Electron microscopy nanoscale characterization of ball milled Cu-Ag powders. Part II: Nanocomposites synthesized by elevated temperature milling or annealing

Acta Materialia 50 (2002) 4711–4726 www.actamat-journals.com

Electron microscopy nanoscale characterization of ball milled Cu-Ag powders. Part II: Nanocomposites synthesized by elevated temperature milling or annealing S. Zghal a, R. Twesten a, Fang Wu b, P. Bellon a b

a Frederick Seitz Materials Research Laboratory, University of Illinois at Urbana-Champaign, IL 61801, USA Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, Urbana, IL 61801, USA

Received 25 October 2001; received in revised form 8 May 2002; accepted 4 July 2002

Abstract Microstructures and phases stabilized at steady state by variable temperature ball milling of Cu50Ag50 powders are characterized using transmission and scanning transmission electron microscopy. Starting from chemically mixed and cold-worked powders obtained by room temperature milling, it is shown that, upon increasing the milling temperature, the material first decomposes into Cu-rich and Ag-rich solid solutions, and then recrystallizes. A similar sequence is observed during the static annealing of the solid solution precursor. In both cases, Cu-Ag nanocomposites are synthesized, at a scale of a few nanometers in the unrecrystallized state, and at a scale ranging from 30 nm after dynamic recrystallization to 75 nm after static recrystallization. These nanocomposites exhibit high hardness values, approaching 6 GPa. Interestingly enough, recrystallization leads to an increase in the hardness of these materials.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Mechanical alloying; Transmission electron microscopy; Alloys (copper, silver); Nanocomposites; Recrystallization

1. Introduction Ball milling (mechanical alloying) has become a widely used technique to process powder materials and to synthesize a large range of nonequilibrium phases, from amorphous materials, to nanocrystalline phases, to extended solid solutions [1–3]. In the case of moderately immiscible elements, such as Cu-Ag for which the heat of mixing is around 6 kJ/g.atom, it is even possible to fully control the phases in presence at steady-state during milling by varying the milling temperature [4,5]. As clearly demonstrated by Klassen and coworkers [5], low temperature milling results in

the stabilization of solid solutions, whereas elevated temperature milling results in the coexistence of two terminal solid solutions. This temperature effect is best explained by considering such alloys during milling as ‘driven systems’: while thermal diffusion tends to promote phase separation because of the positive heat of mixing, this dynamic competes with an atomic mixing forced by the sustained plastic deformation of the material [6]. When deformation results from dislocation glide, and in the absence of plastic flow localization, this forced mixing can be considered as random. At low temperature, thermal diffusion, even enhanced by point defect supersaturation, becomes

1359-6454/02/$22.00.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 2 8 6 - 0

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so sluggish that the evolution of the alloy is controlled by the forced mixing, thus leading to the stabilization of solid solutions. As the temperature increases, the balance between the two dynamics is reversed, leading to phase separation. Driven system approaches have also been successfully used to analyze and rationalize crystal-amorphous and order-disorder reactions during ball milling [7–9]. Atomistic computer simulations and analytical modeling based on the above driven system approach [10] have also revealed that, because of the macroscopic nature of the forced dynamics, phase decomposition should saturate once the phases have reached a certain size, leading to the stabilization of nanocomposites. The idea that precipitate size can be adjusted by ball milling was in fact empirically used early on by Benjamin to strengthen Ni-base alloys with a dispersion of nanometer-size oxide precipitates [11]. By choosing milling conditions that create the right balance between thermal decomposition and plasticityinduced chemical mixing, ball milling offers a simple and yet generic method to synthesize nanocomposites. Recently we have used atom probe field ion microscopy (APFIM) to study the distribution of chemical species in Cu50Ag50 powders milled at variable temperatures [12]. We have indeed found that nanocomposites can be obtained by intermediate temperature milling, leading for instance to domain sizes around 30 nm when milling is performed at 453 K. Another route for the synthesis of nanocomposites has been proposed by He and coworkers [13]: a solid solution precursor is first synthesized by low temperature milling, and the nanocomposite then forms by partial decomposition during heat treatment, for instance during the hot consolidation of the powders. In the preceding paper, here referred to as Part I [14], we have shown that low temperature milling, in addition to the stabilization of a solid solution, results in the formation of heavily coldworked and textured microstructures. Now starting from this state as a precursor and increasing the milling temperature, it is expected that the material can undergo recovery or recrystallization, in addition to the phase decomposition previously discussed. Different nanocomposites can be

