- Email: [email protected]
Creep behaviour of a rapidly solidified Al-5Cr-2Zr alloy between room temperature and 823 K
Scripta Materialia, Vol. 35, No. 12, pp. 1449-1454,1996 Elsevier Science Ltd Copyright 0 1996 Acta Metallurgica Inc. Printed in the USA. All rights reserved 1359-6462/96 $12.00 + .OO
Pergamon
PI1 S1359-6462(96)00318-1
CREEP IBEHAVIOUR OF A RAPIDLY SOLIDIFIED Al-SCr-2Zr ALLtOY BETWEEN ROOM TEMPERATURE AND 823 K A. Brahmi, T. Gerique, M. Lieblich and M. Torralba Dept. Metahlrgia Fisica, Centro National de Investigaciones Metalurgicas (CENIM - CSIC), Avda. Gregorio de1 Amo 8,28040 Madrid, Spain (Received May 2, 1996) (Accepted July 22, 1996) Introduction Rapidly solidified (RS) Al-Cr-Zr alloys are established contenders for applications in the aircraft industry where lower cost, lightweight substitutes for titanium alloys are being sought for use in the temperature range of 473 to 723 K [ 1,2]. In the Al-Cr-Zr family of alloys, Cr provides thermal stability through the precipitation of the AI,$& phase and Zr provides age hardening response through the precipitation of the LIZ-A4Zr phase [3]. Creep resistance is one of the critical properties of any material intended for high temperature applications. Therefore, a precise knowledge of creep behaviour and a clear understanding of the mechanisms controlling creep in these materials are of great importance. The good thermal stability exhibited by the RS Al-SCr-2Zr (wt.%) alloy [4,5] makes it a good candidate for applications where high creep resistance is needed. This paper presents the results of creep behaviour over a wide range of temperatures (0.32 to 0.88 Tm, where Tm = 933 K is the melt temperature of pure aluminium) of an Al-SCr2Zr alloy processed by gas atomization and extrusion and includes a brief discussion on the creep mechanisms that may be involved. Comprehensive analyses will be given in subsequent papers. Experimental Procedure The Al-SCr-2Zr (wt.%) alloy for study was obtained through argon atomization followed by hot extrusion. Atomization was carried out with a confined nozzle atomizer equipped with an induction furnace capable of heating up to about 2000 K. For the production of the Al-SCr-2Zr powder, melt superheats between 140 and 300 K were used and argon pressures of 2.5 + 0.1 MPa were applied. A detailed description of lthe atomization process and characterization of the powder were reported earlier [6]. Powder particles, mainly spherical in shape, with diameters under 25 pm were selected and canned. The Al-Cr-Zr powder was degassed by slowly increasing the temperature up to 773 K, when a final vacuum of 5 x lo-* Pa was achieved. Extrusion was performed on a horizontal hot extrusion press capable of imp,arting extrusion pressures up to about 1600 MPa at temperatures up to 773 K and ram speeds varying between 0.3 and 12 mm s-‘. Filled cans were preheated at 723 K for one hour. Extru-
1449
1450
CREEPBEHAVIOUR
Vol. 35, No. 12
Figure 1. Strain rate verws Young’s modulus compensated flow stress.
