder by the mechanical alloying process (Table 1). In some cases 2 wt.% stearic acid (CH,(CH,),,COOH, PM = 284 g) was added to prevent elemental powders from sticking either to the balls or to the container wall [ll]. 2.2. Processing Attrition of FeAl powder was performed in a vertical mill (attritor of Szegvari type) installed in a glove box under flowing argon containing 100 ppm of oxygen, while mechanical alloying (MA) was carried out in a planetary mill. In the latter the containers were filled under air and hermetically closed during milling. In both cases hardened steel containers and balls (lOOC6 steel, 2 mm diameter for attrition and 4 mm diameter for MA; 50:1-balls to powder ratio) were used. The resulting powders were sintered in a hot uniaxial press (40-80 MPa, under argon). Milling and sintering parameters of all powders used in this study, together with initial
125
Method of processing MA MA Attrition
compositions, are listed in Table 2. Preparing the composite powder by attrition results in significantly lower oxygen content compared with the MA process. This is due to cleaner initial powders and processing under argon. However, significantly higher oxygen content was necessary (MA process) to analyse the transformations of oxidation products during sintering. The analysis of Fe,AlC,,, carbide formation, resulting from the addition of stearic acid as process control agent was described elsewhere [II]. One of the methods of carbide elimination, namely suppressing stearic acid and processing in several steps [l l] was applied to prepare MP94 powder. In the present study the MP84/3 powder was processed in six steps in order to increase the oxygen content so that X-rays could be used to analyse the oxide formation. 2.3. Methods of nnalysis
Phase transformations were studied with X-ray diffraction (XRD) using Co KU. radiation from a Philips diffractometer (1730/10) equipped with Dosophatex system. Microstructure was characterised by scanning electron microscopy (SEM), equipped with standard microanalyser (JEOL JSM840 + Tracer). Additional information has been obtained from transmission electron microscopy (TEM).
Table 2 Processing parameters of powders and samples prepared and analysed in this study Powder symbol
Processing method
Initial composition
IMilling time and velocity
Oxygen (wt.%)
Sintering T (“C)/t (min)
Sample symbol
A38
Attrition
FeAl-20
1.5
1170160
FAG5
MP88j2
MA
5
1200140
FT45
MP84/3
MA
Fe-15.13 g Al-4.87 g sa-0.4 g
10
1270180
FT39
Fe-15.0 g Al-5.0 g (without sa)
Attrition 15 h 420 rev min-’ MA 3 h 200 rev min-’ MA 6 x 0.5 h 200 rev min MA 6hf3x2h 200 rev min-’
Fe-15.13 g Al-4.87 g sa-0.4 g
MA 1.5 h +0.5 h 200 rev min-’
1050/30 1150/30 1240/30 1240/240 lOOOj5 1100/60 1300/60
FTl7 FT78 FT79 FTSO FT52 FT50 FT34
MP94
MA
MA+
MP82/2
sa. stearic acid
aluminothermic reaction
g
-1
10
25
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3. Results 3.1. Attrition
Fie 1. General microstructure of FA65 sample sintered at 1170°C frz; n prealloyed and attrited FeAl powder.
General microstructure of composites prepared by low temperature sintering (T< 1200 “C) of the composite powder obtained by attrition of prealloyed FeAl powders is presented in Fig. 1. It consists of nanosized grains (50 to 100 nm) dispersed in the FeAl grains (2 to 5 pm). This dispersion (Fig. 2(a)) produces concentric rings on the diffraction pattern (Fig. 2(b)), which were indexed as corresponding to the FeA&O, spine1 phase. (Additional high intensity spots correspond to the diffraction from some FeAl grains, which were selected at the same time as numerous spine1 grains.) Fig. 2(c) is a ring pattern for a nanosized dispersion of this spinel. Calculations based on a final oxygen content of 1.45
(4
(b)
Fig. 2. Analysis of the dispersed phase crystallography by TEM (FA65-type sample). a. bright field image of the dispersed phase. b. corresponding diffraction patterns. c. theoretical disposition of rings for FeA120, spinel.
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K. Woiski et al. 1 Materials Science and Engineering r1207 (1996) 1OS 114
wt.% and assumption of the formation of FeAl,O, spine1 from the whole oxygen gives 5.4 vol.% of spine1 dispersed in the Fe-39at.%Al matrix. As low temperature sintering (T < 1200 “C) was sufficient for a complete densification of this composite, high temperature sintering was not investigated in this case. Note that T = 1200 “C is somewhat arbitrary. In fact the transformations discussed in this paper are thermally activated, so depend also on sintering time. Let us define, for further discussion, low temperature sintering as T < 1200 “C and short sintering time (t < 60 min) and high temperature sintering as T> 1200 “C and long sintering time (t > 60 min). 3.2. Mechanical alloying
The composite powders prepared by MA are characterised by a high oxygen content, within the range 5 to 10 wt.%, depending on oxidation conditions. Higher oxygen content, compared with powders prepared by attrition, leads to the necessity of increasing the sintering temperature for complete densification. The microstructures of two samples (FT45 in Fig. 3(a) and FT39 in Fig. 3(b)) are typical of those observed at their respective sintering conditions. FT45 was prepared by one-step milling for 3 h (5 wt.% oxygen), followed by low temperature sintering (1200 “C, 40 min). FT39 was prepared by six-step milling with intermediate sievings (6 x 0.5 = 3 h, IO wt.% oxygen), followed by high temperature sintering (1270 “C, 80 min). The comparison of microstructures (Figs. 3(a) and 3(b)) gives an immediate indication of the strong coarsening of the dispersed particles (dark phase). XRD has identified the coarsened particles in FT39 as a-alumina. However, no information could be obtained for particles in the FT45 sample, probably because of the nanometre size and possibly incomplete crystallisation.
(a)
H
I pm
Consequently, thin foils have been prepared and their analysis has indicated the presence of FeAl,O, spine1 in FT45 (Figs. 4(a) and 4(b)), at least for zones containing uncoarsened particles, as in the left bottom corner of the micrograph in Fig. 4(a) (diffraction patterns from Fig. 4(b) were obtained from this area). Note that diffraction spots corresponding to FeAl are not observed because there was only one FeAl grain in the analysed area and this grain obviously has an orientation out of diffraction conditions. Additional analysis of FT39 sample (Figs. 5(a) and 5(b)) has confirmed the presence of cc-alumina (Fig. 5(c)), as indicated by a coherent indexing of the diffraction pattern (Fig. 5(d)). As no information could be obtained by X-ray analysis on the FT45 sample containing 5 wt.% oxygen and the TEM analysis cannot provide the global information about phase constitution, an additional charge of powder with 10 wt.% oxygen has been prepared by a six-step milling process (MP94) and analysed as a function of increasing sintering time and temperature (Fig. 6). The XRD patterns from this figure again confirm that high temperature sintering leads to the formation of a-alumina. Unfortunately, neither the nature of oxidation products nor any transformation leading to a-alumina could be detected for this already high oxygen level. 3.3. Strong exothemic
reaction
The problem of the detection limit of X-ray analysis has been overcome by further increasing the oxidation products by means of strong exothermic reaction. This reaction takes place immediately after opening a container, if MA is interrupted after 1.5 to 2 h of milling. In fact, for these milling times, Fe and Al powders have a lamellar morphology with unoxidised surfaces of high specific area, so are extremely reactive. The reason for
(b)
i-i
I/m
Fig. 3. General microstructure of samples sintered from powders prepared by MA process. a. FT45 sample containing 5 wt% oxygen and sintered at 1200°C during 40 min. b. FT39 sample containing 10 wt% oxygen and sintered at 1270°C during 80 min, note the significant coarsening of dispersed phase in the latter.
