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Metastable phase characterization in ion implanted and rapidly solidified aluminium-based alloys
1170
Materials Science and Engineering, A134 (1991 ) 1170-1174
Metastable phase characterization in ion implanted and rapidly solidified aluminium-based alloys Q. Li, E. Johnson, A. Johansen, L. D. Yu, S. Steenstrup and L. Sarholt-Kristensen Physics' Laboratory, H.C. Orsted Institute, University of Copenhagen, Universitetsparken 5, DK-2100 Copenhagen 0 (Denmark)
A. Pogrebnyakov Department of Electronic Devices Technology, Minsk Radioengineering Institute, P. Brovka Street 6, Minsk 220600 (U.S.S.R.)
Abstract Metastable phase formation in rapidly solidified polycrystalline and ion implantation processed single crystal A1-Cu and A1-Zn alloys is investigated for various annealing conditions. X-ray diffraction, scanning electron microscopy and electrical resistivity measurements have been employed to characterize the rapidly solidified ribbons, while examination of the ion implanted crystals was carried out using in situ Rutherford backscattering/channelling analysis. Extended solubility is most pronounced in the ion implanted alloys where a single a(A1) phase is formed at an implantation temperature of - - 8 0 °C. Conversely, in the rapidly solidified alloys, the a(A1) phase occurs together with the metastable or equilibrium phases. For the concentrated alloys of both systems, produced by the two techniques, the precipitation temperature is nearly the same.
1. Introduction
2. Experimental details
Both ion implantation and rapid solidification have been widely used to create a new generation of metastable alloys with extended solubility, containing crystalline, quasicrystalline and amorphous phases. In general, these metastable structures are formed in non-equilibrium processes. Ion implantation involves a solid-solid transformation with cooling rates - 1012 K s-1, while a liquid-solid reaction with cooling rates from 10 5 to 10 8 K s-1 is characteristic of rapid solidification [1]. Such differences may influence the formation and variability of the metastable structures in these two processes. In this work we attempt to compare the precipitation process in A1-Cu and A I - Z n alloys produced by both techniques. Previous work [2] has reported that the metastable/stable phases precipitate during implantation of Cu + --,AI, and further coarsening or phase separation is seen after annealing. For rapidly solidified alloys, the formation of G.E zones is suppressed in both systems [3, 4], and the precipitation sequence, temperature and time are strongly altered due to the high cooling rates [5, 6].
Single crystals of pure aluminium were implanted in an isotope separator along a (110) direction with 50 keV Cu + or Z n + ions at - - 80 °C to fluences of 1 x 102°, 2 x 102°, 5 ×102o and 1 × 10 21 m -2, respectively. In situ RBS/channeling analysis was performed as a function of fluence at the ambient implantation temperature using a 550 keV He 2+ beam. Subsequent analyses for various annealing treatments (0.5 h in steps of 50 °C) were carried out in situ after cooling to room temperature. Rapidly solidified ribbons of polycrystalline A1-Cu (2.5 and 17.3 at.%) and A I - Z n (10 and 26 at.%) were produced in an argon atmosphere by single roller melt spinning from a boron nitride crucible. Homogeneous ribbons were obtained with a nozzle to wheel distance of 0.2 mm using a peripheral wheel speed of 25 m s-i. A detailed description of processing and microstructural analysis is given elsewhere [6].