obtained depending upon which reaction takes place first. A similar competition between recrystallization and decomposition also exists if we simply isothermally anneal these precursor powders. This latter situation has in fact been extensively studied by Hornbogen and coworkers (see ref. [15]). When decomposition takes place first, as will be reported here for our powder materials, one is then left with a situation that is similar to the study of recrystallization in duplex alloys. It is directly relevant to the present study to remark that thermomechanical treatments relying on chemical decomposition and recrystallization are commonly used to refine microstructures of duplex alloys leading to superplastic properties [16]. The present study can be seen as an extension of this refining process to produce ‘nano-duplex’ alloys, more commonly referred to as nanocomposites. The Cu-Ag system has been selected for the present study because its thermodynamic properties are well known [17], and because dynamical phase equilibrium during ball milling is conveniently controlled by adjusting the milling temperature. Combining X-ray diffraction (XRD) and differential scanning calorimetry (DSC) measurements, Klassen and coworkers [5] have shown that for any set of control parameters in a milling experiment only one steady-state is reached, regardless of the initial state of the powders. At liquid nitrogen temperature, milling results in the stabilization of a single, random fcc solid solution. At 316 K the steady-state is also a single fcc phase but the degree of chemical mixing is only 70 to 75%, as calculated from comparing integrated intensities in XRD spectra or in DSC traces. Between 331 and 423 K, the steady-state XRD spectra indicate the coexistence of three fcc solid solutions, one Cu-rich, one Ag-rich and one with the nominal composition, Cu50Ag50. At 473 K the steady-state consists of two fcc solid solutions, one Cu-rich, one Ag-rich. In the present paper, we use transmission electron microscopy (TEM), and scanning TEM (STEM) to characterize the microstructure and the phases in presence in Cu50Ag50 powders ball milled at temperatures ranging from near room temperature (315 K) to 503 K. We also investigate the evolution of 315 K milled powders during ther-

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mal annealing at 503 K. Our main aim is to investigate various processing routes that can lead to the synthesis of nanocomposite materials. Nanoindentation is used to measure the hardness of these nanocomposites. The experimental procedures used for milling the powders and for electron microscopy analysis are presented in section 2. The results obtained by TEM, STEM, and nanoindentation are presented in section 3, and then discussed in section 4.

2. Experimental procedures Ball milling of Cu and Ag powders was performed in a SPEX 8000 shaker mill following the procedure used by Klassen and coworkers [5]. Powders of copper and silver were mixed according to the desired nominal (atomic) composition, Cu50Ag50. A purified argon atmosphere was used for all temperatures. The milling temperatures used in this paper range from 315 K (milling performed in the absence of external heating) to 503 K. For elevated milling temperatures, heating tapes were wrapped around the vial to reach the desired temperature. All milling performed at a temperature different from 315 K were preceded by a 24-h milling treatment at 315 K so as to alloy the initial powders, and then milled for 5 h at the chosen temperature. These milling sequences and times are such that in all cases the powders have reached the steady-state characteristic of the milling temperature. These steady-states were determined from XRD and DSC measurements: these results are fully consistent with those reported by Klassen et al. [5] and are therefore not presented here. The evolution of the milled powders during a thermal annealing at 503 K was studied by annealing the powders in the milling vial before any exposure to air, for times ranging from 10 min to 10 h. Samples suitable for transmission electron microscopy were prepared following the procedure described in Part I. The typical foil thickness was 10–15 nm. TEM was performed on a CM12 microscope operated at 120 kV, while STEM was performed on a Vacuum Generator HB-501 microscope operated at 100 kV. Lattice parameter measurements were performed after a careful calibration of the microscope,