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(b)
(a)
Fig. 4. TEM analysis of the dispersed phase in materials prepared by mechanical alloying and low temperature sintering (FT 45-type sample). a. bright field image showing the dispersed phase. b. diffraction patterns, compare with the theoretical disposition of rings for FeAI,O, spine& presented in figure 2~.
(a)
H
loonm
(b)
H
IOOnm
-2:21= l
3 ;4205 ,
. 10 .
0
116 .
40
2212 .
(d) Fig. 5. TEM analysis of the dispersed phase in materials prepared by mechanical alloying and high temperature sintering (FT39-type sample). a. dark field image. b. bright field image. c. diffraction patterns. d. coherent indexing of spots confirming the presence of a-A&O,.
K. Wolski et al. / Materials Science atd Dlgitzeerirzg AZ07 (1996) 105-l 14
I I / 35
30 .
FeAl
A
a-A1203
q
Fe$AIC
FT80 FT79 FT78 FT77
-
- (loo),
'1
e I 40
I
superstructure
b I 45
/
peak
"
I 50
c 11
I
j /. .. ... .. . . . .. .. .j . .. ... .. ... .. ... .. ..i .. ... .. ... .. ... .. .. i .. .. .. ... . . .. .. ...
FT34
j
55
28 [WI : 1iOOWlh
25
x
124O"C, 124O"C, 115O"C, 105O"C,
FT50 : l{OOWlh ... .. ... ... .. .. ...l....................
40MPa, 4OMPa, GOMPa, 80MPa,
240mn. 30mn. 30mn. 30mn.
Fig. 6. Evolution of X-rays diffraction pattern as a function of sintering temperature. Samples FT77-FT80 have been sintered from the same MP94 powder containing 10 wt.% oxygen. Note that any dispersion could be detected for lower sintering temperatures (1050 and 1150°C).
performing this experiment was only to give proof for spine1 formation using XRD and not to prepare a viable material. The XRD patterns of MP82/2 powder, prepared as described above, as well ai of three samples sintered from this powder at increasing time and temperature are presented in Fig. 7. This powder contains free iron, a small amount of FeAl, iron oxide (Fe,O,) and spine1 (FeAl,O,). Note that at this stage aluminium seems to be transformed into FeAl and spine1 but there is still no evidence of the presence of a-alumina. Very short heat treatment of this powder at 1000 “C leads to the decomposition of both iron oxide and FeAl intermetallic and formation of FeAl,O, spine1 in the Fe matrix (Fig. 7, FT52), which additionally contains carbon atoms from stearic acid (2 wt.% stearic acid can result in up to 1.4 wt.% carbon). With further increase in sintering time and temperature (1 h, 1100 “C) the spine1 is progressively transformed into cc-alumina (Fig. 7, FT50) and finally after 1 h at 1200 “C this transformation is almost complete (Fig. 7, FT34). At the same time we observe the formation of Fe& carbide. We consider that this change of composition from FeAl intermetallic to Fe-Fe& cermet (Fig. 7, FT34), due to the presence of stearic acid in the starting powders and complete oxidation of aluminium, does not affect the formation of the spine1 and its transformation into a-alumina.
30
35
45
40
50
55
28 Idegl ' m-
-
FeAl20 4 a.Al
0 2 3
.
o -
Fe203
0 -
Fe3C
Fig. 7. X-rays diffraction patterns of MP82/2 powder prepared by aluminolhermic reaction as well as samples sintercd from this powder at increasing time and temperature. Note the formation of FeAI,O, spine1 at very low sintering temperature (1000°C) and its transformation to a-Al,O, for higher sintering temperatures.
4. Discussion
Let us consider a mixture of Fe and Al powders and a given, small quantity of oxygen. Following the procedure used for the constitution of Ellingham diagrams [13], we can calculate a standard Gibbs free energy change associated with the below-mentioned reactions of oxides formation [14]. Note that these calculations must be done for one mole oxygen, in order to make any comparison possible (Table 3). 4/3Al+
O2 -, 2/3Al,O,
1/2Fe + Al + 0, --f 1/2FeAl,O, 4/3Fe + O2 + 2/3Fe,O, These calculations have also been done with real activities of iron and aluminium in Fe-40at.%Al containing 10 wt.% oxygen at 1300 and 1500 K. The activities of aluminium were measured at 1173 K by Eldrige and Komarek [15]. We have first calculated the activity of iron at 1173 K in Fe-40at.%Al using the Gibbs-Duhem equation [13]. Second, the activities nA, and aFe at 1300 and 1500 K were obtained using the Van? Hoff equation [13]; finally the activity of oxygen was assumed to be equal to its concentration. The
K.
Walslci
Table 3 Free enthaipies of formation (kcal mol-’
et al. I Mmerinls
Sciettce
nml EttgineetTng
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(1996)
10%
111
114
oxygen) of cc-Al,O,, FeAI,O, and Fe,O,, calculated for 1 mol oxygen
Unit activities
Real activities
Product
298 K
1000 K
1300 K
1500 K
1300 K
1500 K
2/3Al,O, 1/2FeAl,O, 2/3Fe,O,
-251.8 -222.3 -118.5
-216.6 - 190.0 -89.6
-200.9 - 175.4 -77.6
-191.1 - 165.7 -69.7
- 174.4 - 151.1 -62.0
- 160.4 -131.7 -54.1
numerical values for Fe-40at.%Al are as follows: at 1300 K OAl= 0.0137, are = 0.343; at 1500 K a Al = 0.0141, aFe = 0.344. The free enthalpy changes associated with the above mentioned reactions were calculated using these real activities and are listed in the Table 3 (last two columns). It can be seen from this table that for each temperature and both methods of calculation (elements in their pure states or including real activities), a-Al,O, has the lowest free energy of formation; therefore, at equilibrium, the oxidation products should transform to this phase. This statement clearly indicates that our system is out of equilibrium not only at the end of the MA process but also after low temperature sintering. The following discussion is essentially based on powders prepared by the MA process and has been thought to give some reasons to explain a temporary formation of Fe,O, and FeAl,O, spinel. 4.2. Fe,O, + FeAlz04 t~ansfomation
The powder particles morphology at the intermediate stage of the MA process displays a lamellar microstructure with highly reactive surfaces as schematised in Fig. 8(a). These particles also contain some oxides from initial powders (1.6 wt.% oxygen) and oxidation products from the oxygen initially present in the container (0.8 wt.% oxygen) [12]. Introducing oxygen at this stage (by stopping the milling and opening the container in
fl
.FeAl204
q
FeAl (+FeAl204)
Fig. 8. Morphology of the composite powder particles, processed by MA during 2 hours: a. at the end of the MA process. b. after aluminothermic reaction.