sion was conducted at 723 K at a ram speed of 1 mm s“. A circular, cross-section die of 8 mm in diameter with a conical angle of 90” was employed which provided an extrusion ratio of 28: 1. Cylindrical tensile specimens of 4.98 mm in diameter and a gauge length of 22.3 mm were machined from the extruded bar while maintaining the tensile axis parallel to the direction of extrusion. Then, the specimens were heat treated at 573 K for 20 minutes to release internal stresses caused during the extrusion process. Creep behaviour was studied by means of tensile strain-rate-change tests performed in air at temperatures ranging from 298 to 823 K and strain rates ranging from 10S6to 10-l s-‘. Results
Figure 1 shows a log-log plot of creep rate E versus Young’s modulus compensated steady state flow stress o/E at temperatures ranging between 298 and 823 K. The E used was that of pure aluminium [7]. Over the whole temperature range investigated, the flow stress increases as strain rate increases and temperature decreases. The data obtained for each temperature fall on a straight line. This permits the description of the flow behaviour of Al-SCr-2Zr by the power law, constitutive equation: ”
where A is a material constant, n is the stress exponent given by the slope of each line and D = Do exp (-Qc/RT) is a diffusion coefficient with Do the frequency factor for diffusion, Qc the activation energy for creep, R the universal gas constant and T the absolute temperature. The calculated values of n for each testing temperature are represented in Figure 2. The stress exponent n has a value of 21 at 298 K, increases to 34 at 423 K, decreases again to 21 at 523 K and has a value between 7 and 10 at 573 to 823 K. At this higher temperature range, given that a constant stress exponent is assumed, activation energies for creep between consecutive temperature pairs can be determined applying equation (1). These are also represented in Figure 2, where the horizontal bars indicate the temperature range for the evaluation of Qc and the vertical bars indicate the range of the values due to the fact that the lines of Figure 1 are not exactly parallel. For temperature pairs between
Vol. 35, No. 12
CREEP BEHAVIOUR
1451
Figure 2. Stress exponent and activation energy versus testing temperature.
573 and 673 K activation energies fall on the range of 135 to 195 kJ mol’, whereas between 673 and 823 K these are in the range of 225 to 300 kJ mol-‘. Figure 3 shows data of Figure 1 replotted as creep rate divided by the effective diffusion coefficient, D,ff. The effective aluminium self-diffusion coefficient, Da, is introduced instead of D in equation (1) to cover the wjlde range of temperatures spanned by the data (0.32 to 0.88 Tm) [8]. D,E includes the additive contributions of the lattice diffusion coefficient DL, DL=l .7 lOa exp(- 142 kJ mol’/RT) m2sW’ [9], and the dislocation pipe diffusion coefficient D,, D,=2.8 1Oaexp(-82 kJ mol’/RT) m2s“ [lo]: I&= fL DL
+ fp D,
(2)
where fL (= 1 - fp = 1) and fly (= 200 ( o/E)~ [ 111) are the fractions of atoms participating in lattice and pipe diffusion, respectively. Figure 3 shows that the data obtained at temperatures below 523 K (0.56 Tm) merge well into a single straight line of mean slope 21 and that for higher temperatures several parallel lines are obtained with a mean slope of 7.5. Discussion The interpretation of the results presented above can be attempted in various ways depending on the
election of the temperature ranges for which distinct, diffusion- controlled, dislocation creep mechanisms would be assumed. Although further studies on creep behaviour of Al-SCr-2Zr are in progress to fully understand the mechanisms responsible for creep and the microstructural features involved, and which will be the subject of future papers, one possible interpretation is presented below. Based on the stress sensitivity of the Al-SCr-2Zr alloy, Figure 2, two temperature ranges can be defined: a high sensitivity region with n > 20 from 298 to 523 K (0.32 to 0.56 Tm), and a low sensitivity region with n close to 8 from 573 to 823 K (0.61 to 0.88 Tm). The latter region can again be divided into two regions based on the results of activation energies calculated between consecutive temperature pairs, Figure 2. In each of these regions, between 573 and 673 K (0.61 to 0.72 Tm), and between 723 and 823 K (0.77 to 0.88 Tm), this alloy would behave differently. For these three temperature ranges Qc was recalculated at constant o/E from the Arrhenius plot of log t? versus TM’presented in Figure 4. From the slopes of the three lines, apparent activation energies of 77, 167 and 273 kJ mol’ for the low, intermediate and high temperature ranges, respectively, are obtained. Once the temperature ranges have been defined, a brief analysis of the mechanisms governing flow of Al-SCr2Zr follows.
1452
CREEP BEHAVIOUR
Vol. 35, No. 12
IO" -7 .5
IO” f
10“
_ -
,o”
_
0 a 0 .
n ‘W
. q
0 lo’e
q
423 473 573 513 623 673 72.3 773 823
-
Figure 3. Effective diffusion compensated strain rate versus Young’s modulus compensated flow stress (fp = 200(0/E)~).
