PHASE SEPARATION BY INTERNAL OXIDATION AND REDUCTION IN A Cu-5at%Ni-ALLOY

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1 Pergamon PII S (98) Scripta Materialia, Vol. 39, No. 1, pp , 1998 Elsevier Science Ltd Copyright 1998 Acta Metallurgica Inc. Printed in the USA. All rights reserved /98 $ PHASE SEPARATION BY INTERNAL OXIDATION AND REDUCTION IN A Cu-5at%Ni-ALLOY Ruth Lüke a, Joachim Bankmann a, Peter-J. Wilbrandt a, Wilfried Erfurth b and Reiner Kirchheim a a Institut für Materialphysik and SFB 345, Universität Göttingen, Hospitalstraße 3-7, Göttingen, Germany b Max Planck Institut für Mikrostrukturphysik, Am Weinberg 2, Halle/Saale, Germany (Received February 19, 1998) (Accepted in revised form March 25, 1998) 1. Introduction The basic idea of our experiments is to separate the constituents of a completely miscible alloy by internal oxidation and following reduction. Figure 1 illustrates the principle of such a phase separation: The preparation procedure starts with a homogeneous alloy AB in the thermodynamic equilibrium. First internal oxidation transfers the homogeneous AB alloy into a two phase system of A matrix and precipitates of B-oxide. This decreases the Gibb s free energy. The treatment will only be feasible, if the following two prerequisites are fulfilled: (1) The oxygen affinity of the component B has to be sufficiently greater than that of A in order to oxidize only B and not the whole material. (2) Oxygen has to diffuse more rapidly in A than B. Otherwise, the B atoms will diffuse to the surface and form an external oxide layer. In a second step a reduction treatment removes the oxygen from the composite. By exposing the sample surface to an oxygen sink an oxygen gradient is established and the oxide precipitates decompose to oxygen dissolved and B precipitated in the matrix. The oxygen then diffusing to the surface is removed by the reducing agent. An interdiffusion of A and B during the reduction treatment is suppressed by the lower reaction temperature during the reduction process. Under this conditions the diffusivity of oxygen has to be much higher than the one of B, whereas at higher temperatures the difference of the diffusivities between O and Ni is usually smaller because of the smaller activation energy for interstitial diffusion. The Gibb s free energy of the new alloy increases and, consequently, the system reaches a state far away from thermodynamic equilibrium. A detailed theoretical description of the internal oxidation and reduction can be found elsewhere [1,2]. The feasibility of the described phase separation procedure is demonstrated for a completely miscible Cu-5at%Ni alloy. The components of this alloy fulfills the requirements mentioned above for chemical behavior and diffusivity [3]. 2. Experimental Procedures The Cu-5at%Ni alloy was prepared by induction melting of pure copper (99,9998 at%) and nickel (99,995 at %) in an Al 2 O 3 -crucible under argon atmosphere. The raw material was cut by spark erosion 73