air) leads to the intensive oxidation of iron and aluminium free surfaces with formation of Fe,O, and Y-Al,O, (Fig. 8(b)), the subsequent heat release which raises the temperature above the aluminium melting point (660 “C) and finally the FeAl exothermic formation [16] takes place within the particles together with further oxidation. During this reaction the powder presented red/orange radiation for some seconds. This radiation corresponds to a temperature of the order of 1000 “C. We consider that, in these conditions of temperature, the oxidation in the vicinity of interfaces between iron and aluminium (points of type A - near the surface, with oxygen coming from air during aluminothermic reaction, and type B - inside the particle, with oxygen introduced during the milling or as surface oxides, see Fig. 8(a)) leads to the formation of FeAl,O, spine1 (Fig. 8(b)), effectively detected by X-ray analysis. The formation of the spine1 at 1000 “C occurs at the same time as the disappearance of Fe,O, and FeAl (FT52 in Fig. 7). In order to explain this observation, let us consider the region noted C in Fig. 8(b). This region is enlarged in Fig. 9(a) where a layer of FeAl,O, at the FeAl-Fe,O, interface was added. The formation of this layer, rather than a-A&O, layer, can be explained as follows. Oxygen atoms are strongly fixed within Fe,O, iron oxide at the periphery of the particle, so aluminium atoms have to diffuse from the bulk to the periphery in order to make the transformation of Fe,O, possible. The reaction will therefore take place at the FeAlFe,O, interface, but will be slowed down by the time the aluminium atoms diffuse from the bulk to this interface? which becomes the reaction zone. Note that for a complete transformation of 1 mol of Fe,O, into a-Al,O,, 2 mol of aluminium are needed (reaction (l)), while for a complete transformation of 1 mol of Fe,O, into FeAl,O,, only 1.5 mol of aluminium is needed (reaction (2)). As the reduction of Fe,O, takes place with a deficit of aluminium in the reaction zone, we have to consider this reduction with a small quantity of aluminium atoms available at any given time (we assume that the time of Al diffusion is much longer the time that Al atoms need to react with Fe,O,). First, the calculations were done considering all elements in their pure state (i.e. unit activity); second, they were done by
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Table 4 Free enthalpy changes (AG) associated with the reaction (2) (FeAI,O, formation) and with the reaction (3) indicating more negative AG values for both temperatures and unit or real activities Unit activities
Real activities
Reaction
1300 K
1500 K
1300 K
1500 K
(2)
- 146.6 - 138.7
-144.1 - 136.6
- 133.5 - 126.2
- 129.1 - 122.4
(3)
considering the real Fe and Al activities in Fe4Oat.%Al. The results show that for each quantity of Al, up to 1.5 mol, the formation of FeAl,O, (reaction (2) written for 1.5 mol Al) is more favourable than the partial reduction to cc-Al,O, (reaction (3) written for 1.5 mol Al). The reaction (2) of FeAl,O, formation is characterised by a more negative free enthalpy change than reaction (3), for both methods (with unit activities and with real activities) (Table 4).
Fig. 10. Arrangement of dispersed particles in rings, observed on the polished surface. Sample prepared by MA process and sintered at 1235°C during 20 min. x in Fe-xAl 0,38
Fe,O, + 2Al--$ A&O, + 2Fe
(1)
Fe,O, + 3/2Al+
(2)
3/4FeAl,O& + 5/4Fe
Fe,O, + 3/2Al-+ 3/4Al,O, + 1/4Fe,O, + 3/2Fe
0,36 0,34 0,32
(3)
0,30
This mechanism, based on the rate-limiting diffusion of aluminium from the bulk to the periphery, implies that the spine1 is preferentially formed in the vicinity of the particle surface. As densification by hot pressing does not significantly affect prior powder morphology, the oxides should be arranged in rings on the polished surfaces. This arrangement has been effectively observed for a sample containing approximately 5 wt.% oxygen and sintered at 1235 “C for 20 min (Fig. IO).
0,28
zone of F&204 formation at FeAl I Fe203 interface
FeA1204 s?hel
0,26
/ 0,24 --
I I
0,22 -0,20
116 I
0
0,2
I
I
I I
1
I I
0,4
0,6
0,8
1
62 I 1 1,2
I I 1,4
oxygen[wl%] I I I
1,6
[sl
oxygen weightfor20g
of
starling
(FecAl)
powders
Fig. 11. Calculated Fe-xA1 stoichiometry as a function of oxygen content for the formation of either FeAl,O, or a-Al,O,. The following assumptions have been made: (i) initial composition is Fe-IS, 13g, Al-4, 87g, both surface oxidised powders containing 0.32 g of oxygen (1.6 wt.%), (ii) higher oxygen content comes from “in situ” oxidation, (iii) no carbon is present, (iv) the whole oxygen is present either as FeA120, or a-Al,O,.
4.3. Stability of FeA1204 dispessim
a.
Fe203
q q
FeAl204
b.
q q
FeAl204 Fe (+ FeAl204)
Fe.41 (cFeAl204)
Fig. 9. Enlarged region C from figure 8b. (a) After aluminothermic reaction, the particle is composed mainly from FeAl in the bulk and a layer of Fe,O, at the surface. A thin layer of FeAl,O, starts to be formed, according to the mechanism described in the text. (b) After low temperature sintering, the diffusion of aluminium from the bulk to the surface leads to the transformation of (i) Fe,O, into FeAl,O, and (ii) FeAl into Fe (with small amount of Al and carbon). (+ FeAl,O,) is a fraction of the spine1 previously formed in the bulk (point p in Fig. 8a).
