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Phase transformations and strengthening during ageing of CuNi10Al3 alloy
Materials Science and Engineering A264 (1999) 279 – 285
Phase transformations and strengthening during ageing of CuNi10Al3 alloy Zdzisl*aw Sierpin´ski *, Janusz Gryziecki Department of Non-Ferrous Alloys, Uni6ersity of Mining and Metallurgy, Mickiewicza A6. 30, PL-30 -059 Krako´w, Poland Received 3 March 1998; received in revised form 29 May 1998
Abstract Microstructure and mechanical properties of precipitation hardened Cu – 10.5 wt.% Ni – 3.15 wt.% Al (CuNi10Al3) after thermal and thermomechanical treatments were investigated. The highest strengthening in this alloy (yield strength larger than 1GPa) has been obtained after thermomechanical treatment (supersaturation, 50% cold reduction, ageing at 773 K). Some alternative procedures to obtain high strength are also proposed. The microstructure shows discontinuous precipitates at grain boundaries. Transmission electron microscopy and X-ray diffraction helped to establish that the main strengthening phase in this alloy is g% (Ni3Al with ordered L12 structure) which precipitates discontinuously and continuously. The b phase (NiAl-B2 structure) also forms during each of the described treatments. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cu – Ni – Al alloys; Precipitation strengthening; g% (Ni3Al) phase
1. Introduction The effect of high precipitation hardening caused by particles of ordered g% (L12 type) phase is frequently used to improve the properties of alloys. The heat resistant alloys of nickel and iron are good examples of the effect [1,2]. The same idea was also tried in some copper alloys [3– 6]. Owing due to their notable corrosion resistance cupronickels have found widespread application in the shipping and chemical industries [7]. However, copper and nickel form only solid solutions. Therefore, the effect of precipitation hardening might be achieved by adding a third component, for example aluminium [8,9]. According to the Cu – Ni – Al ternary phase diagram [9] the precipitation process can occur when the copper based alloy contain 5 – 6 wt.% of Ni and at least 3 wt.% of Al (Fig. 1). For such compositions the b phase (NiAl) is the equilibrium phase at low temperatures. Because of morphology and the structure of the b phase, precipitation strengthening in the copper base * Corresponding author. E-mail address: [email protected] (Z. Sierpin´ski)
alloy with low aluminium and nickel contents is low and not durable [4,10–12]. Cu–Ni–Al alloys with a nickel content higher than 7 wt.% and at least 3 wt.% Al aged at about 770 K show predominant precipitation of g% (Ni3Al) and accordingly allow an improvement in strengthening to be obtained [8,9]. Earlier reports on Cu–Ni–Al alloys [3 –6,10–13] confirmed that the addition of aluminium to cupronickels causes many concurrent precipitation processes like: continuous precipitation of g% (Ni3Al) phase, discontinuous precipitation of rod-like form of the same phase at grain boundaries, continuous precipitation of b (NiAl) phase. Following mechanical testing aged Cu–Ni–Al alloys were characterized by very high strength and a good plasticity. The precipitation effect can be obtained in a wide range of ageing temperatures. Moreover the best strengthening effect in CuNi10Al3 can be achieved when the temperature of heat treatment is about 770 K. At higher temperatures of ageing the effect is smaller. Thermomechanical treatment gives the best result when about 50% plastic deformation is imposed on the material [11]. Heat treated Cu–Ni–Al alloys retain the majority of the physical properties of cupronickels [13]. The present paper reports observations on the correlation between phase transformations and strengthen-
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Fig. 1. Cu – Ni – Al phase diagram at 773 K [9].
ing of the CuNi10Al3 alloy. The aim of this investigation is to design a heat treatment procedure with which it is possibile to affect the microstructure and mechanical properties of this alloy. The Ni and Al additions were kept at a relatively low level as the alloy already capable of high strengthening, is still easy to cast and cold work.