2 74 PHASE SEPARATION BY INTERNAL OXIDATION Vol. 39, No. 1 Figure 1. Phase separation of a completely miscible alloy by internal oxidation and reduction. in 1 mm thick discs. After rolling the discs to a thickness of 130 m they were homogenized by annealing for 20 days at 1273 K in an argon atmosphere. Their microstructure consists of nearly spherical grains with diameters between 100 and 500 m. For the internal oxidation the sheets were put into a mixture of Cu plus Cu 2 O plus Al 2 O 3 powders (1:1:1 by volume) and heated for 10 h at 1273 K in an argon atmosphere. This treatment is known as the Rhines-Pack method [4]. During the annealing the Cu 2 O decomposes into its components and acts as an oxygen source for the oxidation process. The oxygen partial pressure is equal to the decomposition pressure of cuprous oxide (p O Pa at 1273 K [5]). The oxygen diffuses into the sample and oxidizes the Ni to stable NiO. Oxidation of Cu-matrix is impossible at this oxygen partial pressure. For the reduction treatment the samples were annealed in quartz-tubes with a piece of graphite under 100 mbar CO 2 for 1 hour at 773 K. CO 2 and graphite react to CO and CO acts as a sink for oxygen. The fraction of CO depends on temperature and pressure and is determined by the Boudouard- Equilibrium [6]. The carbon monoxide reduces the NiO to Ni. After the oxidation and after the reduction treatment the samples were investigated by Secondary Ion Mass Spectrometry (SIMS) and Transmission Electron Microscopy (TEM) and Energy Dispersive X-Ray Analysis (EDX). The TEM specimens were prepared from small discs punched out of the internally oxidized sheets. After thinning by mechanical polishing they were dimpled to thicknesses between 20 and 50 m and finally thinned by ion-milling. For the TEM observations a Philips EM 420ST and a Philips CM 200 microscopes were used. They are both equipped with an energy dispersive X-ray analysis system allowing a light element analysis down to boron. The microscopes were operated at 120 kv. For the SIMS investigations the oxidized sheets were etched with a solution of 55 vol% CH 3 COOH, 30 vol% HNO 3 and 15 vol% H 3 PO 4. After etching for 30 seconds the samples were cleaned with distilled water. The SIMS investigations were made with a Cameca TOF-SIMS IV. The depth profiles were obtained using the dual beam technique, i.e. different kinds of ions were used for sputtering and analysis. The crater was formed with argon ions of 1 kev. For the analysis a focussed Ga beam (25 kev, beam diameter 200 nm, 0.5 pa beam current) was used. It was scanned over the region of interest an the bottom of the crater [7].

3 Vol. 39, No. 1 PHASE SEPARATION BY INTERNAL OXIDATION 75 Figure 2. NiO precipitates in an internal oxidized Cu 5at% Ni alloy; a: TEM bright field micrograph; b: Results of the EDX measurement along the line in 2a. The count rates are scaled by the maximum count rate of the elements in cps (counts per second). 3. Results and Discussion Figure 2a is a bright field TEM micrograph of an internally oxidized Cu-5at% Ni sample. It shows two NiO precipitates surrounded by the Cu-matrix. The single-crystalline precipitates have diameters of several m. Because of the different sputtering yields of NiO and Cu during the ion-milling a hole was formed between the precipitates. The arrow in the micrograph marks the position of the EDX linescan. In Fig. 2b the count rates of the three elements Cu, Ni and O are plotted for the different positions on the sample. Within the NiO precipitates the count rates of Ni and O are high and the one of Cu is low. Generally, in adjacent areas of NiO and Cu-matrix the count rates of Ni are higher than the count rates of Cu (Ni max: 1726, Cu max: 507). This is due to the greater thickness of the NiO precipitates compared to the matrix. Correspondingly, the varying count rates of Ni and O over the precipitates result from the thickness variations within the precipitates. The low count rates of O compared with Ni are caused by the small detection efficiency for low energy lines [8]. In the bright field TEM micrograph of a reduced specimen a precipitate is surrounded by the Cu-matrix (Fig. 3a). The arrow marks the analyzed area of the EDX linescan. The precipitates consists of nearly pure Ni. The O count rate Fig. 3b has decreased drastically compared to Fig. 2b. As a consequence of the reduction treatment, along the whole line only background intensity is observed. The course of the Ni count rate proves that the reduction treatment leaded to the expected results. Still the Ni count rate varies in the opposite way as the one of Cu. It is only large in the precipitate. The variation of the Cu count rate is similar to the results of the oxidized specimen (Fig. 2b). The count rate is large in the matrix and decreases in the precipitate. The interdiffusion between Cu and Ni could be successfully suppressed. The interface between the matrix and the reduced precipitate is still visible, but it is somewhat blurred compared with Fig. 2a. The higher maximum count rate of Cu indicates that the matrix is thicker than the reduced precipitate. This is considered to be a result of the specimen geometry. The thinned regions constrict themselves during the reduction treatment as a consequence of the surface tension. This effect is not so pronounced in the precipitates because the volume of the precipitates decreases continuously due to the loss of oxygen.