(i) Low temperature range (298-523 K): n > 20 and Qc = 77 kJ mol’. This value of Qc is close to that of the activation energy for pipe diffusion of aluminium, Qp = 82 kJ mol [9]. According to the literature [10,12], pipe diffusion is the rate controlling mechanism for dislocation climb in pure aluminium at temperatures below 0.6 Tm. Figure 3 indicates that, as all the low temperature data merge into a single straight line, Al pipe diffusion is also rate controlling for dislocation creep in Al-SCr-2Zr at these temperatures (DL is negligible compared to fP Dp). Opposite to what occurs with the activation energy encountered, the high value of n cannot be associated with any diffusion-controlled dislocation creep mechanism. High stress sensitivity is commonly found in the literature for ODS alloys, see for example ref. [13], and other dispersion-strengthened aluminium alloys, see for example ref. [ 141. This problem has been faced [ 151 by introducing a tbreshold stress (TVwhich is related to the critical stress for the beginning of plastic deformation, or the yielding stress associated with dispersed particles. As was revealed by hardness measurements [4, 51, in this temperature range the substructure of the Al-SCr-2Zr alloy can be considered invariant. In the invariant substructure model of Sherby et al. [ 12,131, a stress exponent of 8 is proposed when creep is lattice diffusion controlled. However, when pipe diffusion dominates the deformation mechanism, as is lo“ -
“=Y
w md’ . .
OIE -3.2ox1cf’
IO” -
-
1.66Xl(T’
-4.4oxl~ . . qc-77wmr’ IO”
J
-
.
\\ 10“
1 1
I
2
3 I-’
(x
ld
K”)
Figure 4. Arrhenius plot of the steady-state creep versus lO’/T.
4
Vol. 35, No. 12
CREEP BEHAVIOUR
&-
ld’
-
ld’
-
Id’
-
10”
-
1453
3 cl.W
10” -
1
Id
lo*
la
1Q’
(a*oYE Figure 5. Pipe diffusion compensated strain rate versus Young’s modulus compensated effective stress (a-a,,)/E from 298 to 523 K.
the present case!, this stress exponent increased by 2, i.e. n=lO, since fp, and thus Dee, are proportional to (o/E)’ (equation 2). Consequently, assuming the existence of croin the present alloy, the stress exponent was taken equal to 10. Then, using an established procedure [ 16,171, a normalized threshold stress was determined for each temperature, ranging from 2.60~10” at 298 K to 1.43~10” at 523 K. Figure 5 shows the plot of &/Dp versus (a-oo)/E. As can be seen, all the data merge into a straight line except those for the highest temperature of 523 K. This indicates that the existence of a threshold stress would explain quite satisfactorily the creep behaviour of this alloy at least up to 473 K, with 523 K as a transition temperature between the low and the intermediate temperature ranges. (ii) Intermediate temperature range (573-673 K): n close to 8 and Qc = 167 kJ mole’. This value of (Qc is close to that of activation energy for Al lattice self-diffusion (QL = 142 kJ mol-’ [9]). According to the literature [ 14,17-2 11, creep of many aluminium-dispersion-strengthened alloys in this temperature range has been found to be controlled by Al lattice self-diffusion. The present values of n close to 8 and Qc close to that of activation energy for Al lattice self-diffusion are consistent with the substructure invariant creep model of dislocation climb [12]. In Figure 3, plotted for a va.lue of fp = 200 (o/E)*, the data for 573 K appear separated from those for 623 and 673 K. This is because the maximum contribution of pipe diffusion, i.e. the maximum value of fp [ 111, was considered in equation (2). However, given that Qc is close to QL, this contribution should in fact be lower. For lower values of fp, the data approach each other, indicating that dislocation climb seems to be the controlling mechanism in this temperature range. (iii) High temperature range (723-823 K): n close to 8 and Qc = 273 kJ mo1.l. Although the stress exponent suggests that creep may be explained by the substructure invariant model, the activation energy is much higher than the activation energy for self-diffusion of pure aluminium. As can be seen in Figure 3, a different parallel straight line is obtained for each testing temperature. This could be attributed [lo] to the coarsening of dispersoids expected at temperatures higher than the extrusion temperature (723 K). According to the literature, high apparent Qc may also be expected to arise from loss of effectiveness in the stress field of the dislocation networks surrounding the dispersoids 1[20],plastic deformation of the dispersoids along with the matrix [22], or decrease in the diffusion coefficient due to alloying [18]. Any one or a combination of these possibilities could account for the high apparent activation energy for creep obtained. These will be examined in detail in a future study.