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allowing an accuracy of ± 0.002 to ± 0.004 nm, depending upon the sharpness of the diffraction patterns (See Part I for details, as well as for other TEM and STEM procedures). Nanoindentation tests were performed in order to determine the hardness of the various nanostructures we synthesized. Measurements were carried out on a Hysitron Triboscope nanomechanical testing system attached to a Digital Instrument Multimode Atom Force Microscope (AFM). Powder particles were mounted in resin and mechanically polished in preparation for the nanoindentation tests. Atomic force microscopy was used to image the surface and to select smooth areas for nanoindentations (roughness was typically 10 nm or less over 3 µm×3 µm areas). A Berkovich indenter was used, and all reported measurements were obtained using a 500 µN maximum load with a 100 µN/s loading rate. The penetration depth was in the range of 50–80 nm. Different loads, ranging from 200 µN to 2 mN, were also used for comparison, but no indication of indentation size effect was observed in this load range.

3. Results We now present the results obtained upon increasing the milling temperature. Three different microstructures are stabilized at steady state, and we have thus selected three temperatures, 315, 393 and 503 K to present these results. We then present the microstructural evolutions of powders milled at room temperature and subsequently annealed at 503 K. 3.1. Powders milled at 315 K The powders milled at 315 K (Fig. 1) display characteristics similar to those reported in Part I for cryo-milled powders. As for these samples, only one Cu-Ag solid solution is detected in selected area diffraction (SAD) patterns (Fig. 1c), with an average lattice parameter of 0.389 ± 0.004 nm. Notice that the reflections are however more diffuse than after cryomilling, and that Cu2O precipitates are still detected but their density, 4 × 1015 cm⫺3, is smaller (see Part I for details on lattice

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Fig. 1. TEM analysis of Cu50Ag50 powders milled at 315K: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern; (d) is an STEM-DF image (notice the mottled contrast). The arrows in (a) point to Cu2O precipitates. Here and the following TEM analysis, the DF image is formed by selecting one of the most intense part of the (111) reflections, and the aperture used to form the DP has a diameter of 350 nm.

spacing and Cu2O precipitate density measurements). This is probably due to the fact that cryo-milling was performed in air, while for all other milling temperatures milling was performed in an argon purified glove box. In bright field (BF) mode (Fig. 1a), a complex contrast is observed, combining diffraction contrast and strain contrast from small localized defects. Compared to cryomilled powders, the texture appears to be less pronounced (Fig. 1c) and the regions illuminated in DF mode are more ramified (Fig. 1b). In STEM, a small scale contrast is observed in ADF mode (Fig. 1d), as for the cryo-milled samples. This mottled contrast with a 2 nm scale, however, differs from the elongated contrast with

a 5 nm scale observed in cryomilled samples (Fig. 3b in Part I). STEM-EDS analysis indicates that the composition of the sample is still fairly homogeneous (Fig. 2a,b), but the amplitude of the composition fluctuations has clearly increased (Fig. 2c) compared to that of cryomilled powders (Fig. 4c in Part I). 3.2. Powders milled at 393 K While the sample remains textured as for 315 K milling (Fig. 3a,b), the powder is now chemically decomposed. Indeed, more than one Cu-Ag solid solution is found in the SAD pattern (Fig. 3c). The analysis of this SAD pattern is difficult because of

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Fig. 2. STEM chemical characterization of Cu50Ag50 powders milled at 315 K: (a) ADF image, (b) Cu EDS map, (c) 3 nm wide Cu compositional profile along the line shown in white in (b).

the potential overlap of the {200} reflections of Ag-rich phases with the {111} reflections of Curich phases and because of the diffuse character of the reflections. One can nevertheless detect the presence of two main phases: a Ag-rich phase identified from its {111} reflections, yielding an average lattice parameter of 0.402 ± 0.004 nm, and a Cu-rich phase identified from its {200} reflections, yielding an average lattice parameter of 0.366 ± 0.004 nm. These lattice parameters correspond to 83 at.% Ag and 92 at.% Cu, respectively [18]. It is possible that other solid solutions with intermediate compositions are present as well, but no firm conclusion can be reached from Fig. 3c. In addition, Cu2O precipitates are again found (Fig. 3a), with a density of 1×1016 cm⫺3, as measured from STEM-ADF images. In STEM, a fine mottled contrast is observed in DF images, as for the powders milled at 315 K (compare Fig. 1d and Fig. 3d). ADF and EDS images (Fig. 4) both indicate that the composition field appears more or less uniform, and the Cu concentration varies from 60 at.% to 40 at.% (see for instance the second half of the line scan in Fig. 4c), though some regions show clear signs of