The microstructure observations were performed with Neophot optical and Philips CM20 (200 kV) transmission electron microscopes. The structure of the phases were determined with electron and X-ray diffraction techniques. The mechanical property changes were analyzed by Vickers hardness test and tensile test, using flat speci-
2. Experimental procedure The alloy with composition of 10.5 wt.% Ni and 3.15 wt.% Al, balance copper was prepared from 99.99% purity components, cast in a vacuum and homogenized. After solution treatment at 1173 K it was: 1. water quenched, then isothermally aged (treatment A); 2. water quenched, rolled to 50% reduction, then isothermally aged (treatment B); 3. slowly cooled to 773 K, then isothermally aged (treatment C). The mean grain size of the supersaturated alloy was about 80 mm. During each treatment the ageing temperature was kept at 77393 K.
Fig. 2. Hardness changes of CuNi10Al3 alloy during treatments A, B and C. Ageing temperature 773 K.
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Fig. 3. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment A, as a function of ageing time.
Fig. 5. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment C, as a function of ageing time.
mens. The mean error of tensile parameters was estimated to be about 5%. Detailed descriptions of the sample preparation and the experimental method were given elsewhere [13].
faster increase of this parameter during ageing. The maximum strengthening, manifested by yield strength s0.2 exceeding 1GPa, is attained after ageing at 773 K for 5× 103 s. Such a high strength is accompanied by elongation o=4%. Longer ageing results in a drop of both s0.2 and sm and a slow increase of o (Fig. 4). Comparing the changes of mechanical properties during ageing of supersaturated and deformed materials, a much higher relation s0.2/sm in the second case has been observed. The high strength properties, after a short time, appear during treatment C. After a few minutes of ageing the yield strength exceeds 600 MPa, and after 6×105 s it achieves a level of 850 MPa (Fig. 5). The microstructure observations with the use of an optical microscope performed after treatment A showed that the grain boundaries are fringed by a zone of material of different contrast. That type of microstructure suggests the occurrence of a discontinuous type of decomposition. During treatment A these zones of distinct contrast reach 0.1 mm (Fig. 6a). After treatment C these zones were much smaller (Fig. 6c). In the alloy structure subjected to treatment B, a high density of deformation bands is visible (Fig. 6b). The supposition of discontinuous precipitation process at grain boundaries has been confirmed by transmission electron microscopy. During treatment A both large roughly spherical precipitates of diameter up to 300 nm (Fig. 7a) and colonies of rod-like discontinuous precipitates form at grain boundaries (Fig. 7b). The remainder of the grains are filled with evenly distributed much smaller spherical precipitates (Fig. 7a,b). It was found that the highest strengthening is attained when the size of these precipitates is around 20 nm. The selected area diffraction allowed to prove that both rod-like and fine spherical precipitates have the same type of ordered lattice of the L12 type and structure parameters similar to copper (Fig. 7c). The plastic deformation introduces high density of dislocations which tend to form dislocation bands arranged in a regular network (Fig. 8a). During treatment
3. Results Fig. 2 shows the changes of Vickers hardness during ageing of CuNi10Al3 alloy subjected to treatments A, B and C. The water quenched and aged alloy (treatment A) shows slow increase of hardness. After 105 s it reaches 250 HV. A much faster rise in hardness and even to a higher level of 290 HV is obtained for the alloy which was cold rolled before ageing (treatment B). The alloy slowly cooled and aged (treatment C) achieves a maximum hardness of about 230 HV after the shortest ageing time. The significant differences in strengthening during ageing were clearly visible in the tensile characteristics, as summarized in Figs. 3–5. During treatment A an increase in tensile strength sm from 380 MPa in supersaturated state to 720 MPa after ageing for 105 s and yield strength s0.2 from 120 to 550 MPa respectively, corresponds to a decrease in elongation o from 28 to 19% (Fig. 3). The mechanical properties of the alloy deformed 50% and aged (treatment B) show both a higher starting strength (650 MPa) and a
Fig. 4. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment B, as a function of ageing time.
