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Order–disorder transformation in the Ni–4.49 at.% Al alloy
Journal of Alloys and Compounds 414 (2006) 55–59
Order–disorder transformation in the Ni–4.49 at.% Al alloy A.T. Adorno ∗ , W. G´arlipp, M. Cilense, R.A.G. Silva Departamento de F´ısico-Qu´ımica, Instituto de Qu´ımica-UNESP, Caixa Postal 355, 14801-970 Araraquara, SP, Brazil Received 1 June 2005; accepted 29 June 2005 Available online 24 August 2005
Abstract The order–disorder transformation in the Ni–4.49 at.% Al alloy was studied using electrical resistivity measurements, microhardness measurements, differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). The results confirmed the ordering behavior expected for Ni–Al dilute alloys and the suggested relation between resistivity changes and microhardness changes with antiferromagnetic spin ordering. The higher value obtained for the activation energy of vacancy migration was associated with a decrease in the Al concentration gradient near solute-depleted regions. © 2005 Elsevier B.V. All rights reserved. Keywords: Electrical properties; Metals and alloys; Atomic order transformation
1. Introduction Ni–Al alloys have properties that make them candidates for a new generation of high performance alloys and also exhibit phenomena, which are of great theoretical interest [1]. Many high temperature properties required in Ni-based superalloys are achieved by the precipitation of an ordered phase (structure L1) of the Ni3 X (X = Al, Si, etc.) in the fcc phase of these alloys. The morphology of the precipitates related to the precipitate–matrix misfit and the kinetics of precipitation has been investigated. These are usually determined in terms of short-range order, long-range order, clustering or Guinier–Preston Zones. These ordering parameters are determined directly from X-rays or neutron diffraction experiments, and may be discussed qualitatively by measuring the variation of hardness, resistivity and magnetic properties of these alloys with temperature [2]. Williams [3] suggested the existence of short-range order in supersaturated Ni–Al solid solutions after electrical resistivity and hardness studies of isothermally aged sample. Ayub et al. [2] studied the ordering behavior of Ni–5 at.% Al alloy and observed that the tensile strength is related to a magnetic transformation at 220 ◦ C. They also observed that the local order phase changes significantly affect the ∗
Corresponding author. E-mail address: [email protected] (A.T. Adorno).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.07.020
mechanical properties of this alloy. Kozubski et al. [4] studied the “order–order” kinetics in a single crystal of B2-type ordered  -Ni50.5 Al49.5 by means of resistometry at temperatures between 961 and 1105 K. Reversible isothermal relaxations of the electrical resistivity yielded the activation energy EA = 2.3 ± 0.1 eV, the same for ordering and disordering processes and markedly lower than the activation energy for Ni diffusion in NiAl. Cserh´ati et al. [5] performed interdiffusion experiments between different Ni–Al two-phase alloys and found that diffusion of the minor element occurs through a vacancy mechanism in Ni3 Al. Afyouni et al. [6] studied the SRO relaxation in the Ni–6 at.% Al alloy using 4K resistometry and found Em = 1.4 ± 0.3 eV and Ef = 0.9 ± 0.3 eV, respectively, for the mobility and formation of vacancies. In this work, the effect of the Ni–Al interaction on the mechanical properties and on the ordering kinetics of the Ni–4.49 at.% Al alloy was studied, using electrical resistivity measurements, microhardness measurements, differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). 2. Experimental procedure Ni–4.49 at.% Al alloy was prepared in an induction furnace under argon atmosphere using highly pure Ni and Al as starting materials. Results from chemical analysis indicated
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A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
a final alloy composition very close to the nominal one, with Co, Cu, Zn, Fe and Mn as main impurities (concentration less than 100 ppm). Small pieces of 5.00 mm thickness were cold rolled to give samples with 200 mm length, 4.00 mm large and 0.50 mm thickness. These samples were annealed at 1373 K under argon atmosphere for homogenization. The electrical resistivity measurements were made using the four-probe method, a K-3 Leeds & Northrup potentiometer and a HP 6632-A power supply. The samples were equilibrated at 1373 K under argon atmosphere for 2 h and then quenched in water at 273 K. The samples were then heated in successive steps of 25 K for 20 min in each step, up to 925 K and quenched in liquid nitrogen, for the isochronal residual resistivity measurements. The isothermal resistivity changes with time were measured with the samples in a bath of liquid nitrogen. These samples, quenched from 1373 K in cold water, were then aged at the temperatures of 668, 688, 701 and 718 K, chosen from the isochronic curve.