4 76 PHASE SEPARATION BY INTERNAL OXIDATION Vol. 39, No. 1 Figure 3. A reduced precipitate surrounded by Cu-matrix; a: TEM bright field micrograph; b: Results of the EDX measurement along the line in 3a. The count rates are scaled by the maximum count rates of the elements in cps (counts per second). The reduction of thin samples is advantageous because the thin regions can be reduced in a short time. The oxygen need not diffuse through the Cu-matrix, like in bulk samples. The reduction temperature can be kept so low that the diffusion of Cu and Ni is small ( 100 nm at 773 K, 1h). This leads to the observed small interdiffusion between Cu and Ni. Figure 4. SIMS depth profiles; a: NiO precipitate; b: pure NiO; c: reduced precipitate; d: pure Ni.

5 Vol. 39, No. 1 PHASE SEPARATION BY INTERNAL OXIDATION 77 A further proof for the phase separation in the Cu-5at%Ni alloy are the SIMS results in Figure 4. The diagrams show depth profiles obtained at the surface of precipitates in an internal oxidized specimen (Fig. 4a) and a specimen after the reduction treatment (Fig. 4c). As reference for the oxidized and reduced state specimens of pure NiO (Fig. 4b) and pure Ni (Fig. 4d) were used. Regions of 1 1 m 2 were analyzed. The variations of O, NiO 2 and NiO count rates are shown. The intensity of ions does not correspond to the stoichiometric composition of the specimen (here NiO 2 ) as it is typical for SIMS measurements. In the NiO standard the O count rates are the highest ( cps). The count rates of NiO 2 and NiO ions are one and two orders of magnitude smaller than the O count rates. Along the whole depth profile the count rates are nearly constant. To the contrary, in the depth profile of pure Ni the count rates for the three kinds of ions decrease down to zero within the first 600 s. This effect is caused by a surface oxide layer on the specimen. The small count rates result from the absence of oxygen in pure Ni. This is evidenced by the fact, that ionization probabilities of pure metals are small. However, reacting the metal with oxygen leads to an instantaneous increase of the ionization probabilities and, consquently, of the count rates. Therefore, only the profile after 600 s is considered as representative for pure Ni. In the internal oxidized specimen the count rates are similar to the count rates of the NiO standard. The O count rates are the highest, the count rates of NiO 2 and NiO are one and two decades smaller. The differences of one decade between the count rates of the NiO precipitate and the standard result from differences of the beam current. For the comparison of the results the ratios between the different count rates were calculated. The results for of O /NiO 2,O /NiO and NiO 2 /NiO in the specimen and the standard differ by less than 10%. The reduced precipitate (Fig. 4c) corresponds to pure Ni, because the count rate of O is nearly zero. NiO 2 and NiO ions were not detected. Consequently, the significant differences between the depth profile of the NiO precipitate (Fig. 4a) and the reduced precipitate in Fig. 4c indicate that our reduction treatment reduces the NiO to Ni. 4. Conclusions Our TEM/EDX and SIMS investigations demonstrate, that in thin specimens of a homogeneous Cu-5at% Ni alloy Ni and Cu can be separated by internal oxidation and following reduction. However, in bulk samples the reduction process is to slow and the interdiffusion between Ni and Cu-matrix can not be suppressed. In thin samples the NiO precipitates are near the surface. The reduction time is so short that the intermixing of the components by interdiffusion is not possible. The described procedure may be a promising method to produce new material with interesting properties. References 1. J. L. Meijering, Internal Oxidations in Alloys, Wiley, New York (1971). 2. R. Kirchheim and M. Henning, Vaccuumtechnik. 4, 100 (1979). 3. T. B. Massalsky, ed., Binary Alloy Phase Diagrams ASM, Metals Park, OH (1990). 4. F. N. Rhines, Trans AIME. 137, 246 (1940). 5. H. Jang, D. N. Seidman, and K. L. Merkle, Interface Sci. 1, 61 (1993). 6. E. Wiberg, Anorganische Chemie, W. de Gruyter, Berlin (1951). 7. K. Iltgen, C. Bendel, E. Niehuis, and A. Benninghoven, in Proceedings of the 10th International Conference on SIMS 1995, p. 375, Wiley, New York (1997). 8. A. H. Foitzik, J. S. Sears, S.-Q. Xiao, and A. H. Heuer, Ultramicroscopy. 50, 207 (1993).