With increasing milling time the fraction of FeAl formed during the MA process by a solid state reaction is increasing [lo], as well as the’homogeneity of oxidation products dispersion (MP88/2). The proportion of sites like A and B in Fig. 8(a) is increasing and there is less and less free iron; therefore, the probability of FeAl,O, formation is increasing. According to Matteazzi et al. [17], who studied the reduction of Fe,03 by pure aluminium, this spine1 can be formed during the milling. We could not confirm this result but have shown (see Sections 3.1 and 3.2) that low temperature sintering results in the formation of FeAl,O, spine1 for both mechanically alloyed or attrition processed pow~
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Table 5 Free enthalpies of FeAl formation as a function of stoichiometry and temperature; the calculations are based on AH and AS values measured either by Radcliff or by Eldridge and were made on two assumptions: (i) AH and AS are constant in the 1000-1500 K temperature range; (ii) 1 mol of Fe,VAl,, = x x inFe + 1’ x V?~, (e.g. 1 mol Fe-40at.%Al= 0.6 x 55.85 g + 0.4 x 26.98 g = 44.3 g). The values of GFe,AI., are presented according to the formula Gi=e,.my= A%,,,, + xGFe -I- JOG*,,adopted by Barin et al. [14] and facilitating thermodynamic calculations Fee enthalpy for FeAl (kcal mol-‘) G(Fe,Al,.) = AG + sG(Fe) + yG(Al) T= 1000 K
T= 1300 K
T= 1500 K
-
16.41 15.93 15.12 14.60
-
21.43 21.00 20.33 19.88
-
25.08 24.12 24.09 23.16
-
16.55 15.96 14.93 14.33
- 21.80 -21.22 - 20.21 - 19.64
-
25.64 25.08 24.08 23.S3
EMridge
Fe-SOat.%Al Fe-40at.%Al Fe-30at.%Al Fe-25at.%Al
ders. At this stage we cannot give any reason to explain why FeAl,O, is formed rather than u-Al,O, (the mechanism described in Section 4.2 for short MA processing, where free iron still exists and can be oxidised into Fe203, does not apply for long MA processing and attrition) and we are forced to accept it as an experimental fact. Taking into account the exceptional stability of FeAl,O, dispersion within FeAl matrix for low temperature sintering, thermodynamic calculations have been done to verify if this system was stable or not. In fact the transformation FeAl,O, --f a-Al,O, takes place within the FeAl matrix and leads to the modification of the FeAl stoichiometry according to the reaction: &Fe-x,Al
+ BFeAl,O, -+ A,Fe-x,Al
+ CA&O3
A207
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(4)
The calculated evolution of Fe-xA1 stoichiometry as a function of oxygen content for the formation of FeA&O, and cc-Al,O, is presented in Fig. 11 and indicates that for each oxygen content the transformation FeAl,O,-+ a-A&O, leads to Fe-xA1 poorer in aluminium (x2
113
105-114
AG,,, -,~* = A,Fe-xZAl - A,Fe-x,Al, due to the modification of Fe-xA1 stoichiometry (x,
Radcll$f
Fe-SOat.%Al Fe-40at.%Al Fe-30at.%Al Fe-2Sat.%Al
and Engineering
of FeAl,O,
into a-A&O,
and its
We conclude that the dispersion of FeA1,04 is metastable and its transformation into a-A&O, is limited by kinetics factors. This means that increasing the sintering time even for T< 1200 “C or high temperature sintering should result in transformation described by the reaction (4). It has been effectively observed in the latter case. No kinetics calculations have been done, but a qualitative explanation of this transformation can be formulated as follows. The transformation of an FeAl,O, particle in an Fe-x,Al matrix into cr-Al,O, requires: the diffusion of Al from the bulk to the particle, the exchange of Al and Fe atoms and finally the crystallographic rearrangement of the atoms. As each of these steps is thermally activated, increasing the sintering time or temperature accelerate the reaction. The coarsening of cc-A&O,’ probably takes place, before this transformation is complete and affects both spine1 and alumina particles. This coarsening has been described by Ostwald ripening kinetics and will be discussed elsewhere with relation to high temperature slow plastic flow properties [19]. 4.5. Comments on y-A120, The formation of this phase by oxidation of free aluminium surfaces during the milling as well as an intermediate step of the transformation of FeAl,O, into m-A&O, are very likely [20]. However, y-A&O, has never been clearly observed, probably because of poor crystallisation and nanometric size. We consider that a fraction of Y-Al,O,, if effectively formed during the milling, will transform into cc-A&O, independently of transformations described above. In our investigations of spinel-alumina transformation, an attempt has been made to identify the step of yA&O, formation by analysis of diffraction patterns on a series of samples sintered at 1050, 1150 and 1240 “C [12], but only weak arguments for its presence could be found.
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5. Conclusions The aim of this study was to analyse the transformations of the oxidation products within FeAl matrix as a function of sintering conditions. These products have been formed by in situ oxidation during either attrition or MA. It has been shown that: l.Low temperature sintering (T < 1200 “C) results in the formation of FeAl,O, spinel. 2. The FeAl,O, formation can be explained by Fe,O, reduction constrained by the rate-limiting diffusion of aluminium to the reaction zone. 3. From thermodynamic point of view, this spine1 is unstable in FeAl matrix. 4. High temperature sintering (T > 1200 “C) results in the transformation of FeAI,O, spine1 to u-A&O,. 5. Further increase in sintering time and temperature leads to the coarsening of this dispersion.
References [l] N.S. Stoloff, iMet. Trans. A., 24 (993) 561. [2] B. Schmidt, P. Nagpal and I. Baker, in CT. Liu, A.I. Taub, N.S. Stoloff and CC. Koch (eds.), High Temperature Ora’ered Intermetallic
Alloys
III,
Proc.
Mater.
Res.
Sot.,
Boston,
USA,
1988,
MRS, Pittsburgh, 1989, pp. 755-760. [3] M.G. Mendiratta, S.K. Ehlers, D.M. Dimiduk, W.R. Kerr, S. Mazdiyasni and H.A. Lipsitt, in N.S. Stoloff, CC. Koch, C.T. Liu and 0. Izumi (eds.), High Temperature Ordered Intermetallic
and Engineering
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(1996)
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14
Alloys II, Proc. Mater. Res. Sot., Bostotr, USA, 1986, MRS, Pittsburgh, 1987, pp. 393-404. [4] I. Jung, M. Rudy and G. Sauthoff, in N.S. Stoloff, CC. Koch, CT. Liu and 0. Izumi (eds.), High Temperntwe Ordered Iutermetallic
[5] [6] [7] [8] [9] [lo] [Ill
illloys
II,
Proc.
Mater.
Res.
Sot.,
Boston,
USA,
lPS6,
MRS, Pittsburgh, 1987, pp, 263-274. M. Rudy and G. Sauthoff, Mater. Sci. Eug., Sl (1986) 525. E. Arzt, Res. Mechanica, 31 (1991) 399. S. Strothers and K. Vedula, Prog. Pol~~tle/~‘~Iet~~ll., 43 (1987) 597. J.S. Benjamin, Metall. Tra,u., 1 (1970) 2943. J.H. Schneibel, P. Grahe and J. RBsler, Mater. Sci. Eug. A, 153 (1992) 684. K. Wolski, G. Le Ca?x, R. Fillit, F. Thbvenot and J. Le Coze, Mater. Sci. Eng. A207 (1996) 97. K. Wolski, F. Thivenot and J. Le Coze, Control of carbides formation in mechanically alloyed FeAl intermetallic, Eur. J. Solid State
Inorg.
Chem.,
32 (1995)
693.
[12] K. Wolski, Influence de la dispersion de phases &ramiques sur la rCsistance au fluage de l’interm&tallique FeAI, Ph.D., Ecole des Mines de St-Etienne, France, 1994. [13] J. Mackowiak, in CR. Tottle and J.C. Wright (eds.), Physical Chemistry for Metallurgists, Allen & Unwin, 1966. [14] I. Barin, 0. Knacke and 0. Kubaschewski (preface), Thermodynamical Properties of Inorganic Substances (1973) old wppicmet~t (1977), Springer, Berlin, 1977. [15] J. Eldridge and K.L. Komarek, Trarrs. A.I.M.E., 230 (1964) 226. 1161 B.H. Rabin and R.N. Wright, Met. Trans. A, 17 (1991) 277. [17] P. Matteazzi and G. Le Cabr, Hpe7~77e Imeractiows, 68 (1991) 177. [18] S.V. Radcliffe, B.L. Averbach and M. Cohen, Acta Metall., 9 (1961) 169. [19] K. Wolski, F. ThCvenot and J. Le Coze, f/~termetnllics, in press.
