High temperature creep behaviour of an Al-8.5Fe-1.3V-1.7Si alloy reinforced with silicon carbide particulates

Materials Science and Engineering A283 (2000) 172 – 180 www.elsevier.com/locate/msea

High temperature creep behaviour of an Al-8.5Fe-1.3V-1.7Si alloy reinforced with silicon carbide particulates J. C& adek a,*, K. Kucharˇova´ a, S.J. Zhu b b

a Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Z& izˇko6a 22, 616 62 Brno, Czech Republic Department of Mechanical Engineering and Intelligent Systems, The Uni6ersity of Electro-Communications, Chofu, Tokyo, 182 -8585 Japan

Received 12 October 1999; received in revised form 1 December 1999

Abstract The creep behaviour of an Al-8.5Fe-1.3V-1.7Si (the 8009Al type, all numbers indicate wt.%) alloy reinforced with 15 vol.% silicon carbide particulates — the Al-8.5Fe-1.3V-1.7Si-15SiCp composite — is investigated at three temperatures ranging from 623 to 723 K. The measured minimum creep strain rates cover seven orders of magnitude. The creep behaviour is observed to be associated with the true threshold stress that decreases more strongly with increasing temperature than the shear modulus of aluminium. The minimum creep strain rate is controlled by the lattice diffusion in the composite matrix, and the true stress exponent is close to 5. The results are compared with those obtained investigating the creep behaviour of an unreinforced Al-8.5Fe-1.3V-1.7Si alloy in the same temperature range. The creep strength of the composite as characterised by the minimum creep strain rate is found to be up to six orders of magnitude higher in the composite than in the alloy. This creep strengthening is attributed to a much higher true threshold stress in the composite than in the alloy, which is primarily due to finely dispersed alumina particles appearing in the composite matrix as a result of composite fabrication. The creep behaviour is interpreted in terms of athermal detachment of dislocations from interacting particles admitting a temperature dependence of the relaxation factor that characterises the strength of dislocation/particle interaction. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Al-8.5Fe-1.3V-1.7Si-15SiCp composite; Creep; True stress exponent; True threshold stress; Creep strengthening mechanisms

1. Introduction An Al-8.5Fe-1.3V-1.7Si (the 8009Al type, numbers indicate wt.%) alloy processed by fast solidification and powder metallurgy route exhibits remarkable creep resistance up to temperatures around 700 K [1 –5]. This creep resistance is due to high volume fraction (  0.27) of fine incoherent particles of the intermetallic Al12(Fe,V)3Si phase and low coarsening rate of these particles at high temperatures. For the Al-8.5Fe-1.3V1.7Si alloy (simply ‘alloy’ in the following) the true threshold creep behaviour is characteristic. It is well known that the Young’s modulus of an aluminium alloy can be increased significantly by discontinuous reinforcement with hard unshearable ceramic particulates, short fibres or whiskers even at * Corresponding author. Tel.: +420-5-7268374; fax: + 420-541212301. E-mail address: [email protected] (J. C& adek)

temperatures as high as 700 K [6]. This fact motivated Peng et al. [4,7] and Zhu et al. [8] to reinforce the alloy with silicon carbide and/or silicon nitride whiskers. Beside increasing the Young’s modulus, the discontinuous reinforcement can be expected to introduce the load transfer effect (e.g. Refs. [9–12]) that enhances the creep strength, although relatively slightly. As to the Al-8.5Fe-1.3V-1.7Si-15SiCp composite investigated in the present work, (simply ‘composite’ in the following) a far more important effect on the creep strength follows from the fact that as a result of the alloy atomisation (see Section 2), fine alumina particles appear in the composite matrix fabricated by powder metallurgy route [13]. It should be reminded that, generally, the creep strength of the composite is controlled by creep strength of the composite matrix [13–17], though, also the already mentioned load transfer can play a role. Peng et al. [4] have shown that the true threshold stress in creep of an Al-8.5Fe-1.3V-1.7Si alloy reinforced with