1454
CREEP BEHAVIOUR
Vol. 35, No. 12
Summary and Conclusions The flow behaviour of the rapidly solidified Al-SCr-2Zr alloy investigated in this study is well described by the power law creep. Based on the stress sensitivity and the activation energy for creep encountered, three temperature ranges were defined: a low temperature range from 298 to 523 K with n > 20 and Qc = 77 kJ mole’, an intermediate temperature range from 573 to 673 K with n close to 8 and Qc = 167 kJ mol’, and a high temperature range from 723 to 823 K with n close to 8 and Qc = 273 kJ moI’. In the low temperature range, creep rate is controlled by aluminium pipe diffusion. The high value of n > 20 could not be associated with any diffusion controlled dislocation creep mechanism. This behaviour can be rationalized using a substructure invariant model with a stress exponent of 10 and a threshold stress. In the intermediate temperature range, creep rate is controlled by aluminium lattice self-diffusion. The stress exponent and the activation energy obtained are consistent with the substructure invariant creep model. In the high temperature range, although the stress exponent suggests that creep may be explained by the substructure invariant model, the apparent activation energy is much higher than the activation energy for self-diffusion of pure ahnninium, which may be attributed to the coarsening of dispersoids. Acknowledgements Financial support of MAT 94-0779 is greatly acknowledged. Thanks are also due to Prof. O.A. Ruano for helpful discussion, and to G. Caruana and B.J. Ferniindez for experimental assistance. One of the authors (A.B.) would like to thank the Commission of the European Union for a post-doctoral fellowship at CENIM in the framework of the Human Capital and Mobility program. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11, 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22.
J.R. Pickens, High strength P/M Al alloys, p. 200, Metals Handbook, 10th ed., vol. 2, ASM, Metals Park, Ohio (1990). E.J. Lavemia, J.D. Ayers, and T.S. Srivatsan, Intemat. Mater. Reviews 37, 1 (1992). S.P. Midson, H. Jones, and R.A. Buckley, Rapidly Quenched Metals, p. 923, Elsevier, New York (1985). M. Lieblich, G. Caruana, M. Torralba, and H. Jones, Intemat. J. Rapid Solid. 9,45 (1995). G. Caruana, M. Lieblich, M. Torralba, and H. Jones, J. Mater. Sci. Letters. 15,568 (1996). M. Lieblich, G. Caruana, M. Torralba, and H. Jones, Mater. Sci. Technol. 12,25 (1996). W. Koster, Z. Metallkd. 39, 1 (1948). S.L. Robinson and O.D. Sherby, Acta Met. 17, 109 (1969). H.J. Frost and M.F. Ashby, Deformation Mechanism Maps, p. 21, Pergamon Press, London (1982). H. Luthy, A.K. Miller, and O.D. Sherby, Acta Metall. 28, 169 (1980). O.D. Sherby and C.M. Young, Rate Processes in Plastic Deformation of Materials, p. 497, Proceedings from the John E. Dom Symposium, ASM, (1975). O.D. Sherby, R.H. Klundt, and A.K. Miller, Metall. Trans. SA, 843 (1977). J. Lin and O.D. Sherby, Res. Mechanica2,251(1981). F. Carreflo,G. Gonzalez-Doncel, and O.A. Ruano, Mater. Sci. Eng. A&4,217 (1993). K.R. Williams and B. Wilshire, Metal Sci. J. 7, 176 (1973). A.B. Pandey, R.S. Mishra, and Y.R. Mahajan, Metal. Mater. Trans. 27A, 305 (1996). R.S. Mishra, A.J. Paradkar, and K.N. Rao, ActaMetaR. Mater. 41,2243 (1993). D. Legzdina and T.A. Parthasarathy, Metall. Trans. 18A, 1713 (1987). D. Legzdina and T.A. Parthasarathy, Metall. Trans. 21A, 2155 (1990). J. Wolfenstine, G. Gonzalez-Doncel, and O.D. Sherby, J. Mater. Sci. Letters 9,410 (1990). Y.H. Yeh, H. Nakashima, H. Kurishita, G. Goto, and H. Yoshinaga, Mater. Sci. JIM 31,284 (1990). D.L. Yaney and D. Nix, Metall. Trans. 18A, 1713 (1987).