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B the nucleation and growth of spherical precipitates (diameter up to 200 nm after 1.8 × 103 s of ageing) within the shear bands takes place (Fig. 8b). Additionally shear bands serve as a nucleation site for in situ recrystallization (Fig. 8c). On the other hand, similarly like in treatment A, the interiors of grains contain numerous fine spherical precipitates. The overaged alloy (7.8×105 s of ageing) contain a significantly increased volume of recrystallized areas, which explains the lowering of strength at that stage of ageing. Inside these areas the renucleation of spherical, fine precipitates takes place. Recrystallization front does not dissolve the larger b precipitates (Fig. 8d). The microstructure of alloy treated by C mode is similar to the type of microstructure after treatment A.
The main difference lies in the fact that small spherical precipitates were also visible inside the colonies of discontinuous precipitates (Fig. 9a,b,c). Larger b precipitates were found to be situated at the discontinuous precipitation front rather than at the original grain boundary position (Fig. 9b). The dark field image taken with the superlattice reflection (001) confirms that rodlike and fine spherical precipitates are the same phase (Fig. 9c). The uniformly distributed precipitates attained a diameter of about 20 nm after 6×102 s of ageing (Fig. 9). The size and shape of these precipitates remain stable during further ageing. The X-ray diffraction confirmed that ageing of CuNi10Al3 alloys results in precipitation of ordered Ni3Al (g%) and NiAl (b) phases (Fig. 10).
Fig. 6. Optical microscopy images of CuNi10Al3 alloy after: (a) 2.6 ×104 s of treatment A; (b) 5.4 ×104 s of treatment B and; (c) 1.8 × 103 s of treatment C. Magnification × 125.
Fig. 7. The microstructure of CuNi10Al3 alloy after 1.2× 106 s of treatment A (a,b), and accompanying electron diffraction (c). DP cell indicates the area occupied by discontinuous precipitation.
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Fig. 8. The microstructure of CuNi10Al3 alloy after treatment B (a,b) 1.8× 103 s; (c) 3×104 s and; (d) 7.8× 105 s RA and URA indicates recrystallized and unrecrystallized area, respectively.
Fig. 9. The microstructure of CuNi10Al3 alloy after 2.1 × 104 s of treatment C: (a,b) bright field image; (c) dark field image taken from (100) superlattice spot. DP cell indicates area occupied by discontinuous precipitation.
4. Discussion The results of the mechanical tests and microstructure observations performed at different stages of ageing proved that the mechanical property changes are closely correlated with the type and the extent of phase transformation taking place in the material.
During treatment A, based on ageing of the supersaturated alloy, three typical diffusion processes take place: first, the heterogeneous nucleation and growth of b phase and discontinuous form of g% phase at grain boundaries, and then the nucleation and growth of continuous form of g% phase. The continuous precipitation of g% is the slowest process and it controls the
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Fig. 10. The results of X-ray diffraction analysis of CuNi10Al3 alloy after 900 s and 1.2 ×106 s of treatment A: (a) (100) Ni3Al superlatice peak, (b) (110) NiAl peak.
