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Cu–Nb alloys prepared by mechanical alloying and subsequent heat treatment
Journal of Alloys and Compounds 365 (2004) 157–163
Cu–Nb alloys prepared by mechanical alloying and subsequent heat treatment E. Botcharova∗ , J. Freudenberger, L. Schultz IFW Dresden, Institut für Metallische Werkstoffe, PO Box 270016, D-01171 Dresden, Germany Received 14 April 2003; received in revised form 5 June 2003; accepted 5 June 2003
Abstract Nanostructured Cu–Nb composites are prepared by mechanical alloying and a subsequent heat treatment. Up to 10 at.% Nb can be brought into solid solution by milling for 30 h at liquid nitrogen temperature. Subsequent to mechanical alloying a heat treatment was applied to precipitate the niobium from the solid solution in order to enhance both strength and conductivity of the material. The atmosphere present during the milling process strongly influences the phase formation. A high oxygen content within the powder leads to the formation of niobium oxide during the late state of milling and, later, prevents the formation of niobium precipitates by heat treatment. This occurs only in powder with a low oxygen content. A temperature of 400 ◦ C proves to be the optimum for strength, but too low to enhance the conductivity by complete precipitation of the Nb atoms. Almost all the niobium precipitates from the solid solution at 900 ◦ C, but recrystallisation takes place at this temperature. Thus, the optimum heat treatment produces a balance between the mechanical and electrical properties and should be in the range of 600–700 ◦ C. © 2003 Elsevier B.V. All rights reserved. Keywords: Transition metal alloys; Mechanical alloying; Nanostructured materials
1. Introduction A common method of preparing metallic materials and especially, in this case, conductors with high mechanical strength, is to harden the material by insoluble particles. This method becomes efficient only if the distribution of the particles is homogeneous, which is difficult in most cases. A homogeneous distribution either of insoluble particles or alloying atoms can then be used efficiently for precipitation or dispersion hardening. Whenever conductivity is important for the application in mind, copper is the most common metal used, as it shows high conductivity and is, at the same time, cost-effective. Thus copper is suitable as the matrix element in strengthened conductor materials. There are many investigations proving that copper can be successfully dispersion strengthened by oxide or carbide particles. On the other hand the increase of the ultimate tensile strength (UTS)
∗
Corresponding author. Fax: +49-351-465-9541. E-mail address: [email protected] (E. Botcharova).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0925-8388(03)00634-0
at room temperature from ∼250 [1] to ∼600 MPa [2–4] is relatively low. This even holds for the absolute maximum value. Copper-based materials with a much higher UTS can be obtained by means of precipitation hardening, which can be highly effective in metallic materials and, in the case of copper-based materials, one can obtain up to 1000 MPa [1]. In order to apply this hardening mechanism to an alloy, the added elements should show a strongly temperature dependent solubility. If once brought into solution, an element can be precipitated from the matrix by a heat treatment at temperatures where the solubility is much lower than the content of the element in the alloy. The conditions for these heat treatments are chosen with respect to the corresponding phase diagram. The formation of a copper-based solid solution with a thermodynamically insoluble element, as for example Nb, cannot be achieved just by a suitable heat treatment, which can result in a largely enhanced solid solubility [5]. Otherwise it can be achieved through high energy milling or mechanical alloying [6–8]. In a former investigation we showed that ∼10 at.% niobium can be dissolved in a copper matrix by means of mechanical alloying [9]. The
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present study demonstrates the effect of heat treatment (applied subsequent to the mechanically alloying of the powder) on its microstructure and mechanical properties. Furthermore, this study reports on the influence on the alloying process itself of the atmosphere present during the milling and on the phase formation which takes place in the subsequently applied heat treatment.
2. Experimental details Cu–Nb alloys were prepared by mechanical alloying and a subsequent heat treatment. The milling was performed in a PM 4000 Retsch planetary ball mill, using a rotational speed of 200 rpm at a constant rotation direction and a powder-to-ball weight ratio of 1:14. For the purpose of mechanical alloying high purity powders with a particle size of less than 35 m were weighed to give the appropriate composition. This powder mixture was then used for mechanical alloying at low temperatures under either an argon or an air atmosphere. The vial containing balls and powder was cooled in a bath of liquid nitrogen to 77 K and afterwards put onto the planetary mill. In order to ensure that the alloying process takes place almost at liquid nitrogen temperature, the alloying process was interrupted after every 30 min to cool the vial in liquid nitrogen again. After different milling times, a small amount of the powder was taken from the alloy for microstructural investigations. The alloying is regarded to be completed after 30 h of milling. From this alloy, specimens were taken for heat treatments at constant temperatures from 200 up to 900 ◦ C each. The heat treatment was performed under an argon atmosphere for 1 h. The powder was then embedded in conductive resin and treated by standard metallographic techniques for microstructural analysis. These specimens were investigated by scanning electron microscopy (SEM) in a JEOL JSM 6400 device and, for high-resolution SEM, in a LEO 1530. To investigate the lattice structure and parameters (at room temperature) as well as the phase purity of
the samples, a Philips PW 1830 X-ray diffractometer in Bragg-Brentano geometry was used. The X-ray diffraction experiments were performed on crushed powders using Co K␣ (λ = 0.178896 nm) radiation. The scans were taken from 2θ = 20◦ up to 125◦ in steps of θ = 0.05◦ . The lattice parameters were determined from the measured diffraction data using the DBWS Rietveld program. The grain size and the internal strain were obtained from the measured data using Williamson-Hall plots [10]. The impurity content of the powder after milling was determined by chemical analysis. The microhardness of the powder was measured on embedded powder using a Vickers microhardness tester with a load of 245 mN at a loading time of 10 s.