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Table 1 Al-8.5Fe-1.7Si-1.3V alloy reinforced with 15 vol.% silicon carbide particulates – an Al-8.5Fe-1.7Si-1.3V-15SiCp compositea T = 623 K

T= 673 K

T= 723 K

s(MPa)

o; m(s−1)

s(MPa)

o; m(s−1)

s(MPa)

o; m(s−1)

160.0 162.0 165.0 170.0 175.0 185.0 210.0 220.0 242.0 260.0

3.80 1.80 7.80 8.00 5.60 2.90 4.60 1.05 4.70 1.00

122.0 125.0 128.0 131.0 131.0 135.0 142.0 157.0 178.0 200.0

2.80 1.75 3.05 1.55 2.75 1.35 6.10 5.51 3.75 2.15

90.0 91.0 92.0 94.0 98.0 112.0 120.0 130.0 130.0 160.0

2.75 1.40 8.50 7.82 7.05 7.51 5.50 3.50 1.75 5.05

a

E-10 E-9 E-9 E-8 E-7 E-6 E-5 E-4 E-4 E-3

E-10 E-9 E-8 E-7 E-7 E-6 E-6 E-5 E-4 E-3

E-10 E-9 E-9 E-8 E-7 E-6 E-5 E-4 E-4 E-3

Minimum creep strain rates o; m at various temperatures T and applied stresses s.

silicon carbide whiskers (the volume fraction fw $ 0.15) is significantly higher than the true threshold stress in the unreinforced alloy. This can be explained exclusively by the above effect of fine alumina particles present in the composite matrix, but not in the matrix alloy. In a previous paper [18], the creep behaviour of an Al-8.5Fe-1.3V-1.7Si alloy at three temperatures ranging from 623 to 723 K was investigated. The measured minimum creep strain rates covered seven orders of magnitude, the lowest of them were well below 10 − 9 s − 1. The creep behaviour was found to be associated with a true threshold stress. The threshold stress was interpreted in terms of athermal detachment of dislocations from fine incoherent Al12(Fe,V)3Si phase particles. To account for the temperature dependence of the true threshold stress much stronger than that of the shear modulus, the relaxation factor kR characterising the strength of the attractive dislocation/particle interaction [19,20] was assumed to increase with increasing temperature. In the present paper, the results of an investigation of creep behaviour of an Al-8.5Fe-1.3V-1.7Si alloy reinforced with 15 vol.% silicon carbide particulates (SiCp) are reported. The investigation of the discontinuous Al-8.5Fe-1.3V-1.7Si-15SiCp composite was undertaken to show that, at least at temperatures ranging from 623 to 723 K, the measured minimum creep strain rates can be satisfactorily interpreted in terms of athermal detachment of dislocations from fine interacting Al12(Fe,V)3Si phase particles and also fine alumina particles, the latter appearing during the composite fabrication.

the form of a rod 12 mm in diameter. The alloy was atomised (powder size 20 mm) and the powder was mixed with nominally 15 vol.% silicon carbide particulates of the mean diameter of  4.5 mm. The mixed powders were consolidated and extruded at a temperature  830 K to a rod 12 mm in diameter. The resulting mean grain diameter was found slightly less than 1 mm. The structure of the as extruded composite was found to be reasonably homogeneous as compared to the matrix Al-8.5Fe-1.3V-1.3 alloy, although not only

2. Material and experimental procedures The Al-8.5Fe-1.3V-1.7Si alloy processed by fast solidification and powder metallurgy [18] was available in

Fig. 1. Minimum creep strain rates o; m plotted against applied stress s for various temperatures in double logarithmic co-ordinates.

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aluminium alloys as well as aluminium or aluminium alloy matrix composites fabricated by powder metallurgy route small volume fractions of very fine alumina particles are necessarily present. This is apparently the

Fig. 2. Relations between apparent stress exponent mc and applied stress s. m exp means the apparent stress exponent obtained as the c (( ln o; m/( ln s)T derivative (cf. Fig. 1), m calc means the apparent stress c exponent calculated by means of Eq. (2).