strength of the alloy. Other precipitation processes also take part in the strengthening but to a lesser extent. Unfortunately, because of the slow growth rate of continuous g% precipitates treatment A might not be considered as a suitable treatment for industrial application. The material subjected to treatment B, including plastic deformation prior to ageing, is characterized by five times higher initial strength than that in the supersaturated state. The deformed material, except the precipitation processes, is undergoing simultaneous recovery as well as recrystallization. The intensity of the three heat activated parallel processes changes during the several stages of ageing. It is worth noting that there is no discontinuous g% precipitates at shear bands. It suggests that heterogeneous nucleation and growth of b precipitates in these conditions are much faster. It means that a large deformation is the way to generate a new nucleation site for the b phase. Unfortunately, b precipitates can not retard recrystallization front. In spite of recovery and recrystallization, the alloy achieves significant strength properties (s0.2 larger than 1 GPa). The renucleation of g% phase indicates that the recrystallized areas have the possibility to decompose. The low amount of secondary precipitates and the long time needed for their formation is due to the fact that
they have no significant influence on mechanical properties. The simplest and most economical procedure C— like A—is also based exclusively on heat treatment. However, the ageing temperature at this treatment is reached through slow cooling of the material. Therefore, the precipitation processes have a chance to start at temperatures higher than the ageing temperature (773 K). The resulting microstructure is characterized by coarse precipitates. The existence of fine continuous precipitates within discontinuous colonies indicate that the discontinuous precipitation has been a primary process during this treatment started at temperature close to that of sol6us temperature. The newly formed zones of the a phase inside discontinuous precipitation colonies contain enough alloying additions to serve later on, i.e. at much lower temperatures as the matrix for continuous precipitation of the g% phase. The above conclusion about the discontinuous precipitation taking place ahead of the others is also confirmed by observation of b phase precipitates formed only at the front of growing discontinuous colonies. In spite of the high temperatures at which the precipitation processes start during this treatment, no precipitation free zones near grain boundaries—like in aluminium precipitation strengthened alloys—were found. As the precipitation processes start already during cooling the alloy reaches its maximum strength after relatively short time. Moreover, treatment C leads to the limitation of discontinuous precipitation, which contributes in a smaller extent to the total strength than continuous one. The X-ray and electron diffraction confirm that both the g% and b phases form during ageing, independent of treatment conditions.
5. Conclusions 1. The main strengthening process in the CuNi10Al3 alloy, independent of applied heat treatment conditions, is continuous precipitation of ordered g%–Ni3Al phase. The b (NiAl) and discontinuous, rod-like form of g% (Ni3Al) phase also contribute to strengthening but to a smaller extent. 2. The long time needed to obtain a high strengthening, as a result of A treatment (supersaturation, quenching into water, isothermal ageing) may be shortened by the application of treatment C (supersaturation, direct quenching to the ageing temperature, isothermal ageing). 3. During ageing of supersaturated and deformed alloy the intensive nucleation and growth of the b phase in shear bands takes place. The recrystallization front dissolves small, continuous g% precipitates and does not dissolve the larger b precipitates. The repre-
Z. Sierpin´ski, J. Gryziecki / Materials Science and Engineering A264 (1999) 279–285
cipitation of g% phase in recrystallized areas occurs after a long ageing time. 4. Changing the treatment conditions it is possible to control the size, morphological form and proportion of contents of both phases in the alloy and consequently the mechanical properties of the CuNi10Al3 alloy, obtaining even higher strength material. Having in mind well known good physical and chemical properties of Cu–Ni alloys, it confirms good prospects for Cu–Ni– Al future applications.
Acknowledgements This work received financial support from the Polish State Committee of Scientific Research (KBN) under grant 7T08B 001 10, and from the Foundation for Polish Science. We are grateful to W. Baliga for experimental support and to Dr J. Morgiel for several helpful discussions.