3. Results and discussion Fig. 1 shows the plot of isochronic resistivity changes versus quenching temperature and in this curve it is possible to observe four important features: in the range from 350 to 400 K, a resistivity increase; from 400 to 600 K, a small resistivity peak at 493 K; from 600 to 750 K, a quite intense resistivity peak at 670 K; from 800 to 900 K, a small resistivity decrease. The resistivity increase from 350 to 400 K corresponds to the order restored by the quenched-in vacancies before their elimination at the sinks and the intense peak at 670 K corresponds to the order restored by newly created equilibrium vacancies [6]. The small resistivity peak at 493 K occurs at the same temperature in which a hardness change was observed in dilute Ni–Al alloys [3]. This change was attributed to anti-
Fig. 1. Isochronic resistivity changes with quenching temperature.
ferromagnetic spin ordering. When a unit dislocation passes through an array of spins, an anti-ferromagnetic domain wall is produced, and this requires energy. Especially, if the spin order is imperfect, so that pairing of dislocation would not suffice to overcome this energy barrier, one might expect a resultant hardening. It was suggested that this mechanism is a possible contributor to the strength of iron–chromium alloy and that there was no need for the alloy to be macroscopically anti-ferromagnetic, small regions of anti-ferromagnetic bonding arising perhaps from local segregation of solute would suffice. Indeed, small anti-ferromagnetic domains were observed in Ni–Fe alloys at low temperatures. This is a priori a possible interpretation of hardening of the Ni–5 at.% Al alloy below the transition temperature (at 220 ◦ C or 493 K), as pointed out by Ayub et al. [2]. In this way, it is possible to attribute the resistivity peak at 493 K to antiferromagnetic spin ordering in the Ni–4.49 at.% Al alloy. Fig. 2 shows the microhardness changes measurements with quenching temperature and one may observe two steps of hardness decrease, one from 423 to 573 K and other from 573 to 723 K. The first stage seems to confirm what was proposed by Williams [3] and observed in Fig. 1, the anti-ferromagnetic spin ordering. The second stage may be associated to the resistivity peak observed at 670 K in Fig. 1. Fig. 3 shows the transmission electronic micrographs obtained for the Ni–4.49 at.% Al alloy quenched from 670 K (Fig. 3(a)) and from 750 K (Fig. 3(b)), on the resistivity maximum and after it, in the resistivity curve of Fig. 1, together with the corresponding electron diffraction pattern. It is possible to observe a change in the crystalline planes orientation with the heat treatment, probably due to the atomic reordering assisted by vacancy migration. The diffraction pattern in Fig. 3(c) seems to indicate an ordered distribution of the diffracting planes, and in Fig. 3(d) a disordered one, in a distance small enough to give a rounded spots configuration.
Fig. 2. Microhardness changes vs. quenching temperature. Sample initially quenched from 1373 K.
A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
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Fig. 3. Transmission electron micrographs obtained for samples quenched from: (a) 670 K and (b) 750 K. (c and d) Electron diffraction patterns obtained, respectively, at the same conditions.
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A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
Fig. 5. Isothermal resistivity changes with time.
Fig. 4. DSC curve obtained for the Ni–4.49 at.% Al alloy initially quenched from 1373 K. Heating rate 10 K min−1 .