Materials Scienceand EngineeringA207
(1996)
105-114
Oxide transformations during sintering of in situ oxidised FeAl powders K. Wolski, F. Thivenot, J. Le Coze Cextre
SMS,
Ecole
des Mines
de Saint-Etienne
158, cows
Fawiel,
42023
Saint-Etienne
Cedex,
France
Received17 April 1995;in revisedfonn 10 August 1995
Abstract
Powder metallurgy has been used to prepare FeAl powders with in situ created oxidation products. The aim of this study was to analyse the types of oxide formed as a function of sintering conditions. Highly oxidised powders have been used to give evidence of these transformations by X-ray diffraction. It has been shown that low temperature sintering (T< 1200 “C) results in the formation of FeAl,O, spinel, while higher sintering temperatures lead to its transformation into n-A&O3 and coarsening. An attempt, based on thermodynamic calculations, has been made to explain the formation of FeAl,O, spinel. Keywo~&: Tntertnetallics; FeAl powders; Mechanical alloying; Internal oxidation; Oxide dispersion; Oxide transformation
1. Introduction
FeAl intermetallic has a B2 ordered structure for aluminium content between 32-52 at.%. For ductility purposes, FeAl is considered for applications in the composition range restricted to 32-40 at.% Al [l]. This level of aluminium is still high enough to provide excellent oxidation and corrosion resistance at high temperatures (T > 0.5 r,> [2] and together with relatively low density and high melting temperature (5.6 g cmU3, 1340 “C for Fe-40at.%Al) offers a potential for high temperature applications [3]. The creep strength of FeAl is, however, insufficient [4,5]. One of the most powerful methods of improving creep properties is oxide dispersion strengthening (ODS) [6]. The oxides can be introduced as a powder or formed in situ by oxidation [7]. In both cases ODS FeAl has to be prepared in two steps: (i) ball milling, which is necessary to ensure a homogeneous distribution of oxides or oxidation products and leads to a composite powder; (ii) sintering, which is necessary to ensure densification. During sintering, the oxidation products are trans0921-5093/96/$15.000 1996- ElsevierScienceS.A. All rights reserved
formed into oxides. It is worth noting that the first step, leading to the composite powder, can be performed either by milling prealloyed FeAl powders (attrition) [7] or by milling elemental Fe and Al powders (mechanical alloying, MA) [8,9]. The aim of this study is to clarify the transformations of oxidation products that occur during sintering. The starting materials for this study are composite FeAl powders with dispersion of oxidation products prepared either by attrition or by MA. Owing to the relatively short milling times, the formation of FeAl in composite powders prepared by the MA process is not complete [lo]; however, the homogeneity of oxidation products dispersion is sufficient for the purpose of the present investigation. Two aspects of this study have already been published: (i) formation of FeAl by MA process [lo] and (ii) formation of carbides resulting from the use of a process control agent [ll]; consequently these aspects will not be discussed in this paper. This investigation is a part of a wider study aimed at improving high temperature slow plastic flow properties of FeAl intermetallic, by dispersion of nanosized oxide particles [12].
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Table 1 Elemental powders used in this study Power
Fe Al FeAl
Producer
Prolabo Koch Osprey
Diameter (pm)
10 8-15 5-50
Oxygen content (wt.%)
Carbon content (ppm)
certified
measured
certified
measured
0.2 1.0 300 ppm
0.5 5.0 3500 ppm
200
140 400 30
2. Experimental procedure 2.1. Powders
Prealloyed FeAl powder containing 0.35 wt.% oxygen has been used for attrition. Very fine (therefore highly oxidised) powders of elemental Fe and Al were used to prepare the composite powder by the mechanical alloying process (Table 1). In some cases 2 wt.% stearic acid (CH,(CH,),,COOH, PM = 284 g) was added to prevent elemental powders from sticking either to the balls or to the container wall [ll]. 2.2. Processing Attrition of FeAl powder was performed in a vertical mill (attritor of Szegvari type) installed in a glove box under flowing argon containing 100 ppm of oxygen, while mechanical alloying (MA) was carried out in a planetary mill. In the latter the containers were filled under air and hermetically closed during milling. In both cases hardened steel containers and balls (lOOC6 steel, 2 mm diameter for attrition and 4 mm diameter for MA; 50:1-balls to powder ratio) were used. The resulting powders were sintered in a hot uniaxial press (40-80 MPa, under argon). Milling and sintering parameters of all powders used in this study, together with initial
125
Method of processing MA MA Attrition
compositions, are listed in Table 2. Preparing the composite powder by attrition results in significantly lower oxygen content compared with the MA process. This is due to cleaner initial powders and processing under argon. However, significantly higher oxygen content was necessary (MA process) to analyse the transformations of oxidation products during sintering. The analysis of Fe,AlC,,, carbide formation, resulting from the addition of stearic acid as process control agent was described elsewhere [II]. One of the methods of carbide elimination, namely suppressing stearic acid and processing in several steps [l l] was applied to prepare MP94 powder. In the present study the MP84/3 powder was processed in six steps in order to increase the oxygen content so that X-rays could be used to analyse the oxide formation. 2.3. Methods of nnalysis
Phase transformations were studied with X-ray diffraction (XRD) using Co KU. radiation from a Philips diffractometer (1730/10) equipped with Dosophatex system. Microstructure was characterised by scanning electron microscopy (SEM), equipped with standard microanalyser (JEOL JSM840 + Tracer). Additional information has been obtained from transmission electron microscopy (TEM).
Table 2 Processing parameters of powders and samples prepared and analysed in this study Powder symbol
Processing method
Initial composition
IMilling time and velocity
Oxygen (wt.%)
Sintering T (“C)/t (min)
Sample symbol
A38
Attrition
FeAl-20
1.5
1170160
FAG5
MP88j2
MA
5
1200140
FT45
MP84/3
MA
Fe-15.13 g Al-4.87 g sa-0.4 g
10
1270180
FT39
Fe-15.0 g Al-5.0 g (without sa)
Attrition 15 h 420 rev min-’ MA 3 h 200 rev min-’ MA 6 x 0.5 h 200 rev min MA 6hf3x2h 200 rev min-’
Fe-15.13 g Al-4.87 g sa-0.4 g
MA 1.5 h +0.5 h 200 rev min-’
1050/30 1150/30 1240/30 1240/240 lOOOj5 1100/60 1300/60
FTl7 FT78 FT79 FTSO FT52 FT50 FT34
MP94
MA
MA+
MP82/2
sa. stearic acid
aluminothermic reaction
g
-1
10
25
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3. Results 3.1. Attrition
Fie 1. General microstructure of FA65 sample sintered at 1170°C frz; n prealloyed and attrited FeAl powder.
General microstructure of composites prepared by low temperature sintering (T< 1200 “C) of the composite powder obtained by attrition of prealloyed FeAl powders is presented in Fig. 1. It consists of nanosized grains (50 to 100 nm) dispersed in the FeAl grains (2 to 5 pm). This dispersion (Fig. 2(a)) produces concentric rings on the diffraction pattern (Fig. 2(b)), which were indexed as corresponding to the FeA&O, spine1 phase. (Additional high intensity spots correspond to the diffraction from some FeAl grains, which were selected at the same time as numerous spine1 grains.) Fig. 2(c) is a ring pattern for a nanosized dispersion of this spinel. Calculations based on a final oxygen content of 1.45
(4
(b)
Fig. 2. Analysis of the dispersed phase crystallography by TEM (FA65-type sample). a. bright field image of the dispersed phase. b. corresponding diffraction patterns. c. theoretical disposition of rings for FeA120, spinel.