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References [1] K. Kusabiraki, X. Zhang, T. Ooka, ISIJ Int. 35 (1995) 1115. [2] K. Kusabiraki, E. Amada, T. Ooka, ISIJ Int. 36 (1996) 208. [3] W. Leo, G. Wassermann, Metallurgy 21 (1967) 10. [4] Y.-R. Cho, Y.-H. Kim, T.D. Lee, J. Mater. Sci. 26 (1991) 2879. [5] M. Miki, Y. Amano, Trans. J. Intst. Metals 20 (1979) 1. [6] R.N. Caron, S. Shapiro, US Patent number US4016010. [7] G.W. Blunn, E.B.G. Jones, Incra Project, 332A, Hayling Island, UK, 1984, p. 67. [8] W.O. Alexander, J. Inst. Metals LXIII (1938) 163. [9] H. Baker (ed.), ASM Handbook, vol. 3, Alloy Phase Diagrams, ASM, 1992. * atkowski, Arch. Metall. 40 [10] J. Gryziecki, Z. Sierpin´ski, A. L (1995) 437. * atkowski, Arch. Metall. 41 [11] J. Gryziecki, Z. Sierpin´ski, A. L (1996) 78. [12] Z. Sierpin´ski, Rep. Polish Acad. Sci., Secesja, Krako´w XL/2 (1996) 88. [13] Z. Sierpin´ski, International Conference of PhD Students, Miscolc, August 11 – 17, University of Miscolc, Hungary, 1997, p. 168.
Phase transformations and strengthening during ageing of CuNi10Al3 alloy Zdzisl*aw Sierpin´ski *, Janusz Gryziecki Department of Non-Ferrous Alloys, Uni6ersity of Mining and Metallurgy, Mickiewicza A6. 30, PL-30 -059 Krako´w, Poland Received 3 March 1998; received in revised form 29 May 1998
Abstract Microstructure and mechanical properties of precipitation hardened Cu – 10.5 wt.% Ni – 3.15 wt.% Al (CuNi10Al3) after thermal and thermomechanical treatments were investigated. The highest strengthening in this alloy (yield strength larger than 1GPa) has been obtained after thermomechanical treatment (supersaturation, 50% cold reduction, ageing at 773 K). Some alternative procedures to obtain high strength are also proposed. The microstructure shows discontinuous precipitates at grain boundaries. Transmission electron microscopy and X-ray diffraction helped to establish that the main strengthening phase in this alloy is g% (Ni3Al with ordered L12 structure) which precipitates discontinuously and continuously. The b phase (NiAl-B2 structure) also forms during each of the described treatments. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cu – Ni – Al alloys; Precipitation strengthening; g% (Ni3Al) phase
1. Introduction The effect of high precipitation hardening caused by particles of ordered g% (L12 type) phase is frequently used to improve the properties of alloys. The heat resistant alloys of nickel and iron are good examples of the effect [1,2]. The same idea was also tried in some copper alloys [3– 6]. Owing due to their notable corrosion resistance cupronickels have found widespread application in the shipping and chemical industries [7]. However, copper and nickel form only solid solutions. Therefore, the effect of precipitation hardening might be achieved by adding a third component, for example aluminium [8,9]. According to the Cu – Ni – Al ternary phase diagram [9] the precipitation process can occur when the copper based alloy contain 5 – 6 wt.% of Ni and at least 3 wt.% of Al (Fig. 1). For such compositions the b phase (NiAl) is the equilibrium phase at low temperatures. Because of morphology and the structure of the b phase, precipitation strengthening in the copper base * Corresponding author. E-mail address: [email protected] (Z. Sierpin´ski)
alloy with low aluminium and nickel contents is low and not durable [4,10–12]. Cu–Ni–Al alloys with a nickel content higher than 7 wt.% and at least 3 wt.% Al aged at about 770 K show predominant precipitation of g% (Ni3Al) and accordingly allow an improvement in strengthening to be obtained [8,9]. Earlier reports on Cu–Ni–Al alloys [3 –6,10–13] confirmed that the addition of aluminium to cupronickels causes many concurrent precipitation processes like: continuous precipitation of g% (Ni3Al) phase, discontinuous precipitation of rod-like form of the same phase at grain boundaries, continuous precipitation of b (NiAl) phase. Following mechanical testing aged Cu–Ni–Al alloys were characterized by very high strength and a good plasticity. The precipitation effect can be obtained in a wide range of ageing temperatures. Moreover the best strengthening effect in CuNi10Al3 can be achieved when the temperature of heat treatment is about 770 K. At higher temperatures of ageing the effect is smaller. Thermomechanical treatment gives the best result when about 50% plastic deformation is imposed on the material [11]. Heat treated Cu–Ni–Al alloys retain the majority of the physical properties of cupronickels [13]. The present paper reports observations on the correlation between phase transformations and strengthen-
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 0 8 3 - 1
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Fig. 1. Cu – Ni – Al phase diagram at 773 K [9].