3. Results 3.1. Mechanical alloying Mechanical alloying of two ductile elements, in this case copper and niobium, first forms a coarse lamellar structure, as shown in the SEM micrograph of Fig. 1a after 6 h of milling. Upon further milling the lamellar structure becomes finer and finally reaches a state where no lamellae or small areas of pure copper and/or niobium can be detected any more, even by means of high resolution SEM. This state has been reached after 30 h of milling. The corresponding SEM micrograph is shown in Fig. 1b. The EDX analysis (not shown here) of powder milled for 30 h reveals a very homogeneous distribution of Nb within the particles. Depending on the milling atmosphere, different oxygen contents were found in the powders. It has been found that the oxygen has an important impact on the alloying process, as will be discussed later. Whereas the oxygen content in the powder increases with milling time, it is less for the alloy milled under argon compared to that milled in air (Fig. 2). The powder milled under argon reveals an oxygen content of ∼1.1 ma.% after 30 h of milling.
Fig. 1. SEM micrographs of the microstructure of Cu–Nb alloys after 6 h (a) and 30 h (b) of milling, respectively (milling was performed under argon).
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Fig. 4. Lattice parameter of the copper matrix versus milling time. Fig. 2. Oxygen content of the powders versus milling time for milling under different atmospheres.
Powder milled in air ends up with 1.8 ma.% after milling for the same time. Concerning the microstructure as well as its development during the alloying process, no difference can be detected between the powders milled under argon or in air. X-ray powder diffraction patterns of mechanically alloyed Cu–Nb alloys are shown in Fig. 3 for milling times of 6, 14 and 30 h, respectively. With increasing milling time the diffraction reflections corresponding to copper become broader. The corresponding FWHM (full width half maximum) values of the 220 reflection at 2θ≈89◦ are 1.5◦ , 2.9◦ and 4.2◦ after 6, 14 and 30 h, respectively. At the same time, when the copper reflections become broader, the niobium reflections lose intensity. After 30 h of milling the Nb reflections disappear into the background signal. Thus, a long range order within the niobium particles cannot be detected by the applied X-ray powder diffraction technique. Therefore, the alloys are regarded as a solid solution of niobium in copper after 30 h of milling. From the line width of the re-
Fig. 3. X-ray powder diffraction patterns at different stages of milling (milling was performed under argon).
flections corresponding to the copper, the grain size and the internal strain were calculated by applying Williamson-Hall plots [10]. In the case of a linear dependence of β cos θ (β-integral line width) versus sin θ, the slope of the straight line corresponds to the contribution of the internal strain and the intersection with the ordinate yields the contribution of the line broadening by the grain size [9]. The copper grains become progressively smaller during milling and end up with a size of ∼8 nm after 30 h of milling. Parallel to this, the internal strain grows with increasing milling time, corresponding to a dislocation density of 1.5·1012 cm−2 after 30 h of milling. The lattice parameter of the copper matrix, which was calculated using the Rietveld method [11], is displayed as the function of the milling time in Fig. 4. During milling it increases until it reaches its maximum after 18 h. The maximum value for powder milled under argon is 0.3636(4) nm, and is higher than the matrix lattice parameter of powder milled in air, which is only 0.3634(2) nm (a0 (Cu) = 0.36151 nm). The subsequent decrease of the lattice parameter of the copper matrix after 18 h of milling is steeper for powder milled in air than for powder milled under argon. After 30 h of milling the lattice parameter of the copper matrix reaches 0.3633(5) nm for powder milled under argon, and 0.3628(5) nm for powder milled in air. The microhardness of the powder increases very rapidly within the first 10 h of milling (Fig. 5) until it saturates after ∼15 h. In the saturated state the powder milled under argon reveals a microhardness of ∼520 HV0.025 , whereas that milled in air exhibits values of ∼470 HV0.025 . 3.1.1. Heat treatment SEM pictures of different powders treated at selected temperatures are shown in Figs. 6 and 7. Niobium precipitates from the solid solution starting at ∼600 ◦ C, if the oxygen content is low (i.e. powder milled under argon). A heat treatment at 600 ◦ C for 1 h yields a size of the Nb precipitates of 40–80 nm (Fig. 6a). At higher temperatures, coarsening of the precipitates sets in leading to a grain size of 100–150
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Fig. 5. Microhardness of Cu–Nb alloy powder versus milling time.
nm at 700 ◦ C and 150–350 nm at 900 ◦ C (Fig. 6b). No precipitation of niobium can be detected in powders with a high oxygen content (i.e. powder milled in air ) heat treated at temperatures up to 900 ◦ C, as shown in Fig. 7.