108
K. Woiski et al. 1 Materials Science and Engineering r1207 (1996) 1OS 114
wt.% and assumption of the formation of FeAl,O, spine1 from the whole oxygen gives 5.4 vol.% of spine1 dispersed in the Fe-39at.%Al matrix. As low temperature sintering (T < 1200 “C) was sufficient for a complete densification of this composite, high temperature sintering was not investigated in this case. Note that T = 1200 “C is somewhat arbitrary. In fact the transformations discussed in this paper are thermally activated, so depend also on sintering time. Let us define, for further discussion, low temperature sintering as T < 1200 “C and short sintering time (t < 60 min) and high temperature sintering as T> 1200 “C and long sintering time (t > 60 min). 3.2. Mechanical alloying
The composite powders prepared by MA are characterised by a high oxygen content, within the range 5 to 10 wt.%, depending on oxidation conditions. Higher oxygen content, compared with powders prepared by attrition, leads to the necessity of increasing the sintering temperature for complete densification. The microstructures of two samples (FT45 in Fig. 3(a) and FT39 in Fig. 3(b)) are typical of those observed at their respective sintering conditions. FT45 was prepared by one-step milling for 3 h (5 wt.% oxygen), followed by low temperature sintering (1200 “C, 40 min). FT39 was prepared by six-step milling with intermediate sievings (6 x 0.5 = 3 h, IO wt.% oxygen), followed by high temperature sintering (1270 “C, 80 min). The comparison of microstructures (Figs. 3(a) and 3(b)) gives an immediate indication of the strong coarsening of the dispersed particles (dark phase). XRD has identified the coarsened particles in FT39 as a-alumina. However, no information could be obtained for particles in the FT45 sample, probably because of the nanometre size and possibly incomplete crystallisation.
(a)
H
I pm
Consequently, thin foils have been prepared and their analysis has indicated the presence of FeAl,O, spine1 in FT45 (Figs. 4(a) and 4(b)), at least for zones containing uncoarsened particles, as in the left bottom corner of the micrograph in Fig. 4(a) (diffraction patterns from Fig. 4(b) were obtained from this area). Note that diffraction spots corresponding to FeAl are not observed because there was only one FeAl grain in the analysed area and this grain obviously has an orientation out of diffraction conditions. Additional analysis of FT39 sample (Figs. 5(a) and 5(b)) has confirmed the presence of cc-alumina (Fig. 5(c)), as indicated by a coherent indexing of the diffraction pattern (Fig. 5(d)). As no information could be obtained by X-ray analysis on the FT45 sample containing 5 wt.% oxygen and the TEM analysis cannot provide the global information about phase constitution, an additional charge of powder with 10 wt.% oxygen has been prepared by a six-step milling process (MP94) and analysed as a function of increasing sintering time and temperature (Fig. 6). The XRD patterns from this figure again confirm that high temperature sintering leads to the formation of a-alumina. Unfortunately, neither the nature of oxidation products nor any transformation leading to a-alumina could be detected for this already high oxygen level. 3.3. Strong exothemic
reaction
The problem of the detection limit of X-ray analysis has been overcome by further increasing the oxidation products by means of strong exothermic reaction. This reaction takes place immediately after opening a container, if MA is interrupted after 1.5 to 2 h of milling. In fact, for these milling times, Fe and Al powders have a lamellar morphology with unoxidised surfaces of high specific area, so are extremely reactive. The reason for
(b)
i-i
I/m
Fig. 3. General microstructure of samples sintered from powders prepared by MA process. a. FT45 sample containing 5 wt% oxygen and sintered at 1200°C during 40 min. b. FT39 sample containing 10 wt% oxygen and sintered at 1270°C during 80 min, note the significant coarsening of dispersed phase in the latter.
K. Wolski
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109
(b)
(a)
Fig. 4. TEM analysis of the dispersed phase in materials prepared by mechanical alloying and low temperature sintering (FT 45-type sample). a. bright field image showing the dispersed phase. b. diffraction patterns, compare with the theoretical disposition of rings for FeAI,O, spine& presented in figure 2~.
(a)
H
loonm
(b)
H
IOOnm
-2:21= l
3 ;4205 ,
. 10 .
0
116 .
40
2212 .
(d) Fig. 5. TEM analysis of the dispersed phase in materials prepared by mechanical alloying and high temperature sintering (FT39-type sample). a. dark field image. b. bright field image. c. diffraction patterns. d. coherent indexing of spots confirming the presence of a-A&O,.
K. Wolski et al. / Materials Science atd Dlgitzeerirzg AZ07 (1996) 105-l 14
I I / 35
30 .
FeAl
A
a-A1203
q
Fe$AIC
FT80 FT79 FT78 FT77
-
- (loo),
'1
e I 40
I
superstructure
b I 45
/
peak
"
I 50
c 11
I
j /. .. ... .. . . . .. .. .j . .. ... .. ... .. ... .. ..i .. ... .. ... .. ... .. .. i .. .. .. ... . . .. .. ...
FT34
j
55
28 [WI : 1iOOWlh
25
x
124O"C, 124O"C, 115O"C, 105O"C,
FT50 : l{OOWlh ... .. ... ... .. .. ...l....................
40MPa, 4OMPa, GOMPa, 80MPa,
240mn. 30mn. 30mn. 30mn.
Fig. 6. Evolution of X-rays diffraction pattern as a function of sintering temperature. Samples FT77-FT80 have been sintered from the same MP94 powder containing 10 wt.% oxygen. Note that any dispersion could be detected for lower sintering temperatures (1050 and 1150°C).
performing this experiment was only to give proof for spine1 formation using XRD and not to prepare a viable material. The XRD patterns of MP82/2 powder, prepared as described above, as well ai of three samples sintered from this powder at increasing time and temperature are presented in Fig. 7. This powder contains free iron, a small amount of FeAl, iron oxide (Fe,O,) and spine1 (FeAl,O,). Note that at this stage aluminium seems to be transformed into FeAl and spine1 but there is still no evidence of the presence of a-alumina. Very short heat treatment of this powder at 1000 “C leads to the decomposition of both iron oxide and FeAl intermetallic and formation of FeAl,O, spine1 in the Fe matrix (Fig. 7, FT52), which additionally contains carbon atoms from stearic acid (2 wt.% stearic acid can result in up to 1.4 wt.% carbon). With further increase in sintering time and temperature (1 h, 1100 “C) the spine1 is progressively transformed into cc-alumina (Fig. 7, FT50) and finally after 1 h at 1200 “C this transformation is almost complete (Fig. 7, FT34). At the same time we observe the formation of Fe& carbide. We consider that this change of composition from FeAl intermetallic to Fe-Fe& cermet (Fig. 7, FT34), due to the presence of stearic acid in the starting powders and complete oxidation of aluminium, does not affect the formation of the spine1 and its transformation into a-alumina.
30
35
45
40
50
55
28 Idegl ' m-
-
FeAl20 4 a.Al
0 2 3
.
o -
Fe203
0 -
Fe3C
Fig. 7. X-rays diffraction patterns of MP82/2 powder prepared by aluminolhermic reaction as well as samples sintercd from this powder at increasing time and temperature. Note the formation of FeAI,O, spine1 at very low sintering temperature (1000°C) and its transformation to a-Al,O, for higher sintering temperatures.
4. Discussion
Let us consider a mixture of Fe and Al powders and a given, small quantity of oxygen. Following the procedure used for the constitution of Ellingham diagrams [13], we can calculate a standard Gibbs free energy change associated with the below-mentioned reactions of oxides formation [14]. Note that these calculations must be done for one mole oxygen, in order to make any comparison possible (Table 3). 4/3Al+
O2 -, 2/3Al,O,
1/2Fe + Al + 0, --f 1/2FeAl,O, 4/3Fe + O2 + 2/3Fe,O, These calculations have also been done with real activities of iron and aluminium in Fe-40at.%Al containing 10 wt.% oxygen at 1300 and 1500 K. The activities of aluminium were measured at 1173 K by Eldrige and Komarek [15]. We have first calculated the activity of iron at 1173 K in Fe-40at.%Al using the Gibbs-Duhem equation [13]. Second, the activities nA, and aFe at 1300 and 1500 K were obtained using the Van? Hoff equation [13]; finally the activity of oxygen was assumed to be equal to its concentration. The
K.