ing of the CuNi10Al3 alloy. The aim of this investigation is to design a heat treatment procedure with which it is possibile to affect the microstructure and mechanical properties of this alloy. The Ni and Al additions were kept at a relatively low level as the alloy already capable of high strengthening, is still easy to cast and cold work.
The microstructure observations were performed with Neophot optical and Philips CM20 (200 kV) transmission electron microscopes. The structure of the phases were determined with electron and X-ray diffraction techniques. The mechanical property changes were analyzed by Vickers hardness test and tensile test, using flat speci-
2. Experimental procedure The alloy with composition of 10.5 wt.% Ni and 3.15 wt.% Al, balance copper was prepared from 99.99% purity components, cast in a vacuum and homogenized. After solution treatment at 1173 K it was: 1. water quenched, then isothermally aged (treatment A); 2. water quenched, rolled to 50% reduction, then isothermally aged (treatment B); 3. slowly cooled to 773 K, then isothermally aged (treatment C). The mean grain size of the supersaturated alloy was about 80 mm. During each treatment the ageing temperature was kept at 77393 K.
Fig. 2. Hardness changes of CuNi10Al3 alloy during treatments A, B and C. Ageing temperature 773 K.
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Fig. 3. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment A, as a function of ageing time.
Fig. 5. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment C, as a function of ageing time.
mens. The mean error of tensile parameters was estimated to be about 5%. Detailed descriptions of the sample preparation and the experimental method were given elsewhere [13].
faster increase of this parameter during ageing. The maximum strengthening, manifested by yield strength s0.2 exceeding 1GPa, is attained after ageing at 773 K for 5× 103 s. Such a high strength is accompanied by elongation o=4%. Longer ageing results in a drop of both s0.2 and sm and a slow increase of o (Fig. 4). Comparing the changes of mechanical properties during ageing of supersaturated and deformed materials, a much higher relation s0.2/sm in the second case has been observed. The high strength properties, after a short time, appear during treatment C. After a few minutes of ageing the yield strength exceeds 600 MPa, and after 6×105 s it achieves a level of 850 MPa (Fig. 5). The microstructure observations with the use of an optical microscope performed after treatment A showed that the grain boundaries are fringed by a zone of material of different contrast. That type of microstructure suggests the occurrence of a discontinuous type of decomposition. During treatment A these zones of distinct contrast reach 0.1 mm (Fig. 6a). After treatment C these zones were much smaller (Fig. 6c). In the alloy structure subjected to treatment B, a high density of deformation bands is visible (Fig. 6b). The supposition of discontinuous precipitation process at grain boundaries has been confirmed by transmission electron microscopy. During treatment A both large roughly spherical precipitates of diameter up to 300 nm (Fig. 7a) and colonies of rod-like discontinuous precipitates form at grain boundaries (Fig. 7b). The remainder of the grains are filled with evenly distributed much smaller spherical precipitates (Fig. 7a,b). It was found that the highest strengthening is attained when the size of these precipitates is around 20 nm. The selected area diffraction allowed to prove that both rod-like and fine spherical precipitates have the same type of ordered lattice of the L12 type and structure parameters similar to copper (Fig. 7c). The plastic deformation introduces high density of dislocations which tend to form dislocation bands arranged in a regular network (Fig. 8a). During treatment
3. Results Fig. 2 shows the changes of Vickers hardness during ageing of CuNi10Al3 alloy subjected to treatments A, B and C. The water quenched and aged alloy (treatment A) shows slow increase of hardness. After 105 s it reaches 250 HV. A much faster rise in hardness and even to a higher level of 290 HV is obtained for the alloy which was cold rolled before ageing (treatment B). The alloy slowly cooled and aged (treatment C) achieves a maximum hardness of about 230 HV after the shortest ageing time. The significant differences in strengthening during ageing were clearly visible in the tensile characteristics, as summarized in Figs. 3–5. During treatment A an increase in tensile strength sm from 380 MPa in supersaturated state to 720 MPa after ageing for 105 s and yield strength s0.2 from 120 to 550 MPa respectively, corresponds to a decrease in elongation o from 28 to 19% (Fig. 3). The mechanical properties of the alloy deformed 50% and aged (treatment B) show both a higher starting strength (650 MPa) and a
Fig. 4. Changes of yield strength s0.2, tensile strength and elongation of CuNi10Al3 alloy during treatment B, as a function of ageing time.