Fig. 4. Temperature dependences of true threshold stress sTH and true threshold stress to shear modulus G ratio, i.e. sTH/G.

Fig. 3. o; 1/n m plotted against s for the true stress exponent n = 5. The true threshold stress sTH is defined as an applied stress at which the minimum creep strain rate o; m = 0; R is the correlation coefficient.

the Al12(Fe,V)3Si particles but also the silicon carbide particulates (SiCp) were mostly aligned to the extrusion direction as it is usually the case with discontinuous composites processed by powder metallurgy route. The composite matrix exhibited only weak texture. It consisted of the Al12(Fe,V)3Si phase particles of slightly less than 50 nm in mean diameter embedded in the matrix solid solution. The particles of this phase were fairly homogeneously distributed. It is well known that in the

Fig. 5. Relation between o; mb 2/DL and (s− sTH)/G in double logarithmic co-ordinates. DL is the coefficient of lattice self-diffusion in aluminium. Similar relation for unreinforced Al-8.5Fe-1.3V-1.7Si alloy is shown for comparison.

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Table 2 Al-8.5Fe-1.3V-1.7Si-15SiCp compositea T =623 K −1 Q calc ) c (kJ mol 183.4 a

T= 673 K m calc c 7.22

−1 Q calc ) c (kJ mol 187.4

T = 723 K m calc c 6.55

−1 Q calc ) c (kJ mol 192.2

m calc c 6.10

Values of Q calc and m calc for various temperatures and an applied stress of 500 MPa. c c

case of the matrix of the composite investigated in the present work, too. A detailed studies of the structure of similar composites in prior to creep as well as in crept conditions was performed by Peng et al. [4,5,7] and Zhu et al. [8]. From the composite rod, specimens for tensile creep tests 4.0 mm in diameter and 25.0 mm in gauge length were machined. The constant tensile stress creep tests were performed at temperatures 623, 673 and 723 K in purified argon; the testing temperatures were controlled to within 0.5 K. The creep elongation was measured by means of linear variable differential transducers coupled with a digital data acquisition system. The measured minimum creep strain rates covered seven orders of magnitude, the lowest of them were well below 10 − 9 s − 1. All the creep tests were run well into the tertiary stage and interrupted. Generally, no steady state stage was observed, only the minimum creep strain rate o; m could be defined. 3. Results and analysis In Table 1, the minimum creep strain rates o; m measured at various temperatures T and applied stresses s are listed. In Fig. 1, the minimum creep strain rates are plotted against applied stress in double logarithmic co-ordinates. For any temperature under consideration, the o; m(s) relation shows true threshold creep behaviour. In fact, the apparent stress exponent of minimum creep strain rate, mc =(( ln o; m/( ln s)T increases with decreasing applied stress reaching extremely high values at the applied stresses, at which the minimum creep strain rates decrease to  10 − 9 s − 1. This is illustrated in Fig. 2, in which m exp means the value of c mc obtained as the above derivative of ln o; m(s) with respect to ln s. Because of very strong applied stress dependence of o; m (Fig. 1), it was possible to estimate the apparent activation energy of creep, Qc =[( ln o; m/(( − 1/RT)]s at one applied stress only, namely at s= 160 MPa, from the o; m(s,T) creep data obtained in the present investigation. However, it was shown [21,22] that the values of Qc calculated from the relation to be given later in this section — Eq. (3) — are in remarkable agreement with the values of Qc estimated in a conventional way from the temperature dependences of o; m(s). Thus, later, the calculated Qc(s, T) relation will be presented.

