Mechanical properties and microstructure of AlFeX alloys

Materials Science and Engineering, A 186 (1994) 55-64

55

Mechanical properties and microstructure of A1-Fe-X alloys J. C. E h r s t r 6 m Pechiney Centre de Recherches de Voreppe, BP 27, 38340 Voreppe (France)

A. Pineau Ecole des Mines de Paris, Centre des Mat&iaux, BP 87, 91003 Evry Cedex (France) (Received April 12, 1993; in revised form November 15, 1993)

Abstract The mechanical behavior and microstructure of aluminum alloys for high temperature applications, prepared by rapid solidification and powder metallurgy techniques, have been observed. These alloys contain high concentrations of transition elements which form thermally stable precipitates. Tensile tests were performed at room temperature, at 77 K and at high temperature. The microstructure is described in terms of precipitation and dislocation structure. Samples deformed at different temperatures were observed. The behavior of alloys of this type is discussed, introducing the effect of an internal stress. Special attention is devoted to the ductility dip frequently observed in such alloys.

1. Introduction

The development of rapid solidification technologies and powder metallurgy in the past ten or twenty years has enabled metallurgists to produce new kinds of alloys having a variety of until recently unexpected properties. Among these new materials, the Al-Fe-X alloys (where X stands for Zr, Mo, Si, etc.) combine very high strength at room temperature and good mechanical properties up to about 575 K with lightness and better machinability than, for example, titanium alloys. The hardening mechanism of this family of alloys is based upon the introduction of a high volume fraction of finely distributed intermetallic precipitates which are stable at relatively high temperature, owing to the low diffusivity and solubility of the alloying elements. In addition, these materials contain fine subgrains pinned by intermetallic precipitates. The processing route of these alloys is designed to avoid coarsening of the microstructure which could be deleterious to strength and ductility. However, the question arises of the suitability of fine microstructures for high temperature applications. In particular, it has previously been shown that A1-Fe based alloys generally suffer from a decreased ductility or "ductility dip" at 500-550 K [1-10]. Several mechanisms have been invoked to explain this reduced ductility [1, 2, 10]. In the present paper we try to show that these fine microstructures result in a low work hardening rate and a low apparent strain rate sensitivity for this type of 0921-5093/94/$7.00 SSDI 0921-5093(93)09484-Z

alloy. Comparison, using transmission electron microscopy (TEM), of microstructures produced by deformation at different temperatures supports this hypothesis, as well as macroscopic arguments related to the constitutive equations. It will be seen that particular features of the constitutive law provide an alternative explanation for the ductility dip, besides the hypothesis of dynamic strain ageing put forward by Skinner and co-workers [7, 10] and Bouchaud et al. [1].

2. Experimental details

In addition to binary A1-Fe alloys, A I - F e - M o - Z r and A1-Fe-V-Si alloys were investigated. Their composition is given in Table 1. These materials were supplied by Pechiney in the form of extruded bars. The processing route was inert gas or air atomization, canning, degassing and extrusion (Table 1). In the case of the binary alloys, particles having two different average diameters were extruded separately. Tensile and creep tests were performed between 77 and 625 K, using liquid nitrogen, resistance and radiation furnaces. The specimens were held for 15 rain at test temperature before the tests were run. Thin foils for TEM examination were prepared by mechanical polishing down to 150/~m and further electrochemical polishing at - 1 8 °C in a solution of 700 ml ethanol, 1 2 0 m l distilled water, 100 ml 2butoxyethanol and 80 ml HCIO4 (60%). Tensile specimens were quenched within 30 s after unloading and © 1994 - Elsevier Sequoia. All rights reserved

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J. C. Ehrstrrm, A. Pineau

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Mechanical properties and microstructure of Al-Fe-X alloys

TABLE 1. Composition (wt.%) and processing Alloy

Fe (wt.%)

AI-Fe

8.00 8.00

AI-Fe-Mo-Zr

8.00 7.70

A1-Fe-V-Si

Mo (wt.%)

2.10 1.95

8.90

Zr (wt.%)

1.00 0.90

Si (wt.%)

V (wt.%)

0.31 2.10

1.50

Atomization

Particulate size (Fisher)

Extrusion

Air Air

15/zm 40-400 ktm

11 11

Argon Argon

45 ~m 63/zm

6.25 9

Nitrogen

63/~m

9

Fig. 1. Microstructure of the A1-Fe alloy ( 15/~).

kept in liquid nitrogen before observation. Subgrain size measurements were made in the following manner: (1) selection of a (111) diffraction spot; (2) scanning of the foil and taking of micrographs of all illuminated grains in dark field conditions, sometimes after a small adjustment of the foil position; (3) cutting up of the subgrain image and measurement of its surface. Identification of the precipitates was performed by X-ray diffraction, TEM diffraction and semi-quantitative STEM (Philips EM430 microscope) analysis.