Fig. 4 shows the DSC curve obtained for the alloy initially quenched from 1373 K. In this figure, it is possible to observe two well-defined thermal events, one endothermic peak at 795 K and an exothermic peak at 825 K. It is also possible to observe two other (very small) endothermic peaks, one at about 500 K and other at about 650 K. The two latter peaks may be associated, respectively, to anti-ferromagnetic spin ordering and to the order restored by newly created equilibrium vacancies, as already mentioned in the discussion of Fig. 1. The endothermic peak at 795 K may be related to the beginning of the resistivity decrease observed in the last part of the curve in Fig. 1 and the exothermic peak at 825 K to the end of this decrease. This resistivity decrease seems to be due to the disordering process occurring at higher temperatures, which begins at 795 K and ends at 825 K. This seems to confirm what was observed by Ayub et al. [2], that the equilibrium short-range order states in Ni–Al alloys of composition ranging from 2 to 8 at.% Al are obtained at 400 ◦ C (673 K). The degree of short-range order is found to be maximum at this temperature and decreases significantly with increase in temperature, and the alloys are essentially in disordered state above 700 ◦ C (973 K). These results seem to indicate that the order–disorder phenomena occurring in the Ni–4.49 at.% Al alloy were detected with different accuracy by the various experimental techniques here employed. All the considered events were detected by resistivity measurements, but their intensity was variable, with the most pronounced one corresponding to the maximum of short-range ordering. The microhardness measurements, although with an accuracy far from the resistivity measurements, were able to indicate a relationship between local order phase changes and the alloy mechanical properties. The DSC curve showed that, among all events, the transition with the greater heat content is that corresponding to the beginning of the disordering process.
Considering that SRO kinetics is a local probe of atomic mobility and that what is measured is the mobility of the faster element (Al, in the present case) [6], it is possible to evaluate a value for the activation energy associated to Al diffusion in the Ni –4.49 at.% Al alloy from the resistivity measurements shown in Fig. 1, and to get some important insights about the Al diffusion mechanism and the Ni–Al interaction. Fig. 5 shows the plot of isothermal resistivity changes with time obtained for samples aged at the temperatures of 668, 688, 701 and 718 K, chosen from the isochronic curve in Fig. 1. These isothermal curves suggest a first-order kinetics model, with the SRO relaxation described by Rt − R∞ t (1) = exp − Ro − R ∞ τT where τ T is the relaxation time at each ageing temperature and its temperature dependence follows the Arrhenius law EM τT = τo exp − (2) kT where EM is the migration activation energy for the order–disorder transformation. Using Eq. (1) and Fig. 5, the values of τ T for each ageing temperature were obtained and are shown in Table 1. Fig. 6 shows the plot of ln τ T versus the inverse of absolute temperature, for the τ T values in Table 1. The activation energy value, obtained from the slope of the straight line in this figure, is EM = (1.69 ± 0.09) eV. This value is close to that = (1.4 ± 0.3) eV for the obtained by Afyouni et al. [6], EM Ni–6 at.% Al alloy and the small difference may be attributed to the Al content. As EM is associated with Al diffusion and Table 1 Relaxation times obtained for each ageing temperature Temperature (K)
Relaxation time (min)
668 688 701 718
352 134 115 40
A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
59
4. Conclusions The resistivity measurements confirmed the ordering behavior expected for Ni–Al dilute alloys and the suggested relation between resistivity changes and microhardness changes with anti-ferromagnetic spin ordering. The order–disorder phenomena occurring in the Ni–4.49 at.% Al alloy were detected with different accuracy by the various experimental techniques. All the considered events were detected by resistivity measurements, but their intensity was variable, with the most pronounced one corresponding to the maximum of short-range ordering. The higher value obtained for the activation energy of vacancy migration in the Ni–4.49 at.% Al alloy was associated with a decrease in the Al concentration gradient near solute-depleted regions. Fig. 6. Plot of ln τ vs. 1/T.