To determine the true threshold stress for the testing temperatures under consideration, the conventional linear extrapolation technique (e.g. Refs. [21–24]) has been applied. Thus, in Fig. 3, o; 1/n m is plotted against s in double linear co-ordinates for the true stress exponent n equal to 5 for which the best linear fit between o; 1/n m and s was obtained. The values of the threshold stress sTH, obtained extrapolating the relations between o; 1/5 m and s to o; m = 0, are given in the figure together with the correlation coefficients R. Values of the correlation coefficient R range from 0.9942 to 0.9993 (Fig. 3). These sTH(s) data strongly suggest the true stress exponent n close to 5 and the true threshold stress sTH depending strongly on temperature. Nevertheless, sometimes the value of the true stress exponent of 8 is preferred (e.g. Refs. [1,2]) to 5, especially when the measured strain rates do not cover more than four or even than three orders of magnitude. The value of n= 8 cannot be accepted in the present investigation: first because of the above values of the correlation coefficient for n=5 and, second, because of the o; 1/n m versus s relations for n= 8 are not linear and do not define true threshold stresses. This can be easily found using the o; m(T,s) creep data listed in Table 1.

Fig. 6. Relations between the apparent activation energy Qc = Q calc c calculated by means of Eq. (3) and applied stress s. In the figure, also exp the values of Qc =Q c estimated by a conventional way from the o; m(s,T) creep data for s=160 MPa are shown. DHL is the activation enthalpy of lattice self-diffusion in aluminium.

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versus s and sTH/G versus s relations can be well approximated by straight lines. It can be seen that not only sTH but also sTH/G decreases rather strongly with increasing temperature. This phenomenon is common to oxide dispersion strengthened alloys, specifically ODS Al alloys [15] (ODS means oxide dispersion strengthening) as well as to aluminium alloys processed by powder metallurgy [26] and to discontinuous aluminium and aluminium alloy matrix composites fabricated by powder metallurgy route, e.g. Refs. [12,15,27]. Considering all the above results as well as some other reported results [14,21,22], it is justified to plot o; mb 2/DL versus (s − sTH)/G assuming the minimum creep strain rate to be lattice diffusion controlled, and accepting n= 5 to determine values of sTH for the temperatures under consideration; DL is the coefficient of lattice self-diffusion in aluminium [28] and b is the length of Burgers vector in aluminium. Such a plot is shown in Fig. 5. The plot can be approximated by a single straight line. This justifies the assumption of the lattice diffusion in the matrix as the minimum creep strain rate controlling process as well as the assumption of n close to 5. Thus, the minimum creep strain rate as a function of temperature and applied stress can be expressed as (see e.g. Refs. [21,22]). Fig. 7. A comparison of the relation between o; m and s for the composite with a similar relation for unreinforced alloy; temperature 673 K.



s − sTH o; mb 2 =A DL G



n

, n=5

(1)

where A is a dimensionless constant. Combining this creep equation with the definition equation of the apparent stress exponent, mc = (( ln o; m/( ln s)T, on one side and the definition equation of the activation energy of creep, Qc = [(lno; m/((− 1/RT)]s on the other, the following expressions for the apparent stress exponent mc and the apparent activation energy Qc are obtained: mc =

ns s− sTH

(2)

Fig. 8. A comparison of the relation sTH and T for the composite with a similar relation for the unreinforced alloy.

In Fig. 4, the values of the true threshold stress are plotted against temperature. In this figure, also the temperature dependence of the sTH/G ratio is shown; G is the shear modulus of aluminium [25]. Both the sTH

Fig. 9. Calculated values of the relaxation factor kR plotted against temperature. k crit is the ‘critical’ value of this factor that followed R from the analyses of Arzt and Wilkinson [19] and Arzt and Ro¨sler [20].