Fig. 2. Subgrains in a prior particle. AI-Fe alloy (15/~), LT direction, dark field.

3. Results

3.1. Microstructure before deformation As one can see in Fig. 1, the microstructure of the A1-Fe alloy is marked by elongation of prior powder particles in the extrusion direction. These particles are subdivided into subgrains, the subgrain boundaries being pinned by precipitates (Figs. 2 and 3). Both binary alloys have about the same microstructure, the 40-400/zm alloy showing, as expected, a somewhat coarser precipitation. In the case of the 15/zm alloy, quantitative evaluation of the subgrain size was done, the result of which is given in Fig. 4. The precipitates are described in Table 2. No preferential crystallographic relationship between matrix and precipitates was observed. X-ray diffraction pole figures show a strong (111) texture in this binary alloy [11].

§ " 111 Fig. 3. Subgrain boundaries pinned by precipitates. AI-Fe alloy ( 15/z) as extruded.

J. c. Ehrstrdm, A. Pineau

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Mechanical properties and microstructure of Al-Fe-X allovs

The microstructure of the A1-Fe-V-Si alloy is rather similar to that of the AI-Fe alloy. The main difference is that some grains having a very fine precipitation show a less recovered dislocation microstructure. Precipitates consist mainly of the cubic a-AIFeSi phase (lm3, 138 atoms by unit cell, a = 1.26 nm) (Table 3). This metastable phase is reported to be stabilized by the presence of vanadium [15]. No orientation of these precipitates with respect to the surrounding matrix was observed by TEM, although Skinner has reported that in the same alloy obtained by meltspinning, the aA1FeSi phase is coherent with the matrix [16]. The A1-Fe and A1-Fe-V-Si alloys have a relatively homogeneous microstructure, whereas the A1-Fe-

II

I

INITIAL [ ]

173 K

[] 525 K I

M o - Z r alloy contains particles of a zone-A type microstructure [17] which are not deformed at all or are broken into fragments. The volume fraction of this kind of particle has been estimated as more than 5%. They correspond to more rapidly solidified particles and contain mainly a cellular phase called S, and/or stable AI3Fe, as well as a relatively high volume fraction of the coherent or semi-coherent A13Zr metastable phase (Pm3m, a = 0.405 nm) (Fig. 5). This phase is also present in the normal elongated particles but inhomogeneously distributed. Coming back to the S phase, it is shown in Table 3 that this has a crystallographic structure related to the structure of the icosahedral quasicrystalline phase observed by Shechtman and Blech [14], which is itself related to the a-A1FeSi cubic phase. In the A 1 - F e - M o - Z r alloy the S phase is microcrystalline.

I

number of s.-gr. 18.

3.2. M e c h a n i c a l properties

16' 14' 12' 10' 8. 6' 4' 2" 0 0,2

57

0,4

0,6 0,8 sub-grain diameter (.u.m)

1

1,2

Fig. 4. Subgrain size repartition, in the as-extruded state, after deformation at 525 K and 173 K. In the latter case, not only subgrains but also dislocations cells are taken into account.

The yield stress (YS) and the ultimate tensile strength (UTS) of the four alloys are plotted in Fig. 6. At room temperature the A I - F e - M o - Z r alloy has a very high strength, almost comparable with that of 2XXX and 7XXX series alloys. At high temperature, the difference between the four alloys reduces and the average strength is much higher than the strength of conventional aluminum alloys. The ductility, measured by the area reduction at fracture, is shown in Fig. 7 is a function of temperature ( T ) and strain rate g for the binary alloy. Data concerning the quaternary alloys are given in Table 4. It appears that the ductility dip is not always observed. In

TABLE 2. Precipitates present in the AI-Fe-X alloys Alloy

AI-Fe

Particulate type

Deformed

Precipitates

Crystallographic structure

Shape/aspect

AI, Fe

Orthorhombic

50

About 9% ( < 21%)

AI, Fe

Monoclinic

Globular Rods, trapezoid, twinned Globular

50-500

About 9% ( < 13%)

50

?