this effect is enhanced by the increase of Al concentration, a difference in the activation energy may be expected for the Ni–4.49 at.% Al alloy. Considering that in a dilute substitutional alloy the rate of homogenization is controlled by the diffusion of solute atoms to solute-depleted regions and that the rate of solute atoms diffusion is greater than that for the solvent atoms, the faster Al diffusion in Ni–Al dilute alloys may be connected to the fact that solute atoms can attract vacancies to produce a more than random probability of finding a vacancy next to a solute atom. Al atoms can attract more vacancies because Al atoms are larger than Ni atoms and, due to the greater Al–vacancy binding energy, the vacancy will be unable to “escape” from the solute atom. In this way, the Al–vacancy pair can diffuse through the lattice together. Hence, the small difference in the activation energies obtained for the Ni–6 at.% Al and Ni–4.49 at.% Al alloys may be associated with a decrease in the Al concentration gradient near solute-depleted regions.
Acknowledgement The authors thank Termomecˆanica Brasil S/A for the preparation and chemical analysis of the samples.
References [1] M. Sluiter, P.E.A. Turchi, F.J. Oinski, G.M. Stocks, Mater. Sci. Eng. A152 (1992) 1. [2] G. Ayub, F.A. Khwaja, A. Ul Haq, Z. Ahmad, Acta Metall. Mater. 43 (1995) 1457. [3] R.O. Williams, Trans. Am. Inst. Miner. Eng. 215 (1989) 1026. [4] R. Kozubski, D. Kmiec, E. Partyka, M. Danielewski, Intermetallics 11 (2003) 897. [5] C. Cserh´ati, A. Paul, A.A. Kodentsov, M.J.H. van Dal, F.J.J. van Loo, Intermetallics 11 (2003) 291. [6] M. Afyouni, V. Pierron-Bohnes, M.C. Cadeville, Acta Metall. 37 (9) (1989) 2339.
Order–disorder transformation in the Ni–4.49 at.% Al alloy A.T. Adorno ∗ , W. G´arlipp, M. Cilense, R.A.G. Silva Departamento de F´ısico-Qu´ımica, Instituto de Qu´ımica-UNESP, Caixa Postal 355, 14801-970 Araraquara, SP, Brazil Received 1 June 2005; accepted 29 June 2005 Available online 24 August 2005
Abstract The order–disorder transformation in the Ni–4.49 at.% Al alloy was studied using electrical resistivity measurements, microhardness measurements, differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). The results confirmed the ordering behavior expected for Ni–Al dilute alloys and the suggested relation between resistivity changes and microhardness changes with antiferromagnetic spin ordering. The higher value obtained for the activation energy of vacancy migration was associated with a decrease in the Al concentration gradient near solute-depleted regions. © 2005 Elsevier B.V. All rights reserved. Keywords: Electrical properties; Metals and alloys; Atomic order transformation
1. Introduction Ni–Al alloys have properties that make them candidates for a new generation of high performance alloys and also exhibit phenomena, which are of great theoretical interest [1]. Many high temperature properties required in Ni-based superalloys are achieved by the precipitation of an ordered phase (structure L1) of the Ni3 X (X = Al, Si, etc.) in the fcc phase of these alloys. The morphology of the precipitates related to the precipitate–matrix misfit and the kinetics of precipitation has been investigated. These are usually determined in terms of short-range order, long-range order, clustering or Guinier–Preston Zones. These ordering parameters are determined directly from X-rays or neutron diffraction experiments, and may be discussed qualitatively by measuring the variation of hardness, resistivity and magnetic properties of these alloys with temperature [2]. Williams [3] suggested the existence of short-range order in supersaturated Ni–Al solid solutions after electrical resistivity and hardness studies of isothermally aged sample. Ayub et al. [2] studied the ordering behavior of Ni–5 at.% Al alloy and observed that the tensile strength is related to a magnetic transformation at 220 ◦ C. They also observed that the local order phase changes significantly affect the ∗