0921-5093/91/$3.50
3. Results and discussion Figure 1 shows the development of the concentration depth profiles, determined from © Elsevier Sequoia/Printedin The Netherlands
1171
30
(o) 4~
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1 * 1021 m 2
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25
2 . 1020 m-2
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Fig. 1. Concentration depth profiles as a function of fluence for AI crystals implanted at -80°C. (a) Cu÷~AI. (b) Zn + ~AI. random Rutherford backscattering spectra as a function of fluence, for the Cu + --'AI and Zn + --"A1 samples. In both systems the maximum implant concentration increases almost linearly with fluence. At low fluences the two sets of depth distributions are similar with a peak around 40 nm. However, at a fluence of 5 x 1 0 20 m - 2 the peak position in the copper distribution is shifted to a depth of - 60 nm and a second peak appears around 120 nm. This may be associated with precipitation in the A I - C u system, where electron diffraction on polycrystalline aluminium implanted at room temperature with 4 x 102° m - 2 Cu + ions has verified the existence of 0 and probably also 0" phases in the implanted layer [2]. The lattice location of the implanted atoms can be found by combined RBS and channelling analysis. The results indicate nearly full substitution of the implanted zinc for all fluences, in good agreement with the work of Picraux et aL on a dilute alloy [7]. A further increase in supersaturation may be anticipated for even higher fluences. In contrast, for copper in aluminium the substitutional fraction f, decreases from - 0.8 to
- 0 . 7 with increasing fluence for both peak regions. It is comparable to the data of Gerber et al. [8] who, for dilute alloys implanted at liquid nitrogen temperature, found an f, value of 0.82. These results also agree with the D a r k e n - G u r r y plots for aluminium alloys [9], where A1-Zn lies inside the circle of high solubility with f, from 0.8 to 1.0, while A I - C u is just outside this area with J~ from 0.5-0.8 [10, 11]. From lattice parameter determinations of the a(Al) phase in rapidly solidified ribbons, the maximum extension of solid solubility under the present processing parameters is 5.7 at.% Cu and 22 at.% Zn. However, application of the gun technique with cooling rates up to 10 s K s - ] can produce aluminium-based alloys with 17-18 at.% Cu [5] and 38 at.% Zn [12] in solution. In all the as-quenched A1-Cu and A1-Zn ribbons the supersaturated a(AI) phase co-exists with different metastable and/or equilibrium phases. In eutectic A1-Cu ribbons the precipitates are 0 and 0", in good agreement with the ion implantation results of Thackery et al. [2j. Such agreement does not emerge in the A1-Zn alloys where the metastable R phase coexists with the stable /3(Zn) phase in the as-quenched ribbons (Fig. 2(a)), while only the a(Al) phase is formed during implantation. Only samples implanted to the highest fluence were used for subsequent annealing. Figure 3 shows sets of concentration depth profiles as a function of annealing temperature for both Cu + --'A1 and Zn + ~A1. During heating of the samples from the implantation temperature ( - 80 °C) to room temperature, f, decreases from - 0 . 7 to nearly zero for the Cu + ~ A l alloy, while nearly full substitutionality is retained in the Zn + --"AI alloy. For dilute A1-Cu alloys made by implantation at room temperature, the f, value is only - 10% less than at liquid nitrogen temperature [8, 13]. Gerber et al. [8] and Kloska et al. [14] argue that for dilute alloys with (1) positive heat of solution AHsol, or with (2) negative heat of solution and a size mismatch energy AH~z> 10 kJ mol- ~, a decrease in f,, as seen in RBS analysis, is due to an increased interaction between implanted atoms and mobile vacancies, which will cause a displacement of the atoms away from substitutional sites. The f, values for different alloys will hence decrease for case ( 1 ) when A Hso I or AH~iz increases, and for case (2) when only AH~iz increases.
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2e/degrees Fig. 2. X-ray diffraction spectra from the rapidly solidified alloys. (a) AI-17.3at.%Cu. (b) AI-26at.%Zn.