Pergamon
PI1 S1359-6462(96)00318-1
CREEP IBEHAVIOUR OF A RAPIDLY SOLIDIFIED Al-SCr-2Zr ALLtOY BETWEEN ROOM TEMPERATURE AND 823 K A. Brahmi, T. Gerique, M. Lieblich and M. Torralba Dept. Metahlrgia Fisica, Centro National de Investigaciones Metalurgicas (CENIM - CSIC), Avda. Gregorio de1 Amo 8,28040 Madrid, Spain (Received May 2, 1996) (Accepted July 22, 1996) Introduction Rapidly solidified (RS) Al-Cr-Zr alloys are established contenders for applications in the aircraft industry where lower cost, lightweight substitutes for titanium alloys are being sought for use in the temperature range of 473 to 723 K [ 1,2]. In the Al-Cr-Zr family of alloys, Cr provides thermal stability through the precipitation of the AI,$& phase and Zr provides age hardening response through the precipitation of the LIZ-A4Zr phase [3]. Creep resistance is one of the critical properties of any material intended for high temperature applications. Therefore, a precise knowledge of creep behaviour and a clear understanding of the mechanisms controlling creep in these materials are of great importance. The good thermal stability exhibited by the RS Al-SCr-2Zr (wt.%) alloy [4,5] makes it a good candidate for applications where high creep resistance is needed. This paper presents the results of creep behaviour over a wide range of temperatures (0.32 to 0.88 Tm, where Tm = 933 K is the melt temperature of pure aluminium) of an Al-SCr2Zr alloy processed by gas atomization and extrusion and includes a brief discussion on the creep mechanisms that may be involved. Comprehensive analyses will be given in subsequent papers. Experimental Procedure The Al-SCr-2Zr (wt.%) alloy for study was obtained through argon atomization followed by hot extrusion. Atomization was carried out with a confined nozzle atomizer equipped with an induction furnace capable of heating up to about 2000 K. For the production of the Al-SCr-2Zr powder, melt superheats between 140 and 300 K were used and argon pressures of 2.5 + 0.1 MPa were applied. A detailed description of lthe atomization process and characterization of the powder were reported earlier [6]. Powder particles, mainly spherical in shape, with diameters under 25 pm were selected and canned. The Al-Cr-Zr powder was degassed by slowly increasing the temperature up to 773 K, when a final vacuum of 5 x lo-* Pa was achieved. Extrusion was performed on a horizontal hot extrusion press capable of imp,arting extrusion pressures up to about 1600 MPa at temperatures up to 773 K and ram speeds varying between 0.3 and 12 mm s-‘. Filled cans were preheated at 723 K for one hour. Extru-
1449
1450
CREEPBEHAVIOUR
Vol. 35, No. 12
Figure 1. Strain rate verws Young’s modulus compensated flow stress.
sion was conducted at 723 K at a ram speed of 1 mm s“. A circular, cross-section die of 8 mm in diameter with a conical angle of 90” was employed which provided an extrusion ratio of 28: 1. Cylindrical tensile specimens of 4.98 mm in diameter and a gauge length of 22.3 mm were machined from the extruded bar while maintaining the tensile axis parallel to the direction of extrusion. Then, the specimens were heat treated at 573 K for 20 minutes to release internal stresses caused during the extrusion process. Creep behaviour was studied by means of tensile strain-rate-change tests performed in air at temperatures ranging from 298 to 823 K and strain rates ranging from 10S6to 10-l s-‘. Results
Figure 1 shows a log-log plot of creep rate E versus Young’s modulus compensated steady state flow stress o/E at temperatures ranging between 298 and 823 K. The E used was that of pure aluminium [7]. Over the whole temperature range investigated, the flow stress increases as strain rate increases and temperature decreases. The data obtained for each temperature fall on a straight line. This permits the description of the flow behaviour of Al-SCr-2Zr by the power law, constitutive equation: ”
where A is a material constant, n is the stress exponent given by the slope of each line and D = Do exp (-Qc/RT) is a diffusion coefficient with Do the frequency factor for diffusion, Qc the activation energy for creep, R the universal gas constant and T the absolute temperature. The calculated values of n for each testing temperature are represented in Figure 2. The stress exponent n has a value of 21 at 298 K, increases to 34 at 423 K, decreases again to 21 at 523 K and has a value between 7 and 10 at 573 to 823 K. At this higher temperature range, given that a constant stress exponent is assumed, activation energies for creep between consecutive temperature pairs can be determined applying equation (1). These are also represented in Figure 2, where the horizontal bars indicate the temperature range for the evaluation of Qc and the vertical bars indicate the range of the values due to the fact that the lines of Figure 1 are not exactly parallel. For temperature pairs between
Vol. 35, No. 12
CREEP BEHAVIOUR
1451
Figure 2. Stress exponent and activation energy versus testing temperature.