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and Qc = DHL −

nRT G

2





G dsTH n − 1 dG + . s− sTH dT n dT

(3)

In Eq. (3), DHL is the activation enthalpy of lattice self-diffusion in aluminium (DHL =142 kJ mol − 1 [28]). From Eqs. (2) and (3) it follows that both mc and Qc should depend on both the applied stress and temperature, since sTH/G depends on temperature (see Fig. 4). Now, values of the apparent stress exponent mc = m calc can be estimated for various temperatures and c applied stresses; m calc is the apparent stress exponent c calculated by means of Eq. (2). The relations between mc = m exp and s for 623, 673 and 723 K were obtained c from the experimental o; m(T,s) creep data, Fig. 1, and shown in Fig. 2. The values of m calc for the temperac tures under consideration are shown in the same Fig. 2 for comparison. It can be seen that the agreement and m calc is excellent and, thus, strongly between m exp c c supports the validity of the creep Eq. (2) with n= 5. Even to an applied stress as high as 500 MPa the values of m calc correspond, which are still slightly higher than c that of the accepted true stress exponent n, i.e. 5 (Table 2). Similarly, Eq. (3) can be used to calculate the apparent activation energy Qc =Q calc as a function of applied c stress and temperature. As mentioned above, Qc = Q exp c can be estimated in a conventional way from the temperature dependence of o; m(T,s) creep data for one applied stress, namely s = 160 MPa only. However, it should be emphasised once again that the calculated were shown [21 – 23] to be in very values of Qc = Q calc c good agreement with the values of Qc =Q exp deterc mined from the temperature dependences of o; m(s). For the composite investigated in the present work, the relations between Q calc and applied stress s for the c temperatures under consideration are shown in Fig. 6. At any given temperature, the activation energy Q calc c decreases with increasing stress and at any given applied stress it decreases with increasing temperature. Also, in the figure, the value of Qc =Q exp estimated in c an conventional way from the o; m(s, T) creep data for s =160 MPa are shown. The agreement between Q calc c and Q exp for the applied stress under consideration is c very good. At stresses only slightly higher than the respective true threshold stress the apparent activation energy is many times higher than the activation enthalpy DHL. Even at an unrealistically high applied stress as high as 500 MPa the calculated values of the apparent activation energy are still significantly higher than that of DHL for aluminium, which is illustrated in Table 2. At any given applied stress, the energy Q calc c increases with decreasing temperature. All this is in qualitative agreement with the results obtained for an Al-30SiCp composite [21,23] as well as an ODS Al5Mg-30SiCp composite [22], for which the values of the

177

apparent activation energy could be determined performing the conventional analysis of the experimental o; m(T,s) creep data.

4. Discussion

4.1. Effect of discontinuous reinforcement on creep beha6iour of Al-8.5Fe-1.3V-1.7Si alloy In Fig. 7, the relation between o; m and s for the composite at 673 K is compared with a similar relation [18] for the alloy. It can be seen that at an applied stress of  125 MPa the minimum creep strain rate in the composite is six orders of magnitude lower than that in the alloy. In Fig. 5, the relation between o; mb 2/DL and (s− sTH)/G for the composite is compared with a similar relation for the alloy [18]. The true stress exponent n for the alloy is slightly higher than 5 which is not the case of the composite. Nevertheless, from the figure it clearly follows that the creep strengthening associated with the reinforcement is essentially due to the effect of this reinforcement on the true threshold stress, Fig. 7. As already mentioned in Section 1, the effect of reinforcement on the true threshold stress is ‘indirect’. In fact, to reinforce the alloy with silicon carbide particulates, the alloy had to be atomised and, consequently, fine alumina particles appeared in the final composite matrix. With the presence of fine alumina particles in the composite matrix, significantly higher true threshold stress values are associated, Fig. 8. From Fig. 5 it follows that the load transfer effect [10–12] does not play any significant role in creep strengthening of the alloy due to its reinceforcement. In fact, the o; mb 2/DL versus (s− sTH)/G relations for unreinforced Al-8.5Fe-1.3V-1.7Si alloy and Al-8.5Fe-1.3V1.7Si-15SiCp composite are (except a very small difference in the values of the true stress exponent n), practically identical (c.f. Fig. 8 in Ref. [10]). In this respect, the creep behaviour of the present composite differs from that of an Al-8.5Fe-1.3V-1.7Si-15SiCw composite investigated by Peng et al. [4] (see Fig. 8 in ref. [4]). This difference can be hardly attributed to different types of reinforcement, the particulates on one side and the whiskers on the other, and remains to be explained. It should be recalled that the fabrication of the present composite differs from that of the alloy. While the alloy was processed by fast solidification and powder metallurgy, the fabrication of the composite included in addition the atomisation of the alloy, mixing of the alloy powder with SiCp powder, consolidation of the mixture and extrusion, see Ref. [18] and Section 2 of the present paper. Due to the alloy atomisation a small volume fraction of very fine alumina particles