After each of these particular heat treatments, a sample was taken for X-ray powder diffraction analysis in order to trace the phase formation in accordance with the heat treatment. Up to temperatures of 500 ◦ C no significant change of the powder diffraction pattern of the solid solution can be detected (Fig. 8). Higher temperatures cause the formation of Cu2 O as well as NbO. The intensity of the reflections related to these oxides is higher in powder with a higher oxygen content, even after a high temperature heat treatment. At heat treatment temperatures above 600 ◦ C, different scenarios occur, depending on the oxygen content of the powder. In powder with a low oxygen content (i.e. powder milled under argon) and heat treated at 600 ◦ C, niobium reflections can be detected. In contrast to this, reflections related to niobium can only be detected in powder with a high oxygen content when it is heat treated at 900 ◦ C. Due to the heat treatment, other microstructural properties change, such as the lattice parameter of the copper matrix, the dislocation density, the corresponding grain size and finally the microhardness. These properties are shown together in Fig. 9.
Fig. 6. SEM micrographs of Cu–Nb alloys after heat treatment for 1 h at 600 ◦ C (a) and 900 ◦ C (b) (the corresponding samples were milled under argon).
Fig. 7. SEM micrographs of Cu–Nb alloys after heat treatment for 1 h at 600 ◦ C (a) and 900 ◦ C (b) (the corresponding samples were milled in air).
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Fig. 8. X-ray powder diffraction patterns after heat treatment at different temperatures for 1 h (the corresponding samples were milled under argon (a) and in air (b)).
With increasing heat treatment temperature the lattice parameter of the copper matrix decreases until the value converges to that of pure copper. The lattice parameter of the copper matrix of the oxygen-poor sample is quite high. This is caused by the fact that niobium is dissolved into the copper matrix rather than being involved with the formation of phase impurities. This corresponds to a low fraction of phase impurities. Furthermore, either a longer heat treatment or a higher temperature is required to precipitate the niobium from the solid solution, as more niobium is still dissolved. The grain size of the copper matrix of heat treated powder increases with increasing heat treatment temperature. After ageing at 900 ◦ C for 1 h, the copper grain size is ∼30 nm. The dislocation density, which is calculated from the internal strain of the copper lattice,
decreases with increasing heat treatment temperature up to 600 ◦ C at which it reaches an equilibrium state where it stays constant at a level of ∼1·1010 cm−2 . The microhardness of the powder does not change due to heat treatments at temperatures below 300 ◦ C. With increasing heat treatment temperature, a small increase of hardness is observed. The maximum hardness is observed in the case of a powder milled under argon after applying a heat treatment at 400 ◦ C and in the case of a powder milled in air after a heat treatment at 500 ◦ C. A further increase of the heat treatment temperature leads to a decrease of the microhardness of the powder, roughly independent of the oxygen content of the powder. After a heat treatment at 900 ◦ C the microhardness of the powder is ∼400 HV0.025 , in both cases.
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Fig. 9. Lattice parameter of the copper matrix, dislocation density in the copper matrix, grain size of the copper matrix and microhardness of the powder (from top to bottom) as a function of the heat treatment temperature for a 1 h heat treatment.
4. Discussion Copper and niobium are mechanically alloyed to a supersaturated solid solution by means of high energy ball milling. Starting from elementary powders the process of mechanical alloying itself takes ∼30 h. Although the Cu–Nb system does not show a negative heat of mixing, a solid solution of copper and niobium is formed due to the dislocation pumping mechanism [12], which becomes effective due to the fine microstructure of the powder, which is formed during milling [9,13]. The milling atmosphere influences this process. For example, the lattice parameter of the copper matrix varies differently depending on whether milling takes place under argon or in air. In the powder milled under argon the lattice parameter of the copper matrix reaches a higher value than in the case of powder milled in air. According to the
dislocation pumping mechanism the niobium atoms situated at the Cu/Nb phase boundary diffuse along dislocation lines into the copper grains. The driving force for this mechanism is the interaction between the dislocation and the dissolved atom. The effect of the atmosphere on the lattice parameter is caused by oxygen. The oxygen atoms more likely diffuse along the dislocation lines than within the grains. Thus both oxygen and niobium are situated along the dislocation lines where they are enriched. Subsequently, niobium oxide particles are formed instead of the copper niobium solid solution. The more oxygen is present in the powder the more niobium is oxidised and thus less niobium is brought into solid solution. In the powder milled in air the formation of niobium oxide prevails over the dissolution of niobium atoms in copper. This leads to the decrease of the lattice parameter of the copper matrix after 15–18 h of milling (Fig. 4). Furthermore, there are no niobium precipitates found in powder with a high oxygen content after heat treatments up to 900 ◦ C, providing further evidence for the effect of atmosphere on the phase forming process during milling. Within this powder, only niobium oxide can be detected either by SEM or X-ray diffraction. On the other side, there are niobium precipitates present in the heat treated oxygen-poor powder, although there is still some niobium oxide found in these samples, but less than in the others. This niobium oxide is already formed during the late stages of milling, characterised by the almost constant level of the lattice parameter of the copper matrix after 18 h of milling. The oxygen content is not high enough to oxidise the niobium completely and accordingly, niobium precipitates can be detected after heat treatments, even at relatively low temperatures. The decrease of the lattice parameter of the copper matrix with increasing heat treatment temperature indicates an ongoing precipitation of niobium from the supersaturated solid solution in the temperature range between 200 and 600 ◦ C. Above this temperature the lattice parameter of the copper matrix does not change any more and is equal to that of pure copper. Analysing the powder with a high oxygen content heat treated at 900 ◦ C, only small niobium reflections are detectable by means of X-ray analysis, whereas no niobium precipitates can be found by means of the SEM investigations, probably because of the very small precipitate size. This leads to the assumption that niobium, which precipitated from the solid solution, reacts with oxygen atoms and forms niobium oxide. The small amount of niobium precipitates, whose reflections are detected by X-ray analysis, originates from regions where the level of niobium and oxygen atoms becomes unbalanced and finally favors the precipitation of niobium instead of the formation of niobium oxide from the entire niobium content. The development of the microhardness of the heat treated powder is controlled by three processes: (i) precipitation of niobium, (ii) coarsing of the precipitates (in the case of powder with a low oxygen content) or of the niobium oxide particles (in the case of powder with a high oxygen content), and (iii) recrystallisation of the copper matrix and growth of
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the copper grains. These processes, as shown in Fig. 9, take place during the heat treatment and influence the properties of the Cu–Nb alloy. In particular, the formation of niobium precipitates leads to an increase of the microhardness and, conversely, the recrystallisation process, grain growth and precipitation coarsening lead to its decrease. At temperatures below 300 ◦ C, the softening effect of the recrystallisation process compensates for the strengthening effect caused by the formation of Nb or NbO precipitates. In the temperature range between 300 and 400 ◦ C, Nb precipitates are formed. NbO is formed in the case of a high oxygen content of the powders in the range between 300 and 500 ◦ C. At higher temperatures (up to 600 ◦ C) the microhardness decreases as a consequence of the recrystallisation processes. At temperatures above 600 ◦ C the microhardness value is controlled by the precipitation coarsening and grain growth of the copper matrix, which lead to a decrease of the microhardness to 400 HV0.025 after a heat treatment at 900 ◦ C. With respect to the development of the microhardness, an intermediate heat treatment temperature is supposed to generate niobium precipitates with an optimum size, resulting in the best hardening properties. Unfortunately, after a heat treatment at 400 ◦ C, the size of the precipitates cannot be determined as they are too small to be detected by SEM analysis. Furthermore, the lattice parameter of the copper matrix is still higher than that of pure copper. This indicates that a part of the niobium atoms still remains within the solid solution after a heat treatment at 400 ◦ C. This results in an appreciably low electrical conductivity. The optimum heat treatment produces a balance between the mechanical and electrical properties and should be in the range of 600–700 ◦ C.
5. Conclusions The supersaturated solid solution of niobium in copper is formed by mechanical alloying of elementary copper and niobium powder after 30 h of milling. The milling atmosphere influences this process: in the powder milled in air (1.8 ma.% oxygen content after 30 h of milling) niobium oxide particles are formed during the late stage of milling. Thus, almost all the niobium atoms are bound to oxygen atoms. No niobium precipitates can be formed in such powder during the subsequent heat treatments at temperatures up
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to 900 ◦ C, only niobium oxide. In contrast to this, niobium precipitates in powder with a low oxygen content (milled under argon; 1.2 ma.% oxygen content after 30 h of milling) during heat treatment. The decrease of the microhardness of the powder during the heat treatment is negligible. The microhardness value of powder with a low oxygen content treated at 700 ◦ C is ∼470 HV0.025 corresponding to a UTS of ∼1.5 GPa. A significant increase of the electrical properties of heat treated powder is expected after the complete precipitation of niobium from the solid solution. This investigation has shown that the mechanical alloying route with a subsequent heat treatment is a promising way of producing a conductor material with a high mechanical strength, as required for pulsed high-field magnets [14].
Acknowledgements This work was supported by the Bundesministerium für Bildung und Forschung by grant No. 03SC5DRE.