573 and 673 K activation energies fall on the range of 135 to 195 kJ mol’, whereas between 673 and 823 K these are in the range of 225 to 300 kJ mol-‘. Figure 3 shows data of Figure 1 replotted as creep rate divided by the effective diffusion coefficient, D,ff. The effective aluminium self-diffusion coefficient, Da, is introduced instead of D in equation (1) to cover the wjlde range of temperatures spanned by the data (0.32 to 0.88 Tm) [8]. D,E includes the additive contributions of the lattice diffusion coefficient DL, DL=l .7 lOa exp(- 142 kJ mol’/RT) m2sW’ [9], and the dislocation pipe diffusion coefficient D,, D,=2.8 1Oaexp(-82 kJ mol’/RT) m2s“ [lo]: I&= fL DL
+ fp D,
(2)
where fL (= 1 - fp = 1) and fly (= 200 ( o/E)~ [ 111) are the fractions of atoms participating in lattice and pipe diffusion, respectively. Figure 3 shows that the data obtained at temperatures below 523 K (0.56 Tm) merge well into a single straight line of mean slope 21 and that for higher temperatures several parallel lines are obtained with a mean slope of 7.5. Discussion The interpretation of the results presented above can be attempted in various ways depending on the
election of the temperature ranges for which distinct, diffusion- controlled, dislocation creep mechanisms would be assumed. Although further studies on creep behaviour of Al-SCr-2Zr are in progress to fully understand the mechanisms responsible for creep and the microstructural features involved, and which will be the subject of future papers, one possible interpretation is presented below. Based on the stress sensitivity of the Al-SCr-2Zr alloy, Figure 2, two temperature ranges can be defined: a high sensitivity region with n > 20 from 298 to 523 K (0.32 to 0.56 Tm), and a low sensitivity region with n close to 8 from 573 to 823 K (0.61 to 0.88 Tm). The latter region can again be divided into two regions based on the results of activation energies calculated between consecutive temperature pairs, Figure 2. In each of these regions, between 573 and 673 K (0.61 to 0.72 Tm), and between 723 and 823 K (0.77 to 0.88 Tm), this alloy would behave differently. For these three temperature ranges Qc was recalculated at constant o/E from the Arrhenius plot of log t? versus TM’presented in Figure 4. From the slopes of the three lines, apparent activation energies of 77, 167 and 273 kJ mol’ for the low, intermediate and high temperature ranges, respectively, are obtained. Once the temperature ranges have been defined, a brief analysis of the mechanisms governing flow of Al-SCr2Zr follows.
1452
CREEP BEHAVIOUR
Vol. 35, No. 12
IO" -7 .5
IO” f
10“
_ -
,o”
_
0 a 0 .
n ‘W
. q
0 lo’e
q
423 473 573 513 623 673 72.3 773 823
-
Figure 3. Effective diffusion compensated strain rate versus Young’s modulus compensated flow stress (fp = 200(0/E)~).