Globular Globular, (semi-)coherent Globular

50-500 5-20

21-23% Heterogeneous average < 1%

Other phases, incl. AI-Fe, Si AI (Fe, Mo) AI, Zr

Orthorhombic Cubic

Deformed Other A1-Fe, Mo, Zr phases

?

S

Icosahedral?

A type

A1, Fe AI, Zr

Deformed

AI-Fe-V-Si ?

AI-Fe-Mo-Zr

AI-Fe-V-Si

Size(nm)

Estimated vol. fraction

15-20 About 5%

Monoclinic Cubic

Microcrystalline, 200 nm cellular cells Orientated rods 025 x 200 Globular, 5-20 (semi-)coherent

Cubic Cubic ?

Globular 20-200 Microcrystalline 500-1000

About 20% Small

* About 1%

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J.C. Ehrstrrm, A. Pineau

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Mechanical properties and microstructure of A l - F e - X alloys

TABLE 3. Lattice spacing of the S and related phases Phase

S

S

S

-AI-Fe-Si

-AI-Fe-Si

Icosahedric

Morphology

Cellular microcryst,

Cellular microcryst,

Cellular microcryst,

Globular precipit,

Globular precipit,

Globular microcryst.

Alloy Observation Ref.

AI-Fe TEM Jacobs et al. [12]

AI-Fe-Mo-Zr TEM Sainfort [13]

AI-Fe-Mo-Zr TEM Present study

AI-Fe-V-Si TEM Present study

AI-Fe-V-Si X-diff. Present study

AI-Mn X-diff. Shechtman and Blech [14]

4.00 (vw)

4.50 (w) 3.99 (w)

4.44 (10) 3.97 (39)

3.65 3.37 2.47 2.16 2.04 1.49 1.25

3.63 (11) 3.36 (13) 2.46 (7) 2.15 (53) 2.04 (100)

Lattice spacing and intensity a

3.80 (vw)

3.83 (7)

3.64 (vw) 2.52 (vw) 2.08 (s) 1.47 (w) 1.27 (vw) 1.20 (vw) 1.08 (vw)

2.50 (vw) 2.07 (s) 1.45 (w) 1.24 (vw)

2.08 (s*) 1.47 (w) 1.26 (w)

1.06 (vw)

1.08 (w)

(vw) (vw) (vw) (w) (w) (vw) (vw)

3.33 (2) 2.17 2.06 1.50 1.28

(100) (50 + 50) (7) (21)

1.10 (5 +6)

aVW= very weak, w = weak, s = strong, number = relative intensity, *thick ring.

YS (M Pa) 700. 600" 500.

-O- AIFe 151a

400.

•O- AIFe 4011

300"

•ll.- AIFeMoZr

200.

•dk- AIFeVSi

100. 0

0

I

I

I

I

I

I

I

100

200

300

400

500

600

700

(a)

T (K)

UTS (MPa) 800. 700. 600.

Fig. 5. At 3 Zr precipitates in an undeformed particle. A1-Fe-Mo-Zr alloy, dark field with (110) spot.

-O- AIFe 1511

500.

•0 - AIFe 4011

400, •D- AIFeMoZr

300'

fact, regarding the binary alloy, it is present only at low strain rate and for the material having the finest original particle size. The tensile curves always show a very low strain hardening. In order to characterize this strain hardening, an attempt was made to fit the curves by classical laws, such as

(1)

0 = (71 ---Krp n

or, after Kocks [18], (7= (72 - a exp(

-

%)

(2)

ep being the plastic strain (71,/72, a, K and n being constants at a given temperature.

"&- AIFeVSi

200 100'

0(b)

:

;

:

=

,"

I

I

100

200

300

400

500

600

700

T ~K~

Fig. 6. Tensile properties of the A1-Fe-X alloys as a function of temperature: (a) yield strength; (b) ultimate tensile strength.

However the only good fit was obtained by the law: O = (70- K(ep + c0) -N

(3)

with e0 ~<0.2%. Typical values of these parameters for all the alloy are given in Table 5.