References [1] W. Heller, in: R.W. Cahn, P. Haasen, E.J. Krumer (Eds.), Materials Science and Technology, vol. 8; in: K.H. Matucha (Ed.), Structure and Properties of Nonferrous Alloys, VCH, Weinheim, 1996, p. 277. [2] R.K. Islamgaliev, W. Buchgraber, Y.R. Kolobov, N.M. Amirkhanov, A.V. Sergueeva, K.V. Ivanov, G.P. Grabovetskaya, Mater. Sci. Eng. A 319–321 (2001) 872. [3] S. Ochiai, in: Materials Engineering 7. Mechanical Properties of Metallic Composites, Marcel Dekker, New York, 1994. [4] D. Kudashov, M. Heilmaier, U. Martin, H. Oettel, in: Proceedings of the European Powder Metallurgy Congress and Exhibition, Euro PM 2001, Nizza, 22–24 October, 2001, Vol. 4, p. 872. [5] L. Schultz, R. Bormann, J. Appl. Phys. 50 (1979) 418. [6] A. Benghalem, D.G. Morris, Scripta Metall. Mater. 27 (1992) 739. [7] A. Benghalem, D.G. Morris, Mater. Sci. Eng. A 161 (1993) 255. [8] L. Schultz, J. Eckert, in: H. Beck, H.J. Güntherodt (Eds.), Glassy Metals III, Mechanically Alloyed Glassy Metals, Springer, Berlin, 1994, p. 69. [9] E. Botcharova, M. Heilmaier, J. Freudenberger, G. Drew, D. Kudashow, U. Martin, L. Schultz, J. Alloys Comp. 351 (2003) 119. [10] G.K. Williamson, W.H. Hall, Acta Metall. 1 (1953) 22. [11] H.M. Rietveld, J. Appl. Crystallogr. 2 (1969) 65. [12] R.B. Schwarz, Mater. Sci. Forum 269–272 (1998) 665. [13] C. Suryanaryana, Progr. Mater. Sci. 46 (2001) 1. [14] J. Freudenberger, W. Grünberger, E. Botcharova, A. Gaganov, L. Schultz, Adv. Eng. Mater. 4 (2002) 677.
Cu–Nb alloys prepared by mechanical alloying and subsequent heat treatment E. Botcharova∗ , J. Freudenberger, L. Schultz IFW Dresden, Institut für Metallische Werkstoffe, PO Box 270016, D-01171 Dresden, Germany Received 14 April 2003; received in revised form 5 June 2003; accepted 5 June 2003
Abstract Nanostructured Cu–Nb composites are prepared by mechanical alloying and a subsequent heat treatment. Up to 10 at.% Nb can be brought into solid solution by milling for 30 h at liquid nitrogen temperature. Subsequent to mechanical alloying a heat treatment was applied to precipitate the niobium from the solid solution in order to enhance both strength and conductivity of the material. The atmosphere present during the milling process strongly influences the phase formation. A high oxygen content within the powder leads to the formation of niobium oxide during the late state of milling and, later, prevents the formation of niobium precipitates by heat treatment. This occurs only in powder with a low oxygen content. A temperature of 400 ◦ C proves to be the optimum for strength, but too low to enhance the conductivity by complete precipitation of the Nb atoms. Almost all the niobium precipitates from the solid solution at 900 ◦ C, but recrystallisation takes place at this temperature. Thus, the optimum heat treatment produces a balance between the mechanical and electrical properties and should be in the range of 600–700 ◦ C. © 2003 Elsevier B.V. All rights reserved. Keywords: Transition metal alloys; Mechanical alloying; Nanostructured materials
1. Introduction A common method of preparing metallic materials and especially, in this case, conductors with high mechanical strength, is to harden the material by insoluble particles. This method becomes efficient only if the distribution of the particles is homogeneous, which is difficult in most cases. A homogeneous distribution either of insoluble particles or alloying atoms can then be used efficiently for precipitation or dispersion hardening. Whenever conductivity is important for the application in mind, copper is the most common metal used, as it shows high conductivity and is, at the same time, cost-effective. Thus copper is suitable as the matrix element in strengthened conductor materials. There are many investigations proving that copper can be successfully dispersion strengthened by oxide or carbide particles. On the other hand the increase of the ultimate tensile strength (UTS)
∗
Corresponding author. Fax: +49-351-465-9541. E-mail address: [email protected] (E. Botcharova).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0925-8388(03)00634-0
at room temperature from ∼250 [1] to ∼600 MPa [2–4] is relatively low. This even holds for the absolute maximum value. Copper-based materials with a much higher UTS can be obtained by means of precipitation hardening, which can be highly effective in metallic materials and, in the case of copper-based materials, one can obtain up to 1000 MPa [1]. In order to apply this hardening mechanism to an alloy, the added elements should show a strongly temperature dependent solubility. If once brought into solution, an element can be precipitated from the matrix by a heat treatment at temperatures where the solubility is much lower than the content of the element in the alloy. The conditions for these heat treatments are chosen with respect to the corresponding phase diagram. The formation of a copper-based solid solution with a thermodynamically insoluble element, as for example Nb, cannot be achieved just by a suitable heat treatment, which can result in a largely enhanced solid solubility [5]. Otherwise it can be achieved through high energy milling or mechanical alloying [6–8]. In a former investigation we showed that ∼10 at.% niobium can be dissolved in a copper matrix by means of mechanical alloying [9]. The
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E. Botcharova et al. / Journal of Alloys and Compounds 365 (2004) 157–163
present study demonstrates the effect of heat treatment (applied subsequent to the mechanically alloying of the powder) on its microstructure and mechanical properties. Furthermore, this study reports on the influence on the alloying process itself of the atmosphere present during the milling and on the phase formation which takes place in the subsequently applied heat treatment.