In the dilute limit, AHso ~ and AHsL~ for the AICu system [8] are - 3 2 kJmo1-1 and 11 kJ mo1-1, respectively, while for A I - Z n alloys AHsoI is 2 kJ mol-~ [15] and AH~iz is calculated to be 0.76 kJ mol- 1 [8, 16]. Therefore the probability of forming solute-vacancy complexes is larger in AI-Cu than in AI-Zn, and the decrease in fs after heating to room temperature will therefore be most pronounced in the A1-Cu alloy. According to Miedema [15], for concentrated alloys both the size mismatch energy and the heat of formation are concentration dependent. This can explain the decrease of ~ to zero in the annealed AI-Cu sample. After annealing to 200 °C and 250 °C the peak in the copper distribution at a depth - 1 2 0 nm gradually disappears, while the concentration in the large peak increases. It seems to coincide with the formation of semi-coherent 0' and incoherent 0 phases in the surface vicinity of polycrystalline samples [2]. The high disorder level in the surface region of the (110) aluminium channelled spectra seen at all temperatures (Fig. 4(a)), may then partly be due to residual strain in the aluminium
40
80
120
Depth
(nm)
160
200
Fig. 3. Concentration depth profiles of the ion implanted alloys as a function of annealing temperature. (a) Cu ÷~AI. (b) Zn+ ~A1. The fluenceis 1 xl02j m -2. matrix, and partly to aluminium atoms in the incoherent precipitates [13]. During annealing of the Zn ÷ --,A1 sample to 100°C and 150°C a narrow surface peak appears in the zinc distribution (Figs. 3(b) and 4(b)), for which the fs value is only - 0 . 3 . This suggests formation of the incoherent fl(Zn) phase in the surface region. It is correlated to a broadening of the zinc profile and a lowering of the peak concentration, i.e. diffusion of zinc both towards the surface and towards the interior of the sample. After annealing at 200 °C the zinc profile is flat, and it is anticipated that as the pressure in the target chamber is lower than the vapour pressure for zinc (5 x 10 -5 Pa at 200 °C [16]), zinc will be lost at the surface due to sublimation. The evolution of electrical resistivity as a function of isochronal annealing temperature for the rapidly solidified ribbons is shown in Fig. 5(a). The reaction temperatures for eutectic A1-Cu an(t the two A1-Zn alloys fall in the same range as for the respective implanted alloys. This suggests that the reaction temperature will be independent of the processing conditions if the extension of
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Fig. 4. Random and (110) aligned Rutherford backscattering spectra from (a) Cu + ~A1, and (b) Z n + ~A1 for a fluence of 1xlO21m 2.
solid solubility is high enough. Figure 2 shows that after annealing of the eutectic AI-Cu and the A1-26at.% Zn ribbons, there is no precipitation of new metastable phases, and the 0" and R phases formed during solidification are quickly replaced by the equilibrium phases. In the lower concentrated alloys, on the other hand, the supersaturated a(Al) phase transforms to the metastable 0" and R phases, respectively [6]. SEM micrographs (Fig. 5(b)) show that the 0' phase in the A1-2.5at.%Cu ribbon has a plateshaped morphology similar to that seen in the Cu + ~ A1 crystal [2]. 4. Conclusions
(1) The extension of solid solubility is most pronounced for ion implantation processed alloys where a single a(A1) phase is formed at the implantation temperature. In rapidly solidified alloys the supersaturated a(Al) phase co-exists with various metastable and/or equilibrium phases.
Fig. 5. (a) Electrical resistivity as a function of temperature of the rapidly solidified ribbons. (b) SEM micrograph showing the 0 and 0' phases in the AI 2.5at.%Cu alloy annealed at 280 °C for 0.5 h.
(2) During ion implantation and subsequent annealing of the single crystals, the precipitates nucleate and grow preferably in the surface region of the implanted layer. On the contrary, in polycrystalline ribbons these phases are located both in the a(A1) matrix containing excess vacancies and in the refined grain boundaries. (3) In all the rapidly solidified ribbons and the Cu implanted A1 crystal the supersaturated solution decomposes in a continuous reaction, where the sequence of precipitation passes through a succession of metastable phase equifibria to the respective stable phases. A discontinuous reaction controls the precipitation process in the zinc implanted aluminium crystal, where the supersaturated a(A1) solid solution separates directly into the equilibrium phases. (4) For eutectic AI-Cu and high concentrated A1-Zn alloys made by rapid solidification, the precipitation temperatures are similar to those of the ion implanted alloys.