Corresponding author. E-mail address: [email protected] (A.T. Adorno).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.07.020
mechanical properties of this alloy. Kozubski et al. [4] studied the “order–order” kinetics in a single crystal of B2-type ordered  -Ni50.5 Al49.5 by means of resistometry at temperatures between 961 and 1105 K. Reversible isothermal relaxations of the electrical resistivity yielded the activation energy EA = 2.3 ± 0.1 eV, the same for ordering and disordering processes and markedly lower than the activation energy for Ni diffusion in NiAl. Cserh´ati et al. [5] performed interdiffusion experiments between different Ni–Al two-phase alloys and found that diffusion of the minor element occurs through a vacancy mechanism in Ni3 Al. Afyouni et al. [6] studied the SRO relaxation in the Ni–6 at.% Al alloy using 4K resistometry and found Em = 1.4 ± 0.3 eV and Ef = 0.9 ± 0.3 eV, respectively, for the mobility and formation of vacancies. In this work, the effect of the Ni–Al interaction on the mechanical properties and on the ordering kinetics of the Ni–4.49 at.% Al alloy was studied, using electrical resistivity measurements, microhardness measurements, differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). 2. Experimental procedure Ni–4.49 at.% Al alloy was prepared in an induction furnace under argon atmosphere using highly pure Ni and Al as starting materials. Results from chemical analysis indicated
56
A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
a final alloy composition very close to the nominal one, with Co, Cu, Zn, Fe and Mn as main impurities (concentration less than 100 ppm). Small pieces of 5.00 mm thickness were cold rolled to give samples with 200 mm length, 4.00 mm large and 0.50 mm thickness. These samples were annealed at 1373 K under argon atmosphere for homogenization. The electrical resistivity measurements were made using the four-probe method, a K-3 Leeds & Northrup potentiometer and a HP 6632-A power supply. The samples were equilibrated at 1373 K under argon atmosphere for 2 h and then quenched in water at 273 K. The samples were then heated in successive steps of 25 K for 20 min in each step, up to 925 K and quenched in liquid nitrogen, for the isochronal residual resistivity measurements. The isothermal resistivity changes with time were measured with the samples in a bath of liquid nitrogen. These samples, quenched from 1373 K in cold water, were then aged at the temperatures of 668, 688, 701 and 718 K, chosen from the isochronic curve.
3. Results and discussion Fig. 1 shows the plot of isochronic resistivity changes versus quenching temperature and in this curve it is possible to observe four important features: in the range from 350 to 400 K, a resistivity increase; from 400 to 600 K, a small resistivity peak at 493 K; from 600 to 750 K, a quite intense resistivity peak at 670 K; from 800 to 900 K, a small resistivity decrease. The resistivity increase from 350 to 400 K corresponds to the order restored by the quenched-in vacancies before their elimination at the sinks and the intense peak at 670 K corresponds to the order restored by newly created equilibrium vacancies [6]. The small resistivity peak at 493 K occurs at the same temperature in which a hardness change was observed in dilute Ni–Al alloys [3]. This change was attributed to anti-
Fig. 1. Isochronic resistivity changes with quenching temperature.