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B the nucleation and growth of spherical precipitates (diameter up to 200 nm after 1.8 × 103 s of ageing) within the shear bands takes place (Fig. 8b). Additionally shear bands serve as a nucleation site for in situ recrystallization (Fig. 8c). On the other hand, similarly like in treatment A, the interiors of grains contain numerous fine spherical precipitates. The overaged alloy (7.8×105 s of ageing) contain a significantly increased volume of recrystallized areas, which explains the lowering of strength at that stage of ageing. Inside these areas the renucleation of spherical, fine precipitates takes place. Recrystallization front does not dissolve the larger b precipitates (Fig. 8d). The microstructure of alloy treated by C mode is similar to the type of microstructure after treatment A.
The main difference lies in the fact that small spherical precipitates were also visible inside the colonies of discontinuous precipitates (Fig. 9a,b,c). Larger b precipitates were found to be situated at the discontinuous precipitation front rather than at the original grain boundary position (Fig. 9b). The dark field image taken with the superlattice reflection (001) confirms that rodlike and fine spherical precipitates are the same phase (Fig. 9c). The uniformly distributed precipitates attained a diameter of about 20 nm after 6×102 s of ageing (Fig. 9). The size and shape of these precipitates remain stable during further ageing. The X-ray diffraction confirmed that ageing of CuNi10Al3 alloys results in precipitation of ordered Ni3Al (g%) and NiAl (b) phases (Fig. 10).
Fig. 6. Optical microscopy images of CuNi10Al3 alloy after: (a) 2.6 ×104 s of treatment A; (b) 5.4 ×104 s of treatment B and; (c) 1.8 × 103 s of treatment C. Magnification × 125.
Fig. 7. The microstructure of CuNi10Al3 alloy after 1.2× 106 s of treatment A (a,b), and accompanying electron diffraction (c). DP cell indicates the area occupied by discontinuous precipitation.
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Fig. 8. The microstructure of CuNi10Al3 alloy after treatment B (a,b) 1.8× 103 s; (c) 3×104 s and; (d) 7.8× 105 s RA and URA indicates recrystallized and unrecrystallized area, respectively.
Fig. 9. The microstructure of CuNi10Al3 alloy after 2.1 × 104 s of treatment C: (a,b) bright field image; (c) dark field image taken from (100) superlattice spot. DP cell indicates area occupied by discontinuous precipitation.
4. Discussion The results of the mechanical tests and microstructure observations performed at different stages of ageing proved that the mechanical property changes are closely correlated with the type and the extent of phase transformation taking place in the material.