decomposition, as the Cu concentration changes abruptly from 70 at.% to 30 at.% (see for instance the two peaks near the center of the line scan in Fig. 4c), These values, however, do not correspond to the compositions determined from lattice parameter measurements in the previous paragraph. It is therefore concluded that the scale of decomposition is smaller than the thickness of the film (10– 15 nm), which results in an apparent reduction of the amplitude of the composition variations. It is worth noticing that the most visible composition variations in Fig. 4b,c have a characteristic length of about 10 nm, i.e., close to the sample thickness. It is reasonable to propose that this is why these composition variations are fairly visible in Fig. 4b,c. 3.3. Powders milled at 503 K The microstructure of these powders is markedly different. First, nanograins are now clearly identified in BF or DF mode (Fig. 5a,b), with an estimated average grain size of 20 nm. This estimation is in good agreement with the 17 nm value determined by Klassen and coworkers using a William-

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Fig. 3. TEM analysis of Cu50Ag50 powders milled at 393 K: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern; (d) is an STEM-DF image (notice again the mottled contrast).

son-Hall analysis of XRD spectra [5]. These grains are almost randomly oriented, as indicated by the formation of continuous rings in the SAD pattern (Fig. 5c), even though some intensity reinforcements indicate that a weak texture is still present. Two sets of fcc reflections are easily identified, with average lattice parameters 0.407 ± 0.002 nm and 0.362 ± 0.002 nm. These lattice parameters correspond to 95 at.% Ag and 99 at.% Cu respectively [18]. The STEM analytical analysis performed in ADF mode (Fig. 6a) and by EDS (Fig. 6b) reveal the formation of a nanocomposite with several remarkable features. First, the scale of decomposition, around 30 nm as determined from EDS profiles (Fig. 6d), corresponds to twice the grain size. EDS maps (Fig. 6b), combined with

bright field images (Fig. 6c) show that each grain is either Cu-rich or Ag-rich, and that Cu-rich grains are predominantly surrounded by Ag-rich grains. This can also be directly seen from Fig. 6b: the Ag-rich phase forms a clearly connected phase, while the Cu-rich phase is mostly comprised of small isolated agglomerates that contains a few connected Cu-rich grains. This lack of connectivity of the Cu-rich phase, viewed in projection, may be surprising since the volume fractions expected from the measured compositions (via the lattice parameters) are quite similar, 52 and 48% for the Ag-rich and Cu-rich phases, respectively. Furthermore these values are large enough for both phases to percolate, if they were randomly distributed. In order to fully determine the connectivity of the Cu-

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Fig. 4. STEM chemical characterization of Cu50Ag50 powders milled at 393 K: (a) ADF image, (b) Cu EDS map, (c) 3 nm wide Cu compositional profile along the line shown in white in (b).

rich phase, however, a three-dimensional analysis would be required, for instance using three dimensional atom probe field ion microscopy. Finally, it is important to recognize that the shapes of the Curich regions are quite elongated and distorted. These shapes are typical of a material that has undergone plastic deformation. We will see that different shapes are obtained when the nanocomposites are formed by thermal annealing. 3.4. Powders milled at 315 K and annealed at 503 K We now study the evolution of the powders ballmilled near room temperature (315 K) and annealed at the highest milling temperature used in the present work (503 K). It is expected that the solid solution stabilized by ball milling will decompose upon isothermal annealing, thus offering another route for the synthesis of nanocomposites. Decomposition does indeed take place during annealing at 503 K, but two different decomposition paths are observed as the annealing time increases. For short annealing times, (1 h in Fig. 7), the decomposition leads to cube-on-cube