ferromagnetic spin ordering. When a unit dislocation passes through an array of spins, an anti-ferromagnetic domain wall is produced, and this requires energy. Especially, if the spin order is imperfect, so that pairing of dislocation would not suffice to overcome this energy barrier, one might expect a resultant hardening. It was suggested that this mechanism is a possible contributor to the strength of iron–chromium alloy and that there was no need for the alloy to be macroscopically anti-ferromagnetic, small regions of anti-ferromagnetic bonding arising perhaps from local segregation of solute would suffice. Indeed, small anti-ferromagnetic domains were observed in Ni–Fe alloys at low temperatures. This is a priori a possible interpretation of hardening of the Ni–5 at.% Al alloy below the transition temperature (at 220 ◦ C or 493 K), as pointed out by Ayub et al. [2]. In this way, it is possible to attribute the resistivity peak at 493 K to antiferromagnetic spin ordering in the Ni–4.49 at.% Al alloy. Fig. 2 shows the microhardness changes measurements with quenching temperature and one may observe two steps of hardness decrease, one from 423 to 573 K and other from 573 to 723 K. The first stage seems to confirm what was proposed by Williams [3] and observed in Fig. 1, the anti-ferromagnetic spin ordering. The second stage may be associated to the resistivity peak observed at 670 K in Fig. 1. Fig. 3 shows the transmission electronic micrographs obtained for the Ni–4.49 at.% Al alloy quenched from 670 K (Fig. 3(a)) and from 750 K (Fig. 3(b)), on the resistivity maximum and after it, in the resistivity curve of Fig. 1, together with the corresponding electron diffraction pattern. It is possible to observe a change in the crystalline planes orientation with the heat treatment, probably due to the atomic reordering assisted by vacancy migration. The diffraction pattern in Fig. 3(c) seems to indicate an ordered distribution of the diffracting planes, and in Fig. 3(d) a disordered one, in a distance small enough to give a rounded spots configuration.
Fig. 2. Microhardness changes vs. quenching temperature. Sample initially quenched from 1373 K.
A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
57
Fig. 3. Transmission electron micrographs obtained for samples quenched from: (a) 670 K and (b) 750 K. (c and d) Electron diffraction patterns obtained, respectively, at the same conditions.
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A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
Fig. 5. Isothermal resistivity changes with time.
Fig. 4. DSC curve obtained for the Ni–4.49 at.% Al alloy initially quenched from 1373 K. Heating rate 10 K min−1 .
Fig. 4 shows the DSC curve obtained for the alloy initially quenched from 1373 K. In this figure, it is possible to observe two well-defined thermal events, one endothermic peak at 795 K and an exothermic peak at 825 K. It is also possible to observe two other (very small) endothermic peaks, one at about 500 K and other at about 650 K. The two latter peaks may be associated, respectively, to anti-ferromagnetic spin ordering and to the order restored by newly created equilibrium vacancies, as already mentioned in the discussion of Fig. 1. The endothermic peak at 795 K may be related to the beginning of the resistivity decrease observed in the last part of the curve in Fig. 1 and the exothermic peak at 825 K to the end of this decrease. This resistivity decrease seems to be due to the disordering process occurring at higher temperatures, which begins at 795 K and ends at 825 K. This seems to confirm what was observed by Ayub et al. [2], that the equilibrium short-range order states in Ni–Al alloys of composition ranging from 2 to 8 at.% Al are obtained at 400 ◦ C (673 K). The degree of short-range order is found to be maximum at this temperature and decreases significantly with increase in temperature, and the alloys are essentially in disordered state above 700 ◦ C (973 K). These results seem to indicate that the order–disorder phenomena occurring in the Ni–4.49 at.% Al alloy were detected with different accuracy by the various experimental techniques here employed. All the considered events were detected by resistivity measurements, but their intensity was variable, with the most pronounced one corresponding to the maximum of short-range ordering. The microhardness measurements, although with an accuracy far from the resistivity measurements, were able to indicate a relationship between local order phase changes and the alloy mechanical properties. The DSC curve showed that, among all events, the transition with the greater heat content is that corresponding to the beginning of the disordering process.