2. Experimental details Cu–Nb alloys were prepared by mechanical alloying and a subsequent heat treatment. The milling was performed in a PM 4000 Retsch planetary ball mill, using a rotational speed of 200 rpm at a constant rotation direction and a powder-to-ball weight ratio of 1:14. For the purpose of mechanical alloying high purity powders with a particle size of less than 35 m were weighed to give the appropriate composition. This powder mixture was then used for mechanical alloying at low temperatures under either an argon or an air atmosphere. The vial containing balls and powder was cooled in a bath of liquid nitrogen to 77 K and afterwards put onto the planetary mill. In order to ensure that the alloying process takes place almost at liquid nitrogen temperature, the alloying process was interrupted after every 30 min to cool the vial in liquid nitrogen again. After different milling times, a small amount of the powder was taken from the alloy for microstructural investigations. The alloying is regarded to be completed after 30 h of milling. From this alloy, specimens were taken for heat treatments at constant temperatures from 200 up to 900 ◦ C each. The heat treatment was performed under an argon atmosphere for 1 h. The powder was then embedded in conductive resin and treated by standard metallographic techniques for microstructural analysis. These specimens were investigated by scanning electron microscopy (SEM) in a JEOL JSM 6400 device and, for high-resolution SEM, in a LEO 1530. To investigate the lattice structure and parameters (at room temperature) as well as the phase purity of
the samples, a Philips PW 1830 X-ray diffractometer in Bragg-Brentano geometry was used. The X-ray diffraction experiments were performed on crushed powders using Co K␣ (λ = 0.178896 nm) radiation. The scans were taken from 2θ = 20◦ up to 125◦ in steps of θ = 0.05◦ . The lattice parameters were determined from the measured diffraction data using the DBWS Rietveld program. The grain size and the internal strain were obtained from the measured data using Williamson-Hall plots [10]. The impurity content of the powder after milling was determined by chemical analysis. The microhardness of the powder was measured on embedded powder using a Vickers microhardness tester with a load of 245 mN at a loading time of 10 s.
3. Results 3.1. Mechanical alloying Mechanical alloying of two ductile elements, in this case copper and niobium, first forms a coarse lamellar structure, as shown in the SEM micrograph of Fig. 1a after 6 h of milling. Upon further milling the lamellar structure becomes finer and finally reaches a state where no lamellae or small areas of pure copper and/or niobium can be detected any more, even by means of high resolution SEM. This state has been reached after 30 h of milling. The corresponding SEM micrograph is shown in Fig. 1b. The EDX analysis (not shown here) of powder milled for 30 h reveals a very homogeneous distribution of Nb within the particles. Depending on the milling atmosphere, different oxygen contents were found in the powders. It has been found that the oxygen has an important impact on the alloying process, as will be discussed later. Whereas the oxygen content in the powder increases with milling time, it is less for the alloy milled under argon compared to that milled in air (Fig. 2). The powder milled under argon reveals an oxygen content of ∼1.1 ma.% after 30 h of milling.
Fig. 1. SEM micrographs of the microstructure of Cu–Nb alloys after 6 h (a) and 30 h (b) of milling, respectively (milling was performed under argon).
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Fig. 4. Lattice parameter of the copper matrix versus milling time. Fig. 2. Oxygen content of the powders versus milling time for milling under different atmospheres.
Powder milled in air ends up with 1.8 ma.% after milling for the same time. Concerning the microstructure as well as its development during the alloying process, no difference can be detected between the powders milled under argon or in air. X-ray powder diffraction patterns of mechanically alloyed Cu–Nb alloys are shown in Fig. 3 for milling times of 6, 14 and 30 h, respectively. With increasing milling time the diffraction reflections corresponding to copper become broader. The corresponding FWHM (full width half maximum) values of the 220 reflection at 2θ≈89◦ are 1.5◦ , 2.9◦ and 4.2◦ after 6, 14 and 30 h, respectively. At the same time, when the copper reflections become broader, the niobium reflections lose intensity. After 30 h of milling the Nb reflections disappear into the background signal. Thus, a long range order within the niobium particles cannot be detected by the applied X-ray powder diffraction technique. Therefore, the alloys are regarded as a solid solution of niobium in copper after 30 h of milling. From the line width of the re-
Fig. 3. X-ray powder diffraction patterns at different stages of milling (milling was performed under argon).
flections corresponding to the copper, the grain size and the internal strain were calculated by applying Williamson-Hall plots [10]. In the case of a linear dependence of β cos θ (β-integral line width) versus sin θ, the slope of the straight line corresponds to the contribution of the internal strain and the intersection with the ordinate yields the contribution of the line broadening by the grain size [9]. The copper grains become progressively smaller during milling and end up with a size of ∼8 nm after 30 h of milling. Parallel to this, the internal strain grows with increasing milling time, corresponding to a dislocation density of 1.5·1012 cm−2 after 30 h of milling. The lattice parameter of the copper matrix, which was calculated using the Rietveld method [11], is displayed as the function of the milling time in Fig. 4. During milling it increases until it reaches its maximum after 18 h. The maximum value for powder milled under argon is 0.3636(4) nm, and is higher than the matrix lattice parameter of powder milled in air, which is only 0.3634(2) nm (a0 (Cu) = 0.36151 nm). The subsequent decrease of the lattice parameter of the copper matrix after 18 h of milling is steeper for powder milled in air than for powder milled under argon. After 30 h of milling the lattice parameter of the copper matrix reaches 0.3633(5) nm for powder milled under argon, and 0.3628(5) nm for powder milled in air. The microhardness of the powder increases very rapidly within the first 10 h of milling (Fig. 5) until it saturates after ∼15 h. In the saturated state the powder milled under argon reveals a microhardness of ∼520 HV0.025 , whereas that milled in air exhibits values of ∼470 HV0.025 . 3.1.1. Heat treatment SEM pictures of different powders treated at selected temperatures are shown in Figs. 6 and 7. Niobium precipitates from the solid solution starting at ∼600 ◦ C, if the oxygen content is low (i.e. powder milled under argon). A heat treatment at 600 ◦ C for 1 h yields a size of the Nb precipitates of 40–80 nm (Fig. 6a). At higher temperatures, coarsening of the precipitates sets in leading to a grain size of 100–150
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Fig. 5. Microhardness of Cu–Nb alloy powder versus milling time.