decomposition texture as the Ag-rich and Cu-rich {111} reflections are aligned (Fig. 7c). The average lattice parameters are measured to be 0.406 ± 0.004 nm and 0.365 ± 0.004 nm, corresponding to 93 at.% Ag and 94 at.% Cu respectively [18]. This cube-on-cube decomposition texture should not be confused with the deformation texture that developed during room temperature ball milling. This latter texture does not present any major change in BF or DF images (Fig. 7a,b) after this short annealing. It is therefore concluded that the decomposition texture develops within the deformation texture. For longer annealing times (10 h in Fig. 8a,b), the microstructure of the sample has completely transformed: nanograins are clearly visible in BF or DF images, with an estimated average grain size of 40 nm. The compositions are determined to be almost pure Cu and pure Ag. These nanograins are essentially defect-free, except for the frequent presence of nanotwins: an EDS analysis performed over 47 grains containing nanotwins revealed that the twins exist in both Curich (19 grains) and Ag-rich grains (28 grains). The ring patterns observed in SAD (Fig. 8c) indicate that these grains are randomly oriented. The twins

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Fig. 5. TEM analysis of Cu50Ag50 powders milled at 503 K: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern.

analyzed above are typical of annealing twins observed after recrytallization of fcc metals with low stacking fault energies [19], 16 mJ/m2 and 45 mJ/m2 in pure Ag and pure Cu, respectively [20]. At intermediate annealing times (5 h in Fig. 9), bands of textured regions coexist with bands of well-defined nanocrystals (Fig. 9a,b). In diffraction mode features of the two decomposition morphologies are identified: two proportional sets of fcc rings, modulated by the deformation texture, coexist with randomly distributed sharp spot reflections (Fig. 9c). Overall, the new microstructure obtained after static annealing at 503 K is thus

typical of cold-worked materials that have undergone discontinuous static recrystallization [19]. The STEM chemical analysis (Fig. 10a,b) offers a remarkable view of the coexistence of the two modes of decomposition. Note also that, in the unrecrystallized regions, no decomposition is detected by STEM/EDS, whereas diffraction unambiguously reveals that the very same regions are decomposed. We thus conclude that the scale of this decomposition is smaller than the thickness of the foil (10–15 nm). It is interesting to contrast the nanocomposite obtained after annealing at 503 K for 10 h (Fig. 10d) with the one directly obtained

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Fig. 6. STEM images of Cu50Ag50 powders milled at 503 K: (a) ADF, (b) Cu (red) and Ag (green) EDS map, (c) bright field image, (d) 3 nm wide Cu compositional profile along the line shown in white in (b). The average composition of this area is 49.8% Cu.

by elevated temperature ball-milling (Fig. 6b). First, the scale of the thermal nanocomposite, about 75 nm, is more than twice that of the ballmilled nanocomposite, which directly results from the two-fold factor in the grain sizes. Second, the Ag-rich phase still forms a clearly connected matrix, but the Cu-rich regions now take more compact shapes. 3.5. Nanohardness of ball milled or ball milled and annealed Cu50Ag50 powders The measured nanohardness values are reported in Table 1. The lowest hardness value on ballmilled powders, ⬇4.2 GPa, is obtained after milling at liquid nitrogen temperature. Upon increasing the milling temperature, the hardness increases continuously, up to ⬇5.9 GPa at 503 K. Upon annealing the powders milled at 315 K and annealed at 503 K, the hardness also increases with annealing time, reaches a maximum of ⬇6.1 GPa for the 5 h anneal, but then starts to decrease. Using the microstructural characterization of the powders presented in the previous paragraphs, we

conclude that both phase decomposition and recrystallization contribute to the increase of hardness. This last point is unusual and will be discussed below.

4. Discussion Our results on Cu50Ag50 dynamical phase stability as a function of the milling temperature agree well with the main results previously obtained by Klassen and coworkers using XRD and DSC [5], and by us using APFIM [12]: room temperature milling results in the stabilization of a solid solution, whereas milling at elevated temperatures, above 393 K, results in phase decomposition. The statistical analysis of the atom probe results indicates that the solid solution stabilized by cryo-milling is almost random, while the 315 K milling produces a less random solid solution [12]. For the largest bin size used, containing 2000 atoms, the standard deviation of the composition distributions obtained by APFIM reaches 4, 7 and 12% for milling temperatures of 85 K (cryo-milling), 315 K and

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Fig. 7. TEM analysis of Cu50Ag50 powders milled at 315 K and annealed at 503 for 1 h: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern.