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appear in the composite matrix. Presence of these particles in the composite matrix contributes to the true threshold stress of the composite. However, the differences of fabrication routes of the alloy and the composite cannot explain the apparent absence of the load transfer effect in the composite.

4.2. Temperature dependence of the true threshold stress The true threshold creep behaviour of the composite investigated here is associated with the presence of fine incoherent Al12(Fe,V)3Si phase particles as well as the presence of small volume fraction of alumina particles in the composite matrix. Both these types of particles are incoherent with the composite matrix. Hence, they are expected to attract moving dislocations under creep conditions. Then, the controlling step of climb of a dislocation past a particle is represented by detachment of the dislocation from the particle after the climb process had been finished. The stress needed to detach a dislocation from an interacting particle – the detachment stress – expressed as sd = sOB 1 − k 2R

(4)

is then identified with the true threshold stress [19,20]. In Eq. (4), sOB is the Orowan bowing stress and kR is the relaxation factor characterising the strength of the attractive dislocation/particle interaction. The Orowan bowing stress sOB can be expressed by the well known simple formula [29] sOB =

0.84 MGb , l− d

(5)

in which M is the Taylor factor, l is the mean interparticle spacing and d is the mean particle diameter. The Orowan bowing stress is thus proportional to the shear modulus G and, consequently, sOB /G does not depend on temperature. The experimentally determined true threshold stress sTH depends on temperature more strongly than the shear modulus G of the matrix. Therefore, sTH cannot be identified with the detachment stress sd unless a temperature dependence of the relaxation factor kR is admitted. At the present time, the relaxation factor cannot be calculated from the first principles. Arzt and Wilkinson [19] modelled the attractive dislocation/particle interaction in a simple way in terms of the dislocation line energies G and G% in the matrix and in the particle matrix interface, respectively. Accepting the model of these authors and taking into account possible role of impurities (decreasing the energy G% and forming atmospheres around dislocations at the departure sides of interacting particles) [26,27] it has been shown [22,30] that a temperature dependence of the relaxation factor can be realistically expected.

An idea on possible values of kR and especially on the temperature dependence of this factor can be obtained analysing proper o; m(T,s) creep data. This was illustrated for an ODS Al-5Mg-30SiCp composite [22], 2124Al-20SiCp composite [23] and also for the Al8.5Fe-1.3V-1.7Si alloy [18]. To get such an idea for the composite investigated in the present work, Eq. (4) is written in the form (c.f. Ref. [22]) sd = C 1 − k 2R G

(6)