Walslci
Table 3 Free enthaipies of formation (kcal mol-’
et al. I Mmerinls
Sciettce
nml EttgineetTng
A207
(1996)
10%
111
114
oxygen) of cc-Al,O,, FeAI,O, and Fe,O,, calculated for 1 mol oxygen
Unit activities
Real activities
Product
298 K
1000 K
1300 K
1500 K
1300 K
1500 K
2/3Al,O, 1/2FeAl,O, 2/3Fe,O,
-251.8 -222.3 -118.5
-216.6 - 190.0 -89.6
-200.9 - 175.4 -77.6
-191.1 - 165.7 -69.7
- 174.4 - 151.1 -62.0
- 160.4 -131.7 -54.1
numerical values for Fe-40at.%Al are as follows: at 1300 K OAl= 0.0137, are = 0.343; at 1500 K a Al = 0.0141, aFe = 0.344. The free enthalpy changes associated with the above mentioned reactions were calculated using these real activities and are listed in the Table 3 (last two columns). It can be seen from this table that for each temperature and both methods of calculation (elements in their pure states or including real activities), a-Al,O, has the lowest free energy of formation; therefore, at equilibrium, the oxidation products should transform to this phase. This statement clearly indicates that our system is out of equilibrium not only at the end of the MA process but also after low temperature sintering. The following discussion is essentially based on powders prepared by the MA process and has been thought to give some reasons to explain a temporary formation of Fe,O, and FeAl,O, spinel. 4.2. Fe,O, + FeAlz04 t~ansfomation
The powder particles morphology at the intermediate stage of the MA process displays a lamellar microstructure with highly reactive surfaces as schematised in Fig. 8(a). These particles also contain some oxides from initial powders (1.6 wt.% oxygen) and oxidation products from the oxygen initially present in the container (0.8 wt.% oxygen) [12]. Introducing oxygen at this stage (by stopping the milling and opening the container in
fl
.FeAl204
q
FeAl (+FeAl204)
Fig. 8. Morphology of the composite powder particles, processed by MA during 2 hours: a. at the end of the MA process. b. after aluminothermic reaction.
air) leads to the intensive oxidation of iron and aluminium free surfaces with formation of Fe,O, and Y-Al,O, (Fig. 8(b)), the subsequent heat release which raises the temperature above the aluminium melting point (660 “C) and finally the FeAl exothermic formation [16] takes place within the particles together with further oxidation. During this reaction the powder presented red/orange radiation for some seconds. This radiation corresponds to a temperature of the order of 1000 “C. We consider that, in these conditions of temperature, the oxidation in the vicinity of interfaces between iron and aluminium (points of type A - near the surface, with oxygen coming from air during aluminothermic reaction, and type B - inside the particle, with oxygen introduced during the milling or as surface oxides, see Fig. 8(a)) leads to the formation of FeAl,O, spine1 (Fig. 8(b)), effectively detected by X-ray analysis. The formation of the spine1 at 1000 “C occurs at the same time as the disappearance of Fe,O, and FeAl (FT52 in Fig. 7). In order to explain this observation, let us consider the region noted C in Fig. 8(b). This region is enlarged in Fig. 9(a) where a layer of FeAl,O, at the FeAl-Fe,O, interface was added. The formation of this layer, rather than a-A&O, layer, can be explained as follows. Oxygen atoms are strongly fixed within Fe,O, iron oxide at the periphery of the particle, so aluminium atoms have to diffuse from the bulk to the periphery in order to make the transformation of Fe,O, possible. The reaction will therefore take place at the FeAlFe,O, interface, but will be slowed down by the time the aluminium atoms diffuse from the bulk to this interface? which becomes the reaction zone. Note that for a complete transformation of 1 mol of Fe,O, into a-Al,O,, 2 mol of aluminium are needed (reaction (l)), while for a complete transformation of 1 mol of Fe,O, into FeAl,O,, only 1.5 mol of aluminium is needed (reaction (2)). As the reduction of Fe,O, takes place with a deficit of aluminium in the reaction zone, we have to consider this reduction with a small quantity of aluminium atoms available at any given time (we assume that the time of Al diffusion is much longer the time that Al atoms need to react with Fe,O,). First, the calculations were done considering all elements in their pure state (i.e. unit activity); second, they were done by
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Table 4 Free enthalpy changes (AG) associated with the reaction (2) (FeAI,O, formation) and with the reaction (3) indicating more negative AG values for both temperatures and unit or real activities Unit activities
Real activities
Reaction
1300 K
1500 K
1300 K
1500 K
(2)
- 146.6 - 138.7
-144.1 - 136.6
- 133.5 - 126.2
- 129.1 - 122.4
(3)
considering the real Fe and Al activities in Fe4Oat.%Al. The results show that for each quantity of Al, up to 1.5 mol, the formation of FeAl,O, (reaction (2) written for 1.5 mol Al) is more favourable than the partial reduction to cc-Al,O, (reaction (3) written for 1.5 mol Al). The reaction (2) of FeAl,O, formation is characterised by a more negative free enthalpy change than reaction (3), for both methods (with unit activities and with real activities) (Table 4).
Fig. 10. Arrangement of dispersed particles in rings, observed on the polished surface. Sample prepared by MA process and sintered at 1235°C during 20 min. x in Fe-xAl 0,38
Fe,O, + 2Al--$ A&O, + 2Fe
(1)
Fe,O, + 3/2Al+
(2)
3/4FeAl,O& + 5/4Fe
Fe,O, + 3/2Al-+ 3/4Al,O, + 1/4Fe,O, + 3/2Fe
0,36 0,34 0,32
(3)
0,30
This mechanism, based on the rate-limiting diffusion of aluminium from the bulk to the periphery, implies that the spine1 is preferentially formed in the vicinity of the particle surface. As densification by hot pressing does not significantly affect prior powder morphology, the oxides should be arranged in rings on the polished surfaces. This arrangement has been effectively observed for a sample containing approximately 5 wt.% oxygen and sintered at 1235 “C for 20 min (Fig. IO).
0,28
zone of F&204 formation at FeAl I Fe203 interface
FeA1204 s?hel
0,26
/ 0,24 --
I I
0,22 -0,20
116 I
0
0,2
I
I
I I
1
I I
0,4
0,6
0,8
1
62 I 1 1,2
I I 1,4
oxygen[wl%] I I I
1,6
[sl
oxygen weightfor20g
of
starling
(FecAl)
powders
Fig. 11. Calculated Fe-xA1 stoichiometry as a function of oxygen content for the formation of either FeAl,O, or a-Al,O,. The following assumptions have been made: (i) initial composition is Fe-IS, 13g, Al-4, 87g, both surface oxidised powders containing 0.32 g of oxygen (1.6 wt.%), (ii) higher oxygen content comes from “in situ” oxidation, (iii) no carbon is present, (iv) the whole oxygen is present either as FeA120, or a-Al,O,.
4.3. Stability of FeA1204 dispessim
a.