During treatment A, based on ageing of the supersaturated alloy, three typical diffusion processes take place: first, the heterogeneous nucleation and growth of b phase and discontinuous form of g% phase at grain boundaries, and then the nucleation and growth of continuous form of g% phase. The continuous precipitation of g% is the slowest process and it controls the
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Fig. 10. The results of X-ray diffraction analysis of CuNi10Al3 alloy after 900 s and 1.2 ×106 s of treatment A: (a) (100) Ni3Al superlatice peak, (b) (110) NiAl peak.
strength of the alloy. Other precipitation processes also take part in the strengthening but to a lesser extent. Unfortunately, because of the slow growth rate of continuous g% precipitates treatment A might not be considered as a suitable treatment for industrial application. The material subjected to treatment B, including plastic deformation prior to ageing, is characterized by five times higher initial strength than that in the supersaturated state. The deformed material, except the precipitation processes, is undergoing simultaneous recovery as well as recrystallization. The intensity of the three heat activated parallel processes changes during the several stages of ageing. It is worth noting that there is no discontinuous g% precipitates at shear bands. It suggests that heterogeneous nucleation and growth of b precipitates in these conditions are much faster. It means that a large deformation is the way to generate a new nucleation site for the b phase. Unfortunately, b precipitates can not retard recrystallization front. In spite of recovery and recrystallization, the alloy achieves significant strength properties (s0.2 larger than 1 GPa). The renucleation of g% phase indicates that the recrystallized areas have the possibility to decompose. The low amount of secondary precipitates and the long time needed for their formation is due to the fact that
they have no significant influence on mechanical properties. The simplest and most economical procedure C— like A—is also based exclusively on heat treatment. However, the ageing temperature at this treatment is reached through slow cooling of the material. Therefore, the precipitation processes have a chance to start at temperatures higher than the ageing temperature (773 K). The resulting microstructure is characterized by coarse precipitates. The existence of fine continuous precipitates within discontinuous colonies indicate that the discontinuous precipitation has been a primary process during this treatment started at temperature close to that of sol6us temperature. The newly formed zones of the a phase inside discontinuous precipitation colonies contain enough alloying additions to serve later on, i.e. at much lower temperatures as the matrix for continuous precipitation of the g% phase. The above conclusion about the discontinuous precipitation taking place ahead of the others is also confirmed by observation of b phase precipitates formed only at the front of growing discontinuous colonies. In spite of the high temperatures at which the precipitation processes start during this treatment, no precipitation free zones near grain boundaries—like in aluminium precipitation strengthened alloys—were found. As the precipitation processes start already during cooling the alloy reaches its maximum strength after relatively short time. Moreover, treatment C leads to the limitation of discontinuous precipitation, which contributes in a smaller extent to the total strength than continuous one. The X-ray and electron diffraction confirm that both the g% and b phases form during ageing, independent of treatment conditions.
5. Conclusions 1. The main strengthening process in the CuNi10Al3 alloy, independent of applied heat treatment conditions, is continuous precipitation of ordered g%–Ni3Al phase. The b (NiAl) and discontinuous, rod-like form of g% (Ni3Al) phase also contribute to strengthening but to a smaller extent. 2. The long time needed to obtain a high strengthening, as a result of A treatment (supersaturation, quenching into water, isothermal ageing) may be shortened by the application of treatment C (supersaturation, direct quenching to the ageing temperature, isothermal ageing). 3. During ageing of supersaturated and deformed alloy the intensive nucleation and growth of the b phase in shear bands takes place. The recrystallization front dissolves small, continuous g% precipitates and does not dissolve the larger b precipitates. The repre-
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cipitation of g% phase in recrystallized areas occurs after a long ageing time. 4. Changing the treatment conditions it is possible to control the size, morphological form and proportion of contents of both phases in the alloy and consequently the mechanical properties of the CuNi10Al3 alloy, obtaining even higher strength material. Having in mind well known good physical and chemical properties of Cu–Ni alloys, it confirms good prospects for Cu–Ni– Al future applications.
Acknowledgements This work received financial support from the Polish State Committee of Scientific Research (KBN) under grant 7T08B 001 10, and from the Foundation for Polish Science. We are grateful to W. Baliga for experimental support and to Dr J. Morgiel for several helpful discussions.
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