1174
Acknowledgments Financial support from The Daloon Foundation and The Danish Natural Science Research Council is greatly acknowledged.
References 1 D. Tumbull, Met. Trans., 12A (1981) 695. 2 P. A. Thackery and R. S. Nelsen, Phil, Mag., 19 (1969) 169. 3 M. J. Starink and P. van Mourik, in ASM Int. Conf. on Advanced Aluminium and Magnesium Alloys, Amsterdam, 1990, in press. 4 S. Agarwal and H. Herman, Scr. Metall., 7(1973) 503. 5 M.G. Scott and J. A. Leake, Acta Metall., 23 (1975) 503. 6 Q. Li, E. Johnson, A. Johansen and L. Sarholt-Kristensen, in ASM Int. Conf. on Advanced Aluminium and Magnesium Alloys, Amsterdam, 1990, in press.
7 S. T. Picraux, E. Rimini, G. Foti and S. U. Campisano, Phys. Rev. B, 18 (1978) 2078. 8 R. Gerber, O. Meyer and G. C. Xiong, Nucl. lnstrum. Methods, B31 (1988)402. 9 S. T. Pieraux, in J. E Ziegler (ed.), New Uses orlon Accelerators, Plenum, New York, 1975, p. 229. 10 C. Nordling and J. Osterman, Physics Handbook, Studentlitteratur, Lund, 1987, p. 64. 11 M. Hansen, Constitution of Binary Alloys, McGraw-Hill, New York, 1958, p. 1265. 12 I. V. Salli and L. P. Limina, in D. E. Ovsienko (ed.), Growth and Imperfections in Metallic Crystals, Consu|tants Bureau, New York, 1968, p. 251. 13 D. K. Sood and G. Dearnaley, in G. Carter, J. S. Colligon and W. A. Grant (eds.), Application of Ion Beams to Materials, Institute of Physics, London, 1976, p. 196. 14 M. K. Kloska and O. Meyer, Nucl. Instrum. Methods, B19/20 (1987) 140. 15 A. R. Miedema, P. E de Chfitel and F. R. de Boer, Physica, IOOB (1980) 1. 16 C. J. Smithells, Metals Reference Book, Butterworths, London, 1967, pp. 263,708.
Materials Science and Engineering, A134 (1991 ) 1170-1174
Metastable phase characterization in ion implanted and rapidly solidified aluminium-based alloys Q. Li, E. Johnson, A. Johansen, L. D. Yu, S. Steenstrup and L. Sarholt-Kristensen Physics' Laboratory, H.C. Orsted Institute, University of Copenhagen, Universitetsparken 5, DK-2100 Copenhagen 0 (Denmark)
A. Pogrebnyakov Department of Electronic Devices Technology, Minsk Radioengineering Institute, P. Brovka Street 6, Minsk 220600 (U.S.S.R.)
Abstract Metastable phase formation in rapidly solidified polycrystalline and ion implantation processed single crystal A1-Cu and A1-Zn alloys is investigated for various annealing conditions. X-ray diffraction, scanning electron microscopy and electrical resistivity measurements have been employed to characterize the rapidly solidified ribbons, while examination of the ion implanted crystals was carried out using in situ Rutherford backscattering/channelling analysis. Extended solubility is most pronounced in the ion implanted alloys where a single a(A1) phase is formed at an implantation temperature of - - 8 0 °C. Conversely, in the rapidly solidified alloys, the a(A1) phase occurs together with the metastable or equilibrium phases. For the concentrated alloys of both systems, produced by the two techniques, the precipitation temperature is nearly the same.