423 K, respectively (in the last case the alloy is decomposed). To compare our present STEM chemical analysis with these results, we have recorded the composition of the samples discussed in this study using a spot STEM/EDS analysis. Fig. 11 shows the composition histograms obtained for three milling temperatures, 85, 315 and 393 K. The standard deviation increases from 3.3 to 4.8 to 8.6% as the milling temperature increases. A simple estimation of the probed volume is obtained by assuming a probe diameter of 2 nm and a sample thickness of 15 nm. This yields a probed volume

of about 3000 atoms. Even though statistical analysis of the APFIM results is only available for bin sizes up to 2000 atoms, the agreement between the two methods is quite satisfactory. Note also that in the decomposed state (Fig. 11c), the shape of the histogram is significantly different from those obtained from solid solutions (Fig. 11a,b). Klassen and coworkers [5] concluded from the analysis of their XRD results that three solid solutions co-exist at intermediate milling temperatures, e.g., at 393 K or at 423 K. We could not confirm this point, due to the large diffuseness of the diffracted intensity in

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Fig. 8. TEM analysis of Cu50Ag50 powders milled at 315 K and annealed at 503 for 10 h: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern.

the electron diffraction patterns. We are currently trying to address this question by using APFIM. The main body of new results of the present TEM/STEM study deals with the microstructural state of the powders. Indeed, at high milling temperatures (above 453 K), the microstructure observed is typical of recrystallized materials, whereas at low milling temperatures, below 393 K, textured microstructures are obtained (see also Part I). The sequence of steady-states for increasing milling temperatures is similar to the temporal sequence we observe during the isothermal annealing of textured Cu50Ag50 solid solutions: at intermediate milling temperature, or after short

annealing times, decomposition takes place while texture is retained, but at elevated milling temperatures, or at longer annealing times, the material recrystallizes. We will now discuss the interplay between decomposition and recrystallization, first during isothermal annealing, and then during ball milling at elevated temperature. During the annealing of a deformed supersaturated solid solution, both recrystallization and decomposition can take place, and the two processes can affect each other. Using a simple analysis based on the driving forces and the mobilities available for each transformation, Hornbogen and Ko¨ ster [15] predict that, at low annealing tempera-

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Fig. 9. TEM analysis of Cu50Ag50 powders milled at 315 K and annealed at 503 for 5 h: (a) bright field image, (b) dark field image, and (c) selected area diffraction pattern.

tures, decomposition should precede recrystallization. This is indeed in agreement with the sequence reported in section 3.4. Note that for the system studied here the driving force for decomposition is much larger than that for recrystallization: the excess chemical energy stored in a Cu50Ag50 solid solution is around 6 kJ/g.atom, whereas the maximum energy stored in deformed materials is usually estimated around 0.2 kJ/g.atom [19]. The decomposition that we observe at short annealing time takes place at such a small scale that it cannot be measured in STEM/EDS maps. While some dislocations may annihilate during this decomposition, the dislocation density remains probably

high because interfacial dislocations are required to accommodate the ⬇13% lattice mismatch between the decomposed phases. Therefore this decomposed material may still recrystallize. Indeed, at intermediate annealing times (Fig. 9 and Fig. 10a,b), a discontinuous static recrystallization takes place, producing near equiaxed, defect free grains with a typical size around 40 nm. It is also observed that each new grain is either very Agrich or very Cu-rich. The size of these recrystallized grains is, in fact, likely to be limited by chemical decomposition since long range transport of chemical species is required for the grain size to increase. This recrystallization process can be

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Fig. 10. STEM chemical characterization of Cu50Ag50 powders milled at 315 K and annealed at 503 K for 5 h: (a) ADF image and (b) Cu (red) and Ag (green) EDS map; and annealed at 503 K for 10 h: (c) ADF image and (d) Cu (red) and Ag (green) EDS map. Table 1 Nanohardness values measured on ball milled samples at various milling temperatures, or ball milled at 315 K and annealed for various times at 503 K. LN2 refers to the cryo-milled powders studied in Part I [14] Milling temperature or annealing time LN2 315 K 393 K 453 K 503 K 10 min 1h 5h 10 h

