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A Cu–Co composite material is chosen as a model system to study structural evolution and phase formations during severe plastic deformation. The evolving microstructures as a function of the applied strain were characterized at the micro-, nano-, and atomic scale-levels by combining scanning electron microscopy and transmission electron microscopy including energy-filtered transmission electron microscopy and electron energy-loss spectroscopy. The amount of intermixing between the two phases at different strains was examined at the atomic scale using atom probe tomography as complimentary method.

It is shown that Co particles are dissolved in the Cu matrix during severe plastic deformation to a remarkