The effects of addition of Zr and Ag on the mechanical properties of a Cu-0.5wt%Cr alloy have been investigated. The addition of 0.15wt%Zr enhances the strength and resistance to stress relaxation of the Cu-Cr alloy. The increase in strength is caused by both the decrease in inter-precipitate spacing of Cr precipitates and the precipitation of Cu 5 Zr phase. The stress relaxation resistance is improved by the preferentially forming Cu 5 Zr precipitates on dislocations, in addition to Cr precipitates on dislocations. The addition of 0.1wt%Ag to the Cu-Cr and Cu-Cr-Zr alloys improves the strength, stress relaxation resistance and bend formability of these alloys. The increase in strength and stress relaxation resistance is ascribed to the decrease in inter-precipitate spacing of Cr precipitates and the suppression of recovery during aging, and to the Ag-atom-drag effect on dislocation motion.The better bend formability of the Ag-added alloys is explained in terms of the larger post-uniform elongation of the alloys.

The Ostwald ripening of Al 3 Sc precipitates in an Al-0.28 wt pct Sc alloy during aging at 673, 698, and 723 K has been examined by measuring the average size of precipitates by transmission electron microscopy (TEM) and the reduction in Sc concentration in the Al matrix with aging time, t, by electrical resistivity. The coarsening kinetics of Al 3 Sc precipitates obey the t 1/3 time law, as predicted by the Lifshitz-Slyozov-Wagner (LSW) theory. The kinetics of the reduction of Sc concentration with t are consistent with the predicted t Ϫ1/3 time law. Application of the LSW theory has enabled independent calculation of the Al/Al 3 Sc interface energy, ␥, and volume diffusion coefficient, D, of Sc in Al during coarsening of precipitates. The Gibbs-Thompson equation has been used to give a value of ␥ using coarsening data obtained from TEM and electrical resistivity measurements. The value of ␥ estimated from the LSW theory is 218 mJ m Ϫ2 , which is nearly identical to 230 mJ m Ϫ2 from the Gibbs-Thompson equation. The pre-exponential factor and activation energy for diffusion of Sc in Al are determined to be (7.2 Ϯ 6.0) ϫ 10 Ϫ4 m 2 s Ϫ1 and 176 Ϯ 9 kJ mol Ϫ1 , respectively. The values are in agreement with those for diffusion of Sc in Al obtained from tracer diffusion measurements.dependence of the coarsening behavior of precipitates in an Al-0.3 wt pct Sc alloy aged at 573 to 723 K. In addition, the diffusivity of Sc in Al is obtained from coarsening experiments. Iwamura and Miura [10] have examined the coarsening behavior of Al 3 Sc precipitates in an Al-0.2 wt pct Sc alloy at 673 to 763 K on the basis of TEM observations with the numerical model. The radius for coherent/semicoherent transition of the precipitates is determined from TEM images as 15 to 40 nm. The average radius, r, of the Al 3 Sc precipitates obeys the r 3 growth law both in the coherent stage (r Ͻ 15 nm) and in the semicoherent stage (r Ͼ 40 nm). However, in the intermediate stage, where coherent and semicoherent precipitates coexist (15 Ͻ r Ͻ 40 nm), coarsening is delayed.By measuring independently both the growth rate of precipitates and the rate of depletion of the matrix supersaturation during coarsening of the precipitates, independent reliable values of the matrix/precipitate interface energy and the diffusion coefficient of solute in the matrix can be determined from coarsening data alone. [6] This approach has already been applied in several binary systems. [11][12][13][14][15] The interface energy can also be obtained from measurements of the matrix solute concentration during coarsening, together with the mean precipitate size, using the Gibbs-Thompson equation. This direct application of the Gibbs-Thompson equation has been directed at Al-Li [16] and Cu-Ni-P systems. [17] In this work, the coarsening kinetics of the Al 3 Sc precipitates in an Al-0.28 wt pct Sc alloy aged at 673, 698, and 723 K have been investigated. The Al/Al 3 Sc interface energy, ␥, and the diffusivity of Sc in Al have been independently derived from measurements o...

High-resolution transmission electron microscopy and conventional transmission electron microscopy have been used to investigate in detail the transformation processes of twinned 9R copper precipitates to ® nal stable structures via a 3R structure in a thermally aged Fe± Co model alloy. In 9R precipitates of size greater than about 13 nm, the motion of the twin boundaries and the elimination of the regular stacking faults on every third (009) 9R close packed plane were observed to occur nearly simultaneously.Direct evidence was found that 3R is a non-cubic structure obtained when the regular stacking faults on the (009) 9R basal planes are removed. In larger precipitates (of size up to about 26 nm), lattice plane rotations and plane spacing changes in 3R variants took place towards stable fcc and fct structures, suggesting the occurrence of lattice relaxation involving the di usion of atoms. The fct structure had larger lattice constants aˆbˆ0:369 nm and cˆ0:366 nm than bulk fcc copper (aˆ0:361 nm). Precipitates of size 26± 40 nm consisted of nearly twin-related variants with both fcc and fct segments, and aligned with the iron matrix according to the Kurdjumov± Sachs orientation relationship. Larger precipitates were observed to be fcc and fct single crystals.

The precipitation processes in a Cu-0.9 mass%Be alloy single crystal containing only G.P. zones parallel to the matrix (001) plane have been studied by high-resolution transmission electron microscopy. The precipitation sequence found is: G.P. zones ! 00 ! 0 I ! I þ 0 ! . The 00 phase has a two-layer structure of Be atoms separated by a matrix layer parallel to (001) , and a body-centered tetragonal (bct) lattice. The 2 2Þ habit plane and the orientation relationship, ð1 " 1 1 " 2 2Þ == ð0 " 1 1 " 2 2Þ 0 ; ½110 == ½100 0 (OR I), or the ð1 " 1 13Þ habit plane and the orientation relationship, ð1 " 1 13Þ == ð0 " 1 13Þ 0 ; ½110 == ½100 0 (OR II). The structure of 0 successively changes into that of the phase with a ¼ 0:280 nm for OR I or 0.268 nm for OR II. The precipitates of OR I and OR II are elongated along approximately ½1 " 1 11 and ½3 " 3 3 " 2 2 , which are in good agreement with the invariant-line directions.