J. C. EhrstrOm, A. Pineau

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Mechanical properties and microstructure of Al-Fe-X alloys

Here, the higher the value of N, the lower the strain hardening after a few tenths of percent deformation. In Fig. 8, the evolution of n = 0 In o / 0 In e and N at e_ = 1% is plotted against temperature. It appears that (il between the two binary alloys, the one with the

n(ep=1%) 0,12 -

0,08

5.10-2s-1 ~ AIFe 15 microns

20

c15 t i 10 o n 5

~.~,,~ ~lFe 40 miJ'ons 5.10 "2S -1

1

1

~

.

,

~

AIFe 15 microns

r 25 e d

AI Fe 1~

0,06

a43

u

AIFe40microns

0,1

(%) r e35 a 30

59

•

/

~

.

0,04

~

/

0,02

I

I

I

300

350

400

0 250

•

~"~'~

m

5.10 "4s -1

!a)

I

I

450 500 temperature (K)

I

I

I

550

600

650

N 1,2 I 100

I 200

I I 300 400 temperature (K)

I 500

I SO0

I 700

Fig. 7. Ductility of the binary alloys.

0,8

T A B L E 4. Ductility of the quaternary alloys Alloy

AI-Fe-Mo-Zr

A1-Fe-V-Si

Extr. ratio

o,6'

T

Elongation

A r e a reduction

(K)

(%)

(%)

77 293 525

1 8 5

-18 12

9

293

12

23

9

293 525

18 12

--

6.25

ns / ~''''''-'~°.~//~'-FeMoZr j/~''F~/~:'rnicr/°ns

0,4, o, 2

Al~cr

o

I 200

150

o

I 250

I 300

I I I 350 400 450 temperature (K)

I 500

I 550

I 600

I 650

(b) Fig. 8. (a) n = 0 In o / 0 In e at ep = 1% as a function of temperature; N as a function of temperature.

T A B L E 5. Parameters of the work hardening laws Alloy

T (K)

Strain rate (s- 1)

~r0(MPa)

K (MPa)

N

e0

A I - F e 15/~ A I - F e 15 ¢t A1-Fe 15/~ A I - F e 15/~ A I - F e 15/~ A I - F e 40/~ A I - F e 40 p A I - F e 40 p A I - F e 40 ~ AI-Fe40/~ A1-Fe 40/~ A I - F e 40/~ A1-Fe 4 0 / t A1-Fe-Mo-Zr AI-Fe-Mo-Zr AI-Fe-Mo-Zr AI-Fe-V-Si AI-Fe-V-Si

173 295 525 525 525 295 295 425 425 525 525 625 625 295 525 625 295 525

5.10 .4 5.10 -4 5.10 -4 5.10 -3 5.10 -2 5.10 -4 5.10 -2 5.10 -4 5.10 -5 5.10 4 5.10-2 5.10 -4 5.10 .3 5.10 -4 5.10 -4 5.10 .4 5.10 -4 5.10 -4

882 401 233 254 270 415 566 295 278 211 266 152 158 566 302 193 541 289

282 9.47 0.0265 0.0350 0.019 37.22 120.6 6.22 3.70 0.410 1.68 0,077 0.046 40.1 2,72 41,3 19.6 0.060

0.111 0.413 1.13 1.11 1.06 0.259 0.142 0.404 0.448 0.695 0.576 0.844 0.983 0.216 0.462 0.557 0.376 0.998

2.10 -4 2.10 -4 6.10 -4 6.10 -4 6.10 -4 4.10 -4 4.10 -4 4.10 -4 5.10 -4 5.10 -4 6.10 -4 5.10 -4 5.10 .4 2.10 -4 5.10 .4 5.10 -4 3.10 4 6.10 -4