1. Introduction
2. Experimental details
Both ion implantation and rapid solidification have been widely used to create a new generation of metastable alloys with extended solubility, containing crystalline, quasicrystalline and amorphous phases. In general, these metastable structures are formed in non-equilibrium processes. Ion implantation involves a solid-solid transformation with cooling rates - 1012 K s-1, while a liquid-solid reaction with cooling rates from 10 5 to 10 8 K s-1 is characteristic of rapid solidification [1]. Such differences may influence the formation and variability of the metastable structures in these two processes. In this work we attempt to compare the precipitation process in A1-Cu and A I - Z n alloys produced by both techniques. Previous work [2] has reported that the metastable/stable phases precipitate during implantation of Cu + --,AI, and further coarsening or phase separation is seen after annealing. For rapidly solidified alloys, the formation of G.E zones is suppressed in both systems [3, 4], and the precipitation sequence, temperature and time are strongly altered due to the high cooling rates [5, 6].
Single crystals of pure aluminium were implanted in an isotope separator along a (110) direction with 50 keV Cu + or Z n + ions at - - 80 °C to fluences of 1 x 102°, 2 x 102°, 5 ×102o and 1 × 10 21 m -2, respectively. In situ RBS/channeling analysis was performed as a function of fluence at the ambient implantation temperature using a 550 keV He 2+ beam. Subsequent analyses for various annealing treatments (0.5 h in steps of 50 °C) were carried out in situ after cooling to room temperature. Rapidly solidified ribbons of polycrystalline A1-Cu (2.5 and 17.3 at.%) and A I - Z n (10 and 26 at.%) were produced in an argon atmosphere by single roller melt spinning from a boron nitride crucible. Homogeneous ribbons were obtained with a nozzle to wheel distance of 0.2 mm using a peripheral wheel speed of 25 m s-i. A detailed description of processing and microstructural analysis is given elsewhere [6].
0921-5093/91/$3.50
3. Results and discussion Figure 1 shows the development of the concentration depth profiles, determined from © Elsevier Sequoia/Printedin The Netherlands
1171
30
(o) 4~
_ _
1 * 1021 m 2
......
5 . 1 0 2 0 m -2
25
2 . 1020 m-2
v
1020 m-2
20 o
4J
15
ro c_
4J
10
E 8J
u
/
5
o (D
o 40
80 Depth
o~ 4J
120
160
200
(nm)
3O
1. lo~ ~
(b) 25
--
3~ ~/u20 '~ m2
_. . . .
2 * 1020 m-2 /• , /u20m ~2
113
2O o
4J
15
f0 [_
4J
10
/'~-'~"-.
E O (_1
1-
o
0
40
so
Depth
120
160
2oo
(nm)
Fig. 1. Concentration depth profiles as a function of fluence for AI crystals implanted at -80°C. (a) Cu÷~AI. (b) Zn + ~AI. random Rutherford backscattering spectra as a function of fluence, for the Cu + --'AI and Zn + --"A1 samples. In both systems the maximum implant concentration increases almost linearly with fluence. At low fluences the two sets of depth distributions are similar with a peak around 40 nm. However, at a fluence of 5 x 1 0 20 m - 2 the peak position in the copper distribution is shifted to a depth of - 60 nm and a second peak appears around 120 nm. This may be associated with precipitation in the A I - C u system, where electron diffraction on polycrystalline aluminium implanted at room temperature with 4 x 102° m - 2 Cu + ions has verified the existence of 0 and probably also 0" phases in the implanted layer [2]. The lattice location of the implanted atoms can be found by combined RBS and channelling analysis. The results indicate nearly full substitution of the implanted zinc for all fluences, in good agreement with the work of Picraux et aL on a dilute alloy [7]. A further increase in supersaturation may be anticipated for even higher fluences. In contrast, for copper in aluminium the substitutional fraction f, decreases from - 0.8 to