Average (GPa) 4.23 5.32 5.30 5.82 5.87 5.68 5.96 6.11 5.64

Number of measurements 22 7 14 13 10 6 8 14 14

seen as a case of recrystallization of duplex alloys, yet at the nanometer scale. Another point in the discussion on the static recrystallization of ball milled powders is the prediction by Hornbogen and Ko¨ ster [15] that decomposition should delay recrystallization. In our case, the samples have decomposed and have roughly undergone 50% recrystallization after 5 h at 503 K (Fig. 9). The recrystallization temperature

Minimum (GPa) Maximum (GPa) Standard deviation (GPa)

3.45 4.91 4.88 5.31 5.31 5.21 5.53 5.54 5.07

4.83 5.72 5.86 6.43 6.77 6.02. 6.77 6.85 6.26

0.35 0.26 0.24 0.34 0.56 0.27 0.40 0.44 0.36

of non-decomposed Cu50Ag50 solid solutions is not available, but we can attempt to extrapolate it from data on pure elements and dilute solutions. The 50% recrystallization temperature in 98% coldrolled 99.999% pure Cu is around 360 K [21]. It is however known that the addition of 0.01% of Ag can raise the recrystallization temperature by 80 K [22], but this effect saturates when the Ag content increases, leading to maximum increase of

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Fig. 11. Histograms of chemical analysis measurements by STEM/EDS spot analysis performed on powders ball milled at 85 (see Part I), 315, and 393 K. The number of independent measurements is 113, 112 and 201, respectively. The standard deviations of the composition distributions are 3.3, 4.8, and 8.6%, respectively.

150 K for 60% cold-drawn samples [23]. Pure silver, on the other hand, when heavily cold-rolled may recrystallize near room temperature [24], but copper additions can raise this temperature by 200 K [25]. As solubilities larger than 5% are measured in Cu50Ag50 powders after decomposition (see section 3.4), it appears that the 50% recrystallization temperature found for decomposed Cu50Ag50 pow-

ders can well be rationalized by the enhanced solubilities of the phases. From the present results, it is therefore not possible to conclude that decomposition has delayed recrystallization. We now turn to the evolution of steady states in the ball-milled powders as the milling temperature is increased. Similarly to the static case, decomposition is observed to take place first. In the dynamic case, however, decomposition begins at lower temperatures (393 K), most likely because of the supersaturation of point defects produced by nonconservative motion of dislocations [5,9]. This low-temperature decomposition takes place within the deformation texture produced by ball milling. As the milling temperature increases to 503 K, the textured microstructure is replaced by a distribution of small (⬇20 nm) equiaxed grains, which are mostly defect free. This transformation can result from dynamic recovery or dynamic recrystallization. Since a mixture of transformed and untransformed regions is observed for intermediate milling temperatures (453 K, not shown here), it is proposed that discontinuous dynamic recrystallization is the main operating mechanism. This suggestion is supported by the fact that both copper and silver have low stacking fault energies, and thus are not likely to undergo extensive recovery [19]. In fact evidence for dynamic recrystallization in pure silver has been obtained at temperatures as low as 373 K [26]. Several differences are noticeable between the nanocomposite obtained by milling at 503 K and the one obtained by static recrystallization of powders ball milled near room temperature. First the grain size and the scale of decomposition are about two times smaller in the nanocomposite obtained by 503 K milling. Second, in this nanocomposite the shape of the Cu-rich and Ag-rich domains are elongated and ramified (Fig. 6), as opposed to the quite compact shapes observed after static annealing (Fig. 10b,d). This ramified morphology is consistent with the continuous deformation that the material undergoes during milling. A characteristic that is apparently common to both nanocomposites is that the Ag-rich phase is clearly connected, whereas the Cu-rich phase has a tendency to form isolated precipitates, even though the volume fractions of the phases are