Considering that SRO kinetics is a local probe of atomic mobility and that what is measured is the mobility of the faster element (Al, in the present case) [6], it is possible to evaluate a value for the activation energy associated to Al diffusion in the Ni –4.49 at.% Al alloy from the resistivity measurements shown in Fig. 1, and to get some important insights about the Al diffusion mechanism and the Ni–Al interaction. Fig. 5 shows the plot of isothermal resistivity changes with time obtained for samples aged at the temperatures of 668, 688, 701 and 718 K, chosen from the isochronic curve in Fig. 1. These isothermal curves suggest a first-order kinetics model, with the SRO relaxation described by Rt − R∞ t (1) = exp − Ro − R ∞ τT where τ T is the relaxation time at each ageing temperature and its temperature dependence follows the Arrhenius law EM τT = τo exp − (2) kT where EM is the migration activation energy for the order–disorder transformation. Using Eq. (1) and Fig. 5, the values of τ T for each ageing temperature were obtained and are shown in Table 1. Fig. 6 shows the plot of ln τ T versus the inverse of absolute temperature, for the τ T values in Table 1. The activation energy value, obtained from the slope of the straight line in this figure, is EM = (1.69 ± 0.09) eV. This value is close to that = (1.4 ± 0.3) eV for the obtained by Afyouni et al. [6], EM Ni–6 at.% Al alloy and the small difference may be attributed to the Al content. As EM is associated with Al diffusion and Table 1 Relaxation times obtained for each ageing temperature Temperature (K)
Relaxation time (min)
668 688 701 718
352 134 115 40
A.T. Adorno et al. / Journal of Alloys and Compounds 414 (2006) 55–59
59
4. Conclusions The resistivity measurements confirmed the ordering behavior expected for Ni–Al dilute alloys and the suggested relation between resistivity changes and microhardness changes with anti-ferromagnetic spin ordering. The order–disorder phenomena occurring in the Ni–4.49 at.% Al alloy were detected with different accuracy by the various experimental techniques. All the considered events were detected by resistivity measurements, but their intensity was variable, with the most pronounced one corresponding to the maximum of short-range ordering. The higher value obtained for the activation energy of vacancy migration in the Ni–4.49 at.% Al alloy was associated with a decrease in the Al concentration gradient near solute-depleted regions. Fig. 6. Plot of ln τ vs. 1/T.
this effect is enhanced by the increase of Al concentration, a difference in the activation energy may be expected for the Ni–4.49 at.% Al alloy. Considering that in a dilute substitutional alloy the rate of homogenization is controlled by the diffusion of solute atoms to solute-depleted regions and that the rate of solute atoms diffusion is greater than that for the solvent atoms, the faster Al diffusion in Ni–Al dilute alloys may be connected to the fact that solute atoms can attract vacancies to produce a more than random probability of finding a vacancy next to a solute atom. Al atoms can attract more vacancies because Al atoms are larger than Ni atoms and, due to the greater Al–vacancy binding energy, the vacancy will be unable to “escape” from the solute atom. In this way, the Al–vacancy pair can diffuse through the lattice together. Hence, the small difference in the activation energies obtained for the Ni–6 at.% Al and Ni–4.49 at.% Al alloys may be associated with a decrease in the Al concentration gradient near solute-depleted regions.
Acknowledgement The authors thank Termomecˆanica Brasil S/A for the preparation and chemical analysis of the samples.
References [1] M. Sluiter, P.E.A. Turchi, F.J. Oinski, G.M. Stocks, Mater. Sci. Eng. A152 (1992) 1. [2] G. Ayub, F.A. Khwaja, A. Ul Haq, Z. Ahmad, Acta Metall. Mater. 43 (1995) 1457. [3] R.O. Williams, Trans. Am. Inst. Miner. Eng. 215 (1989) 1026. [4] R. Kozubski, D. Kmiec, E. Partyka, M. Danielewski, Intermetallics 11 (2003) 897. [5] C. Cserh´ati, A. Paul, A.A. Kodentsov, M.J.H. van Dal, F.J.J. van Loo, Intermetallics 11 (2003) 291. [6] M. Afyouni, V. Pierron-Bohnes, M.C. Cadeville, Acta Metall. 37 (9) (1989) 2339.