where C= sOB/G is a temperature independent constant equal to 0.84 Mb/(l −d), since sOB 8 G, see Eq. (5). Accepting the value of 0.85 for kR at 673 K (see Refs. [12,13]) and setting sd/G=sTH/G=6.146×10 − 3 for this temperature, the constant C= sTH/(G 1 −k 2R) is obtained equal to 1.167×10 − 2. Accepting this value of C, values of kR equal to 0.75, 0.85 and 0.91 are obtained for 623, 673 and 723 K, respectively. These values of the relaxation factor kR and, consequently, also the temperature dependence of this factor seem reasonable. It is worthwhile pointing out that the factor kR apparently approaches a constant value close to the ‘critical’ value of 0.94 with the temperature approaching 800 K, Fig. 9. The ‘critical’ value of kR followed from the analyses of Arzt and Wilkinson [19] and Arzt and Ro¨sler [20]. As pointed out previously by the present authors [22], the above assumption on temperature dependence of the relaxation factor kR seems to be the only, although perhaps still a somewhat speculative explanation of the observed temperature dependence of the normalised true threshold stress, sTH/G, as obtained analysing the experimental o; m(T, s) creep data for discontinuous dispersion strengthened aluminium alloy matrix composites. However, as already mentioned, this conclusion is supported by extensive considerations on the possible role of impurities on dislocations bypassing interacting particles by localised climb and final detachment [13,26,27]. Peng et al. [4] interpreted their results of an investigation of creep in an Al-8.5Fe-1.3V-1.7Si-15SiCw composite at temperatures ranging from 573 to 723 K alternatively in terms of the true threshold stress concept and in terms of the concept of thermally activated detachment of dislocations [31] from fine Al12(Fe,V)3Si phase particles. Both the alternatives were found equally acceptable. The present analysis of o; m(T,s) creep data for the composite under consideration strongly supports the interpretation of its creep behaviour in terms of the former concept. This unambiquity of the interpretation given in the present paper follows from almost twice as broad interval of the measured (clearly defined) minimum creep strain rates as compared to the interval of the creep strain rates measured by Peng et al. [4,5]). In a previous paper [32],

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it was shown that to estimate accurately the true stress exponent n, the measured minimum creep strain rates should cover not less than five orders of magnitude (see also Ref. [33]). Of course, the same holds for the true threshold stress determination when the true stress exponent n is chosen considering a specific creep model.

5. Conclusions In the present work, the creep behaviour of an Al-8.5Fe-1.3V-1.7Si (the 8009 Al type) alloy reinforced with 15 vol.% silicon carbide particulates – a discontinuous Al-8.5Fe-1.3V-1.7Si-15SiCp composite-was investigated at three temperatures ranging from 623 to 723 K using the isothermal constant tensile stress creep test technique. The measured minimum creep strain rates covered seven orders of magnitude. The main results can be listed as follows. 1. The creep behaviour is associated with the true threshold stress sTH decreasing with increasing temperature more strongly than the shear modulus G of aluminium. 2. The minimum creep strain rate o; m is controlled by lattice diffusion in the composite matrix and the true stress exponent is close to 5. 3. The apparent applied stress exponent mc depends strongly on both applied stress and temperature and is generally several times higher than the true stress exponent n= 5. Also the apparent activation energy of creep, Qc, depends strongly on applied stress as well as on temperature, generally being several times higher than the activation enthalpy DHL of lattice self-diffusion in aluminium. This behaviour of mc and Qc is accounted for by the strong temperature dependence of the sTH/G ratio. 4. The threshold creep behaviour of the composite is interpreted in terms of athermal detachment of dislocations from fine incoherent Al12(Fe,V)3Si phase and alumina particles assuming a temperature dependence of the relaxation factor kR, characterising the strength of the attractive dislocation/particle interaction. 5. The true threshold stress of the composite is by a factor ranging from 1.36 to 1.53 (depending on temperature) higher than that of the unreinforced alloy. This is explained by an effect of fine alumina particles appearing in the composite matrix during the composite fabrication by powder metallurgy route. 6. The creep strength of the composite, as characterised by the minimum creep strain rate, is up to six orders of magnitude higher (the minimum creep strain rate is up to six orders of magnitude lower) than that of the Al-8.5Fe-1.3V-1.7Si alloy. This

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creep strengthening is attributed to the higher true threshold stress in the composite than in the unreinforced alloy.

Acknowledgements Two of the authors (J.C& . and K.K.) thank the Institute of Physics of Materials, Academy of Sciences of the Czech Republic for financial support (Project No. 7/96 K). S.J.Z. (on leave from Dalian University of Technology, Dalian, P.R. China) thanks for financial support from the University of ElectroCommunications, Tokyo. The authors are grateful to Eva Najvarova´ for assistance in manuscript preparation.

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