- 0 . 7 with increasing fluence for both peak regions. It is comparable to the data of Gerber et al. [8] who, for dilute alloys implanted at liquid nitrogen temperature, found an f, value of 0.82. These results also agree with the D a r k e n - G u r r y plots for aluminium alloys [9], where A1-Zn lies inside the circle of high solubility with f, from 0.8 to 1.0, while A I - C u is just outside this area with J~ from 0.5-0.8 [10, 11]. From lattice parameter determinations of the a(Al) phase in rapidly solidified ribbons, the maximum extension of solid solubility under the present processing parameters is 5.7 at.% Cu and 22 at.% Zn. However, application of the gun technique with cooling rates up to 10 s K s - ] can produce aluminium-based alloys with 17-18 at.% Cu [5] and 38 at.% Zn [12] in solution. In all the as-quenched A1-Cu and A1-Zn ribbons the supersaturated a(AI) phase co-exists with different metastable and/or equilibrium phases. In eutectic A1-Cu ribbons the precipitates are 0 and 0", in good agreement with the ion implantation results of Thackery et al. [2j. Such agreement does not emerge in the A1-Zn alloys where the metastable R phase coexists with the stable /3(Zn) phase in the as-quenched ribbons (Fig. 2(a)), while only the a(Al) phase is formed during implantation. Only samples implanted to the highest fluence were used for subsequent annealing. Figure 3 shows sets of concentration depth profiles as a function of annealing temperature for both Cu + --'A1 and Zn + ~A1. During heating of the samples from the implantation temperature ( - 80 °C) to room temperature, f, decreases from - 0 . 7 to nearly zero for the Cu + ~ A l alloy, while nearly full substitutionality is retained in the Zn + --"AI alloy. For dilute A1-Cu alloys made by implantation at room temperature, the f, value is only - 10% less than at liquid nitrogen temperature [8, 13]. Gerber et al. [8] and Kloska et al. [14] argue that for dilute alloys with (1) positive heat of solution AHsol, or with (2) negative heat of solution and a size mismatch energy AH~z> 10 kJ mol- ~, a decrease in f,, as seen in RBS analysis, is due to an increased interaction between implanted atoms and mobile vacancies, which will cause a displacement of the atoms away from substitutional sites. The f, values for different alloys will hence decrease for case ( 1 ) when A Hso I or AH~iz increases, and for case (2) when only AH~iz increases.
1172 6KL
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(el) fu
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250°6 200%
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4J
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(_ 4~
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I 35
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C o
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30
- -
35
1oo% SO°t RT \
\\
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45
40
80 Depth
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120
160
200
(nm)
30
(b) .p
- 50°6
25
...... I00OC .... 150% - - _ _ _ 200oC
ro cO
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~ , .
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annealing16(P£ Ofi~r
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50
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2e/degrees Fig. 2. X-ray diffraction spectra from the rapidly solidified alloys. (a) AI-17.3at.%Cu. (b) AI-26at.%Zn.