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Mechanical properties and microstructure of A l - F e - X alloys

coarser particle size has a higher strain hardening rate; and (ii) the strain hardening rate of the A1F e - M o - Z r alloy is much less temperature dependent than the strain hardening rates of the two other alloys. At high temperature the asymptotic stress a 0 is reached after 1% deformation, whereas at room temperature fracture occurs before a 0 is reached. Thus, at high temperature the UTS equals a0. Variation of the minimum creep strain rate as a function of applied stress a a is reported in Fig. 9, together with the results of tensile tests, the plotted stress being the UTS (UTS = a 0 at 150, 250 and 350 °C). It is worth noting the continuity of the creep and tensile tests carried out at the same temperature. The approximate activation energy was determined (Q = 250 kJ mol-~), and this was used to calculate the Zener parameter, Z = g e x p ( Q / R T ). Figure 10 shows that this parameter is able to unify the whole set of data obtained on the binary alloy. This result is not surprising since during a tensile test at high temperature, because of the absence of work hardening, a constant stress results essentially from a constant strain rate, which is virtually identical to a creep test. In fact, the apparent activation energy was found to vary from 160 to 345 kJ mol- 1, according to the temperature. The values are given in Table 6, together with

the strain rate sensitivity parameter m = 1/nv, where

n v = (0 In g/O In a) at constant temperature, and

Q =RT2(O In g/O T)

1 E-02

• 623K

creep

D 623K

tension

1E-05

• 523K

creep

1 E-06

4) 523K

tension

• 423K

creep

1 E-03

m~pA

1 E-04

O

1 E-07

(5)

at constant stress. The values of Q and nv = 1/m are m u c h higher than those observed for dislocation creep, i.e. n v = 3 - 5 and Q = 1 4 2 k J m o 1 - 1 [15], the latter being the activation energy for self diffusion in aluminum. Similar results have been found by Guyot [19] in similar materials, i.e. sintered aluminum powders. A classical way of rationalizing these observations is to introduce an internal stress (see ref. 20), writing

=A( a - qi) n,',,e x p ( - Qo/RT )

(6)

instead of g = A a "V e x p ( - Q / R T )

Assuming that nv,,= 5, as indicated in Table 6, the values of a i and those of Q0 were calculated. All the

details have been given elsewhere [11]. Table 6 shows that the values of ai are quite large compared to the strength of the materials, since they are close to the yield strength.

stress 300

strain rate (s-1 ) 1E-01

(4)

MPa)

,:,°

250

2O0

13

13 creep; AI-Fe 15p

~P D~"

150

cg"

"

& tension; AI-Fe 151x " tension; AI-Fe 40ix

1 E-08

100 =_j

1E-09 1E-10

A 423K

.* I

10

tension

50

I

lO0 stress (MPa)

~O

1000

Fig. 9. Strain rate as a function of stress at high temperature for the binary alloy. For creep tests the stress is the initial applied stress and the strain rate is the minimum strain rate. For tensile tests the stress is the UTS (UTS = a0) and the strain rate is the initial applied strain rate.

0 25

I

I

I

i

i

i

i

30

35

40

45

50

55

60

In (Z)

Fig. 10. Relation between high temperature tension test and creep test. Z = e exp ( Q / R T ) with Q = 250 kJ tool- 1.

TABLE 6. Parameter of the constitutive equation at high temperature for the binary alloys Temperature (K)

150

nv = 1/m Q (kJ mo1-1) n o~ (MPa) Q0 (kJ mol- 1 )

26 160 5 258 111

175

200

225

177

190

246

250 22.7 250 5 187 128

300 290

350 12.5 345 5 98 144

J. C. Ehrstr6m, A. Pineau

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Mechanical properties and microstructure of AI-Fe-X alloys

3.3. Microstructure after deformation The microstructure of the tensile specimens was examined by TEM (Figs. 11 and 12). These observations show that after high temperature deformation the microstructure of the binary alloy is essentially unmodified. It remains of a recovered type, with subgrain boundaries which are still pinned by precipitates, as shown in Fig. 11. Quantitative measurement of the grain size after deformation at 250°C and = 5 10 -4 S - l, reported in Fig. 4, shows that the substructure of the material is not changed and the subgrain size remains at an average value of 0.6/zm. This indicates that under these conditions the material exhibits extensive dynamic recovery during deformation. However, at lower temperature, the microstruc-

61

ture of the material is considerably modified. In particular, a very dense dislocation forest is produced, interacting with precipitates located in the subgrains (Fig. 12). After large plastic strains (/> 10%) at low temperature (173 K), dislocation cells are formed. This is the reason why we have taken into account not only subgrains but also the dislocation cells in Fig. 4. The microstructure of the A1-Fe-Mo-Zr alloy is less recovered after high temperature deformation than it is for the binary alloys. It has been observed that zones with higher dislocation density are associated with AI3Zr precipitates ]11].

4. Discussion

4.1. Origins of the strength of Al-Fe alloys

50 n m

"lm

Fig. 11. Subgrain boundaries in the A I - F e alloy (15 At) after 3% deformation at 525 K.

Different strengthening mechanisms are operating in these A1-Fe-X alloys. Rapid solidification can result in a highly supersaturated solid solution with corresponding strengthening. X-ray diffraction measurements of the lattice parameter [l 1] and calculations using the Cottrell expressions derived for solid solution strengthening [21, 22] show that this type of strengthening must be negligible. Rapid solidification also produces fine precipitation, which induces Orowan strengthening. Premkumar [23] assumes that this is responsible for the strength of AI-Fe-Ni alloys. In the alloys that we have studied the distance L between precipitates is of the order of 0.1 to 0.5/~m (microstructures are rather heterogeneous), except for the A13Zr precipitates which are heterogeneously distributed and, further, probably shearable because of their extremely small size. Their contribution can thus be ignored. Estimation of the Orowan strengthening gives Ao = 3.06 tzb/L

(7)

where kt is the shear modulus, b the Biirgers vector, and 3.06 the Taylor factor. Thus the strengthening effect Ao could be of the order of 150 MPa. But, as we have seen previously, the precipitates are mainly located at the subgrain boundaries, so that this value is probably an overestimation. Several authors [24-26] have attributed the strengthening of materials of the same type to their fine grain size. The well-known Hall-Petch relationship between grain size and yield strength is written as YS = of + k d -a

Fig. 12. Microstructure of the A I - F e alloy ( 15 At) after 3% deformation at 173 K.

(8)

where of is a friction stress, k is almost a constant and a = 0 . 5 . For pure aluminum, o f = 1 6 M P a and k = 70 M P a / z m - ~[27]. However, in the case of subgrains instead of grains, it has been suggested that a is closer to 1 than to 0.5

J.C. Ehrstr6m, A. Pineau

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Mechanical properties and microstructure of Al-Fe-X alloys

[28-31], while the value of k is increased owing to the presence of particles along the subgrains [32]. Taking the values of at, k and a given in ref. 30 on a similar alloy, we obtain YS = 50 + 150 d -l = 300 MPa

(9)

with d = 0.6 #m, which corresponds to the mean value of the subgrain size (Fig. 4). This value of the strengthening effect associated with the subgrain is larger than the Orowan effect. Therefore, one can conclude that the strengthening of AI-Fe alloys is mainly due to fine subgrain size which is a consequence of the pinning of boundaries by precipitates.

4.2. Work hardening behavior as a function of temperature Below room temperature, work hardening occurs by multiplication of dislocations. This results in the formation of a dislocation forest, as observed by TEM (Fig. 12). At very low temperature (173 K) no dynamic recovery takes place, and dislocation cells are formed as in conventional materials. These deformation substructures are associated with relatively large values for the work hardening coefficient, n (n >/0.10). At elevated temperature, typically 250°C, the situation is quite different with the work hardening coefficient being considerably reduced (n~<0.04). Examination of the tensile curves shows that a stationary stress is already reached after about 1% plastic strain. This macroscopic feature is associated with dynamic recovery, as observed by TEM (Figs. 4 and 11). Looking at the different behaviors of the two binary alloys, it appears that this early transition to high temperature deformation modes is a consequence of the fine grain size. This is best illustrated in Fig. 13 where the parameter

1,1

•

1

0,8 0,7 0,6 0,5

50

I 200

I 250

I 300

(10)