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almost identical (see section 3.3). Two main factors could explain the formation of isolated Cu-rich precipitates in a Ag-rich matrix. First, since copper has higher elastic constants, the modulus effect leads to an elastic repulsion between copper precipitates, but to an attractive interaction between silver precipitates. These elastic interactions will influence atomic diffusion and thus the decomposition morphology. In fact two-phase morphologies very similar to the ones observed in this study after annealing at 503 K (Fig. 10c,d) have been reported by Finel [27], who used atomistic computer simulations to follow the decomposition of a model alloy close to Cu-Ag. Second, since Cu has a higher yield stress than Ag, one expects that the more deformable Ag-rich phase coats the more deformation-resistant Cu-rich phase. This factor, of course, would only contribute to the stabilization of the morphology of the nanocomposite directly obtained by milling (Fig. 6). It is interesting to compare and to extend the present results to other binary systems, as well as to other forcing conditions. While Cu-Fe [13] and Cu-Co [28] nanocomposites have been obtained by thermal annealing of ball milled solid solutions, we anticipate that it should be possible to synthesize directly nanocomposites in these systems by elevated temperature ball milling. Preliminary results indicate that even for alloys with large positive heats of mixing, e.g., Ni-Ag, it is still possible to obtain nanocomposites by ball milling [29], even though full solid solutions cannot be obtained by low temperature milling of this alloy system [30]. We also want to stress that the present strategy for synthesizing nanocomposites, which relies on a dynamical balance between thermal decomposition and some forced mixing introduced by the external processing, can be extended to other nonequilibrium processing routes [31]: we have recently shown that Cu-Ag nanocomposites can be synthesized by processing Cu/Ag multilayers with energetic ions when the irradiation temperature is properly adjusted to achieve a similar dynamical balance between thermal decomposition and radiation-induced chemical mixing [32]. We finally discuss the nanoindentation results. From our microstructural analysis and from the results summarized in Table 1, we can conclude

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that both phase separation and recrystallization lead to the strengthening of the powders. The first point, the effect of phase separation, is not surprising since it is well known that ageing of quenched concentrated Ag-Cu alloys leads to large hardness increase [33]. The second point is more surprising since recrystallization is commonly associated with softening. We have established, however, that the grain size after static or dynamic recrystallization is much smaller than the initial grain size of the textured and cold-worked powders. It is proposed that the formation of a large number of high-angle heterophase boundaries and annealing twins during recrystallization result in a hardness increase that exceeds the softening due to dislocation annihilation. We finally note that if we calculate the yield stress of these powders from the measured hardness by applying the conventional 1/3 factor, the resulting values are near the theoretical shear strengths of Cu and Ag, 3.6 and 2.3 GPa, respectively [20].

5. Conclusion We have shown that, starting from solid solution powder precursors, Cu50Ag50 nanocomposites can be synthesized by milling at elevated temperature or by thermal annealing. In both cases, decomposition of the initial solid solution precedes recrystallization of the heavily deformed powder precursor. At low enough milling temperature (393 K), or at short enough annealing time (1 h at 503 K), decomposition takes place at a scale of 10 nm or less, within the deformation texture produced by room temperature milling. At higher milling temperature (503 K), or for longer annealing times (10 h at 503 K), dynamic and static recrystallization takes place, respectively. The resulting nanocomposites present remarkably homogeneous nanostructures, with a decomposition scale ranging from 30 nm after dynamic recrystallization to 75 nm after static recrystallization, in direct relationship with the small grain size achieved, 20 and 40 nm, respectively. These nanocomposites exhibit very high hardness values, close to 6 GPa. Both decomposition and recrystallization are observed to result in hardening of the materials. The unusual

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recrystallization-induced hardening results from the very small grain sizes produced by dynamic or static recrystallization.

Acknowledgements Helpful discussions with Professors R. S. Averback, J.-P. Chevalier, U. Herr, and Drs. R. Enrique, M. Hy¨ tch, T. Klassen, are gratefully acknowledged. This research was supported by the US Department of Energy, Division of Materials Sciences, under the Award No. DEFG0291ER45439, through the Materials Research Laboratory at the University of Illinois at UrbanaChampaign, and by the National Science Foundation, Grant No. DMR-9733582.

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