In the dilute limit, AHso ~ and AHsL~ for the AICu system [8] are - 3 2 kJmo1-1 and 11 kJ mo1-1, respectively, while for A I - Z n alloys AHsoI is 2 kJ mol-~ [15] and AH~iz is calculated to be 0.76 kJ mol- 1 [8, 16]. Therefore the probability of forming solute-vacancy complexes is larger in AI-Cu than in AI-Zn, and the decrease in fs after heating to room temperature will therefore be most pronounced in the A1-Cu alloy. According to Miedema [15], for concentrated alloys both the size mismatch energy and the heat of formation are concentration dependent. This can explain the decrease of ~ to zero in the annealed AI-Cu sample. After annealing to 200 °C and 250 °C the peak in the copper distribution at a depth - 1 2 0 nm gradually disappears, while the concentration in the large peak increases. It seems to coincide with the formation of semi-coherent 0' and incoherent 0 phases in the surface vicinity of polycrystalline samples [2]. The high disorder level in the surface region of the (110) aluminium channelled spectra seen at all temperatures (Fig. 4(a)), may then partly be due to residual strain in the aluminium
40
80
120
Depth
(nm)
160
200
Fig. 3. Concentration depth profiles of the ion implanted alloys as a function of annealing temperature. (a) Cu ÷~AI. (b) Zn+ ~A1. The fluenceis 1 xl02j m -2. matrix, and partly to aluminium atoms in the incoherent precipitates [13]. During annealing of the Zn ÷ --,A1 sample to 100°C and 150°C a narrow surface peak appears in the zinc distribution (Figs. 3(b) and 4(b)), for which the fs value is only - 0 . 3 . This suggests formation of the incoherent fl(Zn) phase in the surface region. It is correlated to a broadening of the zinc profile and a lowering of the peak concentration, i.e. diffusion of zinc both towards the surface and towards the interior of the sample. After annealing at 200 °C the zinc profile is flat, and it is anticipated that as the pressure in the target chamber is lower than the vapour pressure for zinc (5 x 10 -5 Pa at 200 °C [16]), zinc will be lost at the surface due to sublimation. The evolution of electrical resistivity as a function of isochronal annealing temperature for the rapidly solidified ribbons is shown in Fig. 5(a). The reaction temperatures for eutectic A1-Cu an(t the two A1-Zn alloys fall in the same range as for the respective implanted alloys. This suggests that the reaction temperature will be independent of the processing conditions if the extension of
1173 9OO 80O
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2 0 - (Q)
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o 100
150
200
250
300
350
400
450
/ 5-
500
800
--
± 100
500 Temperature (°C)
Channels
900
o
(b)
o R~dom RT b Random '50% c <110> 150%
700
ddlO~
RT
400
450
600 500 400 300 200 100 0 100
150
200
250
300
350
500
Channels
Fig. 4. Random and (110) aligned Rutherford backscattering spectra from (a) Cu + ~A1, and (b) Z n + ~A1 for a fluence of 1xlO21m 2.
solid solubility is high enough. Figure 2 shows that after annealing of the eutectic AI-Cu and the A1-26at.% Zn ribbons, there is no precipitation of new metastable phases, and the 0" and R phases formed during solidification are quickly replaced by the equilibrium phases. In the lower concentrated alloys, on the other hand, the supersaturated a(Al) phase transforms to the metastable 0" and R phases, respectively [6]. SEM micrographs (Fig. 5(b)) show that the 0' phase in the A1-2.5at.%Cu ribbon has a plateshaped morphology similar to that seen in the Cu + ~ A1 crystal [2]. 4. Conclusions
(1) The extension of solid solubility is most pronounced for ion implantation processed alloys where a single a(A1) phase is formed at the implantation temperature. In rapidly solidified alloys the supersaturated a(Al) phase co-exists with various metastable and/or equilibrium phases.
Fig. 5. (a) Electrical resistivity as a function of temperature of the rapidly solidified ribbons. (b) SEM micrograph showing the 0 and 0' phases in the AI 2.5at.%Cu alloy annealed at 280 °C for 0.5 h.
(2) During ion implantation and subsequent annealing of the single crystals, the precipitates nucleate and grow preferably in the surface region of the implanted layer. On the contrary, in polycrystalline ribbons these phases are located both in the a(A1) matrix containing excess vacancies and in the refined grain boundaries. (3) In all the rapidly solidified ribbons and the Cu implanted A1 crystal the supersaturated solution decomposes in a continuous reaction, where the sequence of precipitation passes through a succession of metastable phase equifibria to the respective stable phases. A discontinuous reaction controls the precipitation process in the zinc implanted aluminium crystal, where the supersaturated a(A1) solid solution separates directly into the equilibrium phases. (4) For eutectic AI-Cu and high concentrated A1-Zn alloys made by rapid solidification, the precipitation temperatures are similar to those of the ion implanted alloys.
1174
Acknowledgments Financial support from The Daloon Foundation and The Danish Natural Science Research Council is greatly acknowledged.
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