Fe203
q q
FeAl204
b.
q q
FeAl204 Fe (+ FeAl204)
Fe.41 (cFeAl204)
Fig. 9. Enlarged region C from figure 8b. (a) After aluminothermic reaction, the particle is composed mainly from FeAl in the bulk and a layer of Fe,O, at the surface. A thin layer of FeAl,O, starts to be formed, according to the mechanism described in the text. (b) After low temperature sintering, the diffusion of aluminium from the bulk to the surface leads to the transformation of (i) Fe,O, into FeAl,O, and (ii) FeAl into Fe (with small amount of Al and carbon). (+ FeAl,O,) is a fraction of the spine1 previously formed in the bulk (point p in Fig. 8a).
With increasing milling time the fraction of FeAl formed during the MA process by a solid state reaction is increasing [lo], as well as the’homogeneity of oxidation products dispersion (MP88/2). The proportion of sites like A and B in Fig. 8(a) is increasing and there is less and less free iron; therefore, the probability of FeAl,O, formation is increasing. According to Matteazzi et al. [17], who studied the reduction of Fe,03 by pure aluminium, this spine1 can be formed during the milling. We could not confirm this result but have shown (see Sections 3.1 and 3.2) that low temperature sintering results in the formation of FeAl,O, spine1 for both mechanically alloyed or attrition processed pow~
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Science
Table 5 Free enthalpies of FeAl formation as a function of stoichiometry and temperature; the calculations are based on AH and AS values measured either by Radcliff or by Eldridge and were made on two assumptions: (i) AH and AS are constant in the 1000-1500 K temperature range; (ii) 1 mol of Fe,VAl,, = x x inFe + 1’ x V?~, (e.g. 1 mol Fe-40at.%Al= 0.6 x 55.85 g + 0.4 x 26.98 g = 44.3 g). The values of GFe,AI., are presented according to the formula Gi=e,.my= A%,,,, + xGFe -I- JOG*,,adopted by Barin et al. [14] and facilitating thermodynamic calculations Fee enthalpy for FeAl (kcal mol-‘) G(Fe,Al,.) = AG + sG(Fe) + yG(Al) T= 1000 K
T= 1300 K
T= 1500 K
-
16.41 15.93 15.12 14.60
-
21.43 21.00 20.33 19.88
-
25.08 24.12 24.09 23.16
-
16.55 15.96 14.93 14.33
- 21.80 -21.22 - 20.21 - 19.64
-
25.64 25.08 24.08 23.S3
EMridge
Fe-SOat.%Al Fe-40at.%Al Fe-30at.%Al Fe-25at.%Al
ders. At this stage we cannot give any reason to explain why FeAl,O, is formed rather than u-Al,O, (the mechanism described in Section 4.2 for short MA processing, where free iron still exists and can be oxidised into Fe203, does not apply for long MA processing and attrition) and we are forced to accept it as an experimental fact. Taking into account the exceptional stability of FeAl,O, dispersion within FeAl matrix for low temperature sintering, thermodynamic calculations have been done to verify if this system was stable or not. In fact the transformation FeAl,O, --f a-Al,O, takes place within the FeAl matrix and leads to the modification of the FeAl stoichiometry according to the reaction: &Fe-x,Al
+ BFeAl,O, -+ A,Fe-x,Al
+ CA&O3
A207
(1996)
(4)
The calculated evolution of Fe-xA1 stoichiometry as a function of oxygen content for the formation of FeA&O, and cc-Al,O, is presented in Fig. 11 and indicates that for each oxygen content the transformation FeAl,O,-+ a-A&O, leads to Fe-xA1 poorer in aluminium (x2
113
105-114
AG,,, -,~* = A,Fe-xZAl - A,Fe-x,Al, due to the modification of Fe-xA1 stoichiometry (x,
Radcll$f
Fe-SOat.%Al Fe-40at.%Al Fe-30at.%Al Fe-2Sat.%Al
and Engineering
of FeAl,O,
into a-A&O,
and its
We conclude that the dispersion of FeA1,04 is metastable and its transformation into a-A&O, is limited by kinetics factors. This means that increasing the sintering time even for T< 1200 “C or high temperature sintering should result in transformation described by the reaction (4). It has been effectively observed in the latter case. No kinetics calculations have been done, but a qualitative explanation of this transformation can be formulated as follows. The transformation of an FeAl,O, particle in an Fe-x,Al matrix into cr-Al,O, requires: the diffusion of Al from the bulk to the particle, the exchange of Al and Fe atoms and finally the crystallographic rearrangement of the atoms. As each of these steps is thermally activated, increasing the sintering time or temperature accelerate the reaction. The coarsening of cc-A&O,’ probably takes place, before this transformation is complete and affects both spine1 and alumina particles. This coarsening has been described by Ostwald ripening kinetics and will be discussed elsewhere with relation to high temperature slow plastic flow properties [19]. 4.5. Comments on y-A120, The formation of this phase by oxidation of free aluminium surfaces during the milling as well as an intermediate step of the transformation of FeAl,O, into m-A&O, are very likely [20]. However, y-A&O, has never been clearly observed, probably because of poor crystallisation and nanometric size. We consider that a fraction of Y-Al,O,, if effectively formed during the milling, will transform into cc-A&O, independently of transformations described above. In our investigations of spinel-alumina transformation, an attempt has been made to identify the step of yA&O, formation by analysis of diffraction patterns on a series of samples sintered at 1050, 1150 and 1240 “C [12], but only weak arguments for its presence could be found.
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5. Conclusions The aim of this study was to analyse the transformations of the oxidation products within FeAl matrix as a function of sintering conditions. These products have been formed by in situ oxidation during either attrition or MA. It has been shown that: l.Low temperature sintering (T < 1200 “C) results in the formation of FeAl,O, spinel. 2. The FeAl,O, formation can be explained by Fe,O, reduction constrained by the rate-limiting diffusion of aluminium to the reaction zone. 3. From thermodynamic point of view, this spine1 is unstable in FeAl matrix. 4. High temperature sintering (T > 1200 “C) results in the transformation of FeAI,O, spine1 to u-A&O,. 5. Further increase in sintering time and temperature leads to the coarsening of this dispersion.
References [l] N.S. Stoloff, iMet. Trans. A., 24 (993) 561. [2] B. Schmidt, P. Nagpal and I. Baker, in CT. Liu, A.I. Taub, N.S. Stoloff and CC. Koch (eds.), High Temperature Ora’ered Intermetallic
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[12] K. Wolski, Influence de la dispersion de phases &ramiques sur la rCsistance au fluage de l’interm&tallique FeAI, Ph.D., Ecole des Mines de St-Etienne, France, 1994. [13] J. Mackowiak, in CR. Tottle and J.C. Wright (eds.), Physical Chemistry for Metallurgists, Allen & Unwin, 1966. [14] I. Barin, 0. Knacke and 0. Kubaschewski (preface), Thermodynamical Properties of Inorganic Substances (1973) old wppicmet~t (1977), Springer, Berlin, 1977. [15] J. Eldridge and K.L. Komarek, Trarrs. A.I.M.E., 230 (1964) 226. 1161 B.H. Rabin and R.N. Wright, Met. Trans. A, 17 (1991) 277. [17] P. Matteazzi and G. Le Cabr, Hpe7~77e Imeractiows, 68 (1991) 177. [18] S.V. Radcliffe, B.L. Averbach and M. Cohen, Acta Metall., 9 (1961) 169. [19] K. Wolski, F. ThCvenot and J. Le Coze, f/~termetnllics, in press.