(i) Low temperature range (298-523 K): n > 20 and Qc = 77 kJ mol’. This value of Qc is close to that of the activation energy for pipe diffusion of aluminium, Qp = 82 kJ mol [9]. According to the literature [10,12], pipe diffusion is the rate controlling mechanism for dislocation climb in pure aluminium at temperatures below 0.6 Tm. Figure 3 indicates that, as all the low temperature data merge into a single straight line, Al pipe diffusion is also rate controlling for dislocation creep in Al-SCr-2Zr at these temperatures (DL is negligible compared to fP Dp). Opposite to what occurs with the activation energy encountered, the high value of n cannot be associated with any diffusion-controlled dislocation creep mechanism. High stress sensitivity is commonly found in the literature for ODS alloys, see for example ref. [13], and other dispersion-strengthened aluminium alloys, see for example ref. [ 141. This problem has been faced [ 151 by introducing a tbreshold stress (TVwhich is related to the critical stress for the beginning of plastic deformation, or the yielding stress associated with dispersed particles. As was revealed by hardness measurements [4, 51, in this temperature range the substructure of the Al-SCr-2Zr alloy can be considered invariant. In the invariant substructure model of Sherby et al. [ 12,131, a stress exponent of 8 is proposed when creep is lattice diffusion controlled. However, when pipe diffusion dominates the deformation mechanism, as is lo“ -
“=Y
w md’ . .
OIE -3.2ox1cf’
IO” -
-
1.66Xl(T’
-4.4oxl~ . . qc-77wmr’ IO”
J
-
.
\\ 10“
1 1
I
2
3 I-’
(x
ld
K”)
Figure 4. Arrhenius plot of the steady-state creep versus lO’/T.
4
Vol. 35, No. 12
CREEP BEHAVIOUR
&-
ld’
-
ld’
-
Id’
-
10”
-
1453
3 cl.W
10” -
1
Id
lo*
la
1Q’
(a*oYE Figure 5. Pipe diffusion compensated strain rate versus Young’s modulus compensated effective stress (a-a,,)/E from 298 to 523 K.
the present case!, this stress exponent increased by 2, i.e. n=lO, since fp, and thus Dee, are proportional to (o/E)’ (equation 2). Consequently, assuming the existence of croin the present alloy, the stress exponent was taken equal to 10. Then, using an established procedure [ 16,171, a normalized threshold stress was determined for each temperature, ranging from 2.60~10” at 298 K to 1.43~10” at 523 K. Figure 5 shows the plot of &/Dp versus (a-oo)/E. As can be seen, all the data merge into a straight line except those for the highest temperature of 523 K. This indicates that the existence of a threshold stress would explain quite satisfactorily the creep behaviour of this alloy at least up to 473 K, with 523 K as a transition temperature between the low and the intermediate temperature ranges. (ii) Intermediate temperature range (573-673 K): n close to 8 and Qc = 167 kJ mole’. This value of (Qc is close to that of activation energy for Al lattice self-diffusion (QL = 142 kJ mol-’ [9]). According to the literature [ 14,17-2 11, creep of many aluminium-dispersion-strengthened alloys in this temperature range has been found to be controlled by Al lattice self-diffusion. The present values of n close to 8 and Qc close to that of activation energy for Al lattice self-diffusion are consistent with the substructure invariant creep model of dislocation climb [12]. In Figure 3, plotted for a va.lue of fp = 200 (o/E)*, the data for 573 K appear separated from those for 623 and 673 K. This is because the maximum contribution of pipe diffusion, i.e. the maximum value of fp [ 111, was considered in equation (2). However, given that Qc is close to QL, this contribution should in fact be lower. For lower values of fp, the data approach each other, indicating that dislocation climb seems to be the controlling mechanism in this temperature range. (iii) High temperature range (723-823 K): n close to 8 and Qc = 273 kJ mo1.l. Although the stress exponent suggests that creep may be explained by the substructure invariant model, the activation energy is much higher than the activation energy for self-diffusion of pure aluminium. As can be seen in Figure 3, a different parallel straight line is obtained for each testing temperature. This could be attributed [lo] to the coarsening of dispersoids expected at temperatures higher than the extrusion temperature (723 K). According to the literature, high apparent Qc may also be expected to arise from loss of effectiveness in the stress field of the dislocation networks surrounding the dispersoids 1[20],plastic deformation of the dispersoids along with the matrix [22], or decrease in the diffusion coefficient due to alloying [18]. Any one or a combination of these possibilities could account for the high apparent activation energy for creep obtained. These will be examined in detail in a future study.