nm at 700 ◦ C and 150–350 nm at 900 ◦ C (Fig. 6b). No precipitation of niobium can be detected in powders with a high oxygen content (i.e. powder milled in air ) heat treated at temperatures up to 900 ◦ C, as shown in Fig. 7.
After each of these particular heat treatments, a sample was taken for X-ray powder diffraction analysis in order to trace the phase formation in accordance with the heat treatment. Up to temperatures of 500 ◦ C no significant change of the powder diffraction pattern of the solid solution can be detected (Fig. 8). Higher temperatures cause the formation of Cu2 O as well as NbO. The intensity of the reflections related to these oxides is higher in powder with a higher oxygen content, even after a high temperature heat treatment. At heat treatment temperatures above 600 ◦ C, different scenarios occur, depending on the oxygen content of the powder. In powder with a low oxygen content (i.e. powder milled under argon) and heat treated at 600 ◦ C, niobium reflections can be detected. In contrast to this, reflections related to niobium can only be detected in powder with a high oxygen content when it is heat treated at 900 ◦ C. Due to the heat treatment, other microstructural properties change, such as the lattice parameter of the copper matrix, the dislocation density, the corresponding grain size and finally the microhardness. These properties are shown together in Fig. 9.
Fig. 6. SEM micrographs of Cu–Nb alloys after heat treatment for 1 h at 600 ◦ C (a) and 900 ◦ C (b) (the corresponding samples were milled under argon).
Fig. 7. SEM micrographs of Cu–Nb alloys after heat treatment for 1 h at 600 ◦ C (a) and 900 ◦ C (b) (the corresponding samples were milled in air).
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Fig. 8. X-ray powder diffraction patterns after heat treatment at different temperatures for 1 h (the corresponding samples were milled under argon (a) and in air (b)).
With increasing heat treatment temperature the lattice parameter of the copper matrix decreases until the value converges to that of pure copper. The lattice parameter of the copper matrix of the oxygen-poor sample is quite high. This is caused by the fact that niobium is dissolved into the copper matrix rather than being involved with the formation of phase impurities. This corresponds to a low fraction of phase impurities. Furthermore, either a longer heat treatment or a higher temperature is required to precipitate the niobium from the solid solution, as more niobium is still dissolved. The grain size of the copper matrix of heat treated powder increases with increasing heat treatment temperature. After ageing at 900 ◦ C for 1 h, the copper grain size is ∼30 nm. The dislocation density, which is calculated from the internal strain of the copper lattice,
decreases with increasing heat treatment temperature up to 600 ◦ C at which it reaches an equilibrium state where it stays constant at a level of ∼1·1010 cm−2 . The microhardness of the powder does not change due to heat treatments at temperatures below 300 ◦ C. With increasing heat treatment temperature, a small increase of hardness is observed. The maximum hardness is observed in the case of a powder milled under argon after applying a heat treatment at 400 ◦ C and in the case of a powder milled in air after a heat treatment at 500 ◦ C. A further increase of the heat treatment temperature leads to a decrease of the microhardness of the powder, roughly independent of the oxygen content of the powder. After a heat treatment at 900 ◦ C the microhardness of the powder is ∼400 HV0.025 , in both cases.
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Fig. 9. Lattice parameter of the copper matrix, dislocation density in the copper matrix, grain size of the copper matrix and microhardness of the powder (from top to bottom) as a function of the heat treatment temperature for a 1 h heat treatment.