is plotted against temperature after normalization (after Jensrud [24]). In this equation /~ is the shear modulus. In Fig. 13, it is observed that the high temperature regime occurs at lower temperature for the small grain size material. The reason for this behavior is that the density of dislocation sinks formed by subgrain boundaries is high, so that dislocations have a reduced mean free path and have no opportunity to build a dense forest. They are annihilated by the rise in subgrain boundaries. It is worth noting that similar macroscopic observations concerning the mechanical properties of AI-Si alloys have been reported by Stewart and Martin [33] and Humphreys and Kalu [34]. In these materials, the authors attribute the decrease in strain hardening rate above 150 °C to recovery taking place in the vicinity of Si particles. A special situation is observed in the A I - F e - M o - Z r alloy. Figure 13 suggests that the high temperature regime begins at lower temperature compared with the binary alloys. However, Fig. 8 shows that the strain hardening rate of the material at elevated temperature is higher than that of the other alloys. TEM observations show that, in the A I - F e - M o - Z r alloy deformed at elevated temperature, recovery was not fully obtained, in particular in the zones where AI3Zr precipitates were present. This suggests that this phase might be extremely efficient for maintaining a significant work hardenability at elevated temperature, like the ~,' phase in Ni-based superalloys which has the same crystallographic texture. Another feature of the high temperature behavior of the binary AI-Fe alloys is their low strain rate sensitivity, as observed in Fig. 9 (see also Table 6). It was shown that this situation can be interpreted as a consequence of the presence of a high internal stress o i. A straightforward explanation for this high internal stress is the fact that the major part of the strengthening at high temperature is due to the fine subgrain size, especially as oi is found to be of the same order of magnitude as the yield stress.

0

0,9

0,4

r = YS(T)/YS(300 K) ×/~(300 K)//~(T)

I 350

I 400 T (K)

I 450

I 500

I 550

I 600

I

650

Fig. 13. Normalized yield stress as a function of temperature (see text). Slope of the lines taken from Jensrud [24].

4.3. High temperature ductility In Fig. 7 it was observed that the fine grained binary alloy exhibited a ductility dip at about 250 °C when the material was deformed at low strain rate (5 10 -4 s - l ) . This has also been observed by other investigators [1, 10]. These authors attributed the observed ductility dip to dynamic strain ageing (DSA) associated with the diffusion of irons atoms in solid solution. This explanation cannot account for the grain size effect observed in the binary alloys (Fig. 7), as X-ray measurements showed that the amount of iron in solid solution was the same in both alloys I11]. Moreover, in the present study, the jerky flow which is generally

J. C. EhrstrOm, A. Pineau

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Mechanical properties and microstructure of AI-Fe-X alloys

observed over certain temperature and strain ranges when DSA is present was never observed. T h e s e observations suggest that the origin of the ductility dip might rather be the temperature dependence of both the work hardening coefficient, n, and the strain rate sensitivity, rn = 1 Inv. With regard to the theory of necking of tensile bars, it has been shown by Hart [35] that uniform strain is an increasing function of both n and m coefficients. In all cases, it was observed that the strain rate sensitivity was quite small, as has already been discussed. As far as the work hardening coefficient is concerned, Fig. 8 shows that the values of n in the fine grain material are lower than in the large grain alloy. Similar values of m and n are, however, observed at high and medium strain rates. T h e effect of strain rate on ductility is therefore not explained by the above observation. We believe that diffusion enhanced damage of the material is possible at a relatively low temperature such as 525 K, because of the fineness of the microstructure, the very fine substructure promoting pipe diffusion. However, further work is clearly necessary to determine this particular point unambiguously.

5. Conclusions (1) A I - F e - X alloys have interesting mechanical properties up to 550 K. T h e A I - F e - M o - Z r alloy also exhibits high strength at r o o m temperature. (2) T h e ductility dip generally observed with these alloys is restricted to the cases of the alloys of the finest microstructure, and to low strain rates. (3) T h e strengthening of these materials is mainly attributed to their fine subgrain size, which does not vary during high temperature deformation because their boundaries are pinned by the precipitates. Additional Orowan and dislocation forest strengthening takes place at low temperature. (4) Owing to the fine subgrain size, dynamic recovery occurs at relatively low temperature and plastic strain, which results in a low strain hardening rate. Recovery is apparently inhibited by coherent or semicoherent A13Zr precipitates in the A I - F e - M o - Z r alloy. (5) T h e low strain rate sensitivity (m = 0.05) is attributed to the existence of a high internal stress, close to the yield stress of the materials. This internal stress originates from the fine subgrain size (0.6/zm). (6) Dynamic strain ageing does not account for the whole set of data concerning the intermediate temperature ductility dip of this type of alloy. It is proposed that this ductility dip is due to aspects of the constitutive behavior: low strain rate sensitivity and d e p e n d e n c e of the strain hardening coefficient on grain size. However, the strain rate effect is not explained.

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Acknowledgment T h e grant provided by Pechiney to J.C.E. is gratefully acknowledged.

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