1454
CREEP BEHAVIOUR
Vol. 35, No. 12
Summary and Conclusions The flow behaviour of the rapidly solidified Al-SCr-2Zr alloy investigated in this study is well described by the power law creep. Based on the stress sensitivity and the activation energy for creep encountered, three temperature ranges were defined: a low temperature range from 298 to 523 K with n > 20 and Qc = 77 kJ mole’, an intermediate temperature range from 573 to 673 K with n close to 8 and Qc = 167 kJ mol’, and a high temperature range from 723 to 823 K with n close to 8 and Qc = 273 kJ moI’. In the low temperature range, creep rate is controlled by aluminium pipe diffusion. The high value of n > 20 could not be associated with any diffusion controlled dislocation creep mechanism. This behaviour can be rationalized using a substructure invariant model with a stress exponent of 10 and a threshold stress. In the intermediate temperature range, creep rate is controlled by aluminium lattice self-diffusion. The stress exponent and the activation energy obtained are consistent with the substructure invariant creep model. In the high temperature range, although the stress exponent suggests that creep may be explained by the substructure invariant model, the apparent activation energy is much higher than the activation energy for self-diffusion of pure ahnninium, which may be attributed to the coarsening of dispersoids. Acknowledgements Financial support of MAT 94-0779 is greatly acknowledged. Thanks are also due to Prof. O.A. Ruano for helpful discussion, and to G. Caruana and B.J. Ferniindez for experimental assistance. One of the authors (A.B.) would like to thank the Commission of the European Union for a post-doctoral fellowship at CENIM in the framework of the Human Capital and Mobility program. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11, 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22.
J.R. Pickens, High strength P/M Al alloys, p. 200, Metals Handbook, 10th ed., vol. 2, ASM, Metals Park, Ohio (1990). E.J. Lavemia, J.D. Ayers, and T.S. Srivatsan, Intemat. Mater. Reviews 37, 1 (1992). S.P. Midson, H. Jones, and R.A. Buckley, Rapidly Quenched Metals, p. 923, Elsevier, New York (1985). M. Lieblich, G. Caruana, M. Torralba, and H. Jones, Intemat. J. Rapid Solid. 9,45 (1995). G. Caruana, M. Lieblich, M. Torralba, and H. Jones, J. Mater. Sci. Letters. 15,568 (1996). M. Lieblich, G. Caruana, M. Torralba, and H. Jones, Mater. Sci. Technol. 12,25 (1996). W. Koster, Z. Metallkd. 39, 1 (1948). S.L. Robinson and O.D. Sherby, Acta Met. 17, 109 (1969). H.J. Frost and M.F. Ashby, Deformation Mechanism Maps, p. 21, Pergamon Press, London (1982). H. Luthy, A.K. Miller, and O.D. Sherby, Acta Metall. 28, 169 (1980). O.D. Sherby and C.M. Young, Rate Processes in Plastic Deformation of Materials, p. 497, Proceedings from the John E. Dom Symposium, ASM, (1975). O.D. Sherby, R.H. Klundt, and A.K. Miller, Metall. Trans. SA, 843 (1977). J. Lin and O.D. Sherby, Res. Mechanica2,251(1981). F. Carreflo,G. Gonzalez-Doncel, and O.A. Ruano, Mater. Sci. Eng. A&4,217 (1993). K.R. Williams and B. Wilshire, Metal Sci. J. 7, 176 (1973). A.B. Pandey, R.S. Mishra, and Y.R. Mahajan, Metal. Mater. Trans. 27A, 305 (1996). R.S. Mishra, A.J. Paradkar, and K.N. Rao, ActaMetaR. Mater. 41,2243 (1993). D. Legzdina and T.A. Parthasarathy, Metall. Trans. 18A, 1713 (1987). D. Legzdina and T.A. Parthasarathy, Metall. Trans. 21A, 2155 (1990). J. Wolfenstine, G. Gonzalez-Doncel, and O.D. Sherby, J. Mater. Sci. Letters 9,410 (1990). Y.H. Yeh, H. Nakashima, H. Kurishita, G. Goto, and H. Yoshinaga, Mater. Sci. JIM 31,284 (1990). D.L. Yaney and D. Nix, Metall. Trans. 18A, 1713 (1987).