4. Discussion Copper and niobium are mechanically alloyed to a supersaturated solid solution by means of high energy ball milling. Starting from elementary powders the process of mechanical alloying itself takes ∼30 h. Although the Cu–Nb system does not show a negative heat of mixing, a solid solution of copper and niobium is formed due to the dislocation pumping mechanism [12], which becomes effective due to the fine microstructure of the powder, which is formed during milling [9,13]. The milling atmosphere influences this process. For example, the lattice parameter of the copper matrix varies differently depending on whether milling takes place under argon or in air. In the powder milled under argon the lattice parameter of the copper matrix reaches a higher value than in the case of powder milled in air. According to the
dislocation pumping mechanism the niobium atoms situated at the Cu/Nb phase boundary diffuse along dislocation lines into the copper grains. The driving force for this mechanism is the interaction between the dislocation and the dissolved atom. The effect of the atmosphere on the lattice parameter is caused by oxygen. The oxygen atoms more likely diffuse along the dislocation lines than within the grains. Thus both oxygen and niobium are situated along the dislocation lines where they are enriched. Subsequently, niobium oxide particles are formed instead of the copper niobium solid solution. The more oxygen is present in the powder the more niobium is oxidised and thus less niobium is brought into solid solution. In the powder milled in air the formation of niobium oxide prevails over the dissolution of niobium atoms in copper. This leads to the decrease of the lattice parameter of the copper matrix after 15–18 h of milling (Fig. 4). Furthermore, there are no niobium precipitates found in powder with a high oxygen content after heat treatments up to 900 ◦ C, providing further evidence for the effect of atmosphere on the phase forming process during milling. Within this powder, only niobium oxide can be detected either by SEM or X-ray diffraction. On the other side, there are niobium precipitates present in the heat treated oxygen-poor powder, although there is still some niobium oxide found in these samples, but less than in the others. This niobium oxide is already formed during the late stages of milling, characterised by the almost constant level of the lattice parameter of the copper matrix after 18 h of milling. The oxygen content is not high enough to oxidise the niobium completely and accordingly, niobium precipitates can be detected after heat treatments, even at relatively low temperatures. The decrease of the lattice parameter of the copper matrix with increasing heat treatment temperature indicates an ongoing precipitation of niobium from the supersaturated solid solution in the temperature range between 200 and 600 ◦ C. Above this temperature the lattice parameter of the copper matrix does not change any more and is equal to that of pure copper. Analysing the powder with a high oxygen content heat treated at 900 ◦ C, only small niobium reflections are detectable by means of X-ray analysis, whereas no niobium precipitates can be found by means of the SEM investigations, probably because of the very small precipitate size. This leads to the assumption that niobium, which precipitated from the solid solution, reacts with oxygen atoms and forms niobium oxide. The small amount of niobium precipitates, whose reflections are detected by X-ray analysis, originates from regions where the level of niobium and oxygen atoms becomes unbalanced and finally favors the precipitation of niobium instead of the formation of niobium oxide from the entire niobium content. The development of the microhardness of the heat treated powder is controlled by three processes: (i) precipitation of niobium, (ii) coarsing of the precipitates (in the case of powder with a low oxygen content) or of the niobium oxide particles (in the case of powder with a high oxygen content), and (iii) recrystallisation of the copper matrix and growth of
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the copper grains. These processes, as shown in Fig. 9, take place during the heat treatment and influence the properties of the Cu–Nb alloy. In particular, the formation of niobium precipitates leads to an increase of the microhardness and, conversely, the recrystallisation process, grain growth and precipitation coarsening lead to its decrease. At temperatures below 300 ◦ C, the softening effect of the recrystallisation process compensates for the strengthening effect caused by the formation of Nb or NbO precipitates. In the temperature range between 300 and 400 ◦ C, Nb precipitates are formed. NbO is formed in the case of a high oxygen content of the powders in the range between 300 and 500 ◦ C. At higher temperatures (up to 600 ◦ C) the microhardness decreases as a consequence of the recrystallisation processes. At temperatures above 600 ◦ C the microhardness value is controlled by the precipitation coarsening and grain growth of the copper matrix, which lead to a decrease of the microhardness to 400 HV0.025 after a heat treatment at 900 ◦ C. With respect to the development of the microhardness, an intermediate heat treatment temperature is supposed to generate niobium precipitates with an optimum size, resulting in the best hardening properties. Unfortunately, after a heat treatment at 400 ◦ C, the size of the precipitates cannot be determined as they are too small to be detected by SEM analysis. Furthermore, the lattice parameter of the copper matrix is still higher than that of pure copper. This indicates that a part of the niobium atoms still remains within the solid solution after a heat treatment at 400 ◦ C. This results in an appreciably low electrical conductivity. The optimum heat treatment produces a balance between the mechanical and electrical properties and should be in the range of 600–700 ◦ C.
5. Conclusions The supersaturated solid solution of niobium in copper is formed by mechanical alloying of elementary copper and niobium powder after 30 h of milling. The milling atmosphere influences this process: in the powder milled in air (1.8 ma.% oxygen content after 30 h of milling) niobium oxide particles are formed during the late stage of milling. Thus, almost all the niobium atoms are bound to oxygen atoms. No niobium precipitates can be formed in such powder during the subsequent heat treatments at temperatures up
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to 900 ◦ C, only niobium oxide. In contrast to this, niobium precipitates in powder with a low oxygen content (milled under argon; 1.2 ma.% oxygen content after 30 h of milling) during heat treatment. The decrease of the microhardness of the powder during the heat treatment is negligible. The microhardness value of powder with a low oxygen content treated at 700 ◦ C is ∼470 HV0.025 corresponding to a UTS of ∼1.5 GPa. A significant increase of the electrical properties of heat treated powder is expected after the complete precipitation of niobium from the solid solution. This investigation has shown that the mechanical alloying route with a subsequent heat treatment is a promising way of producing a conductor material with a high mechanical strength, as required for pulsed high-field magnets [14].
Acknowledgements This work was supported by the Bundesministerium für Bildung und Forschung by grant No. 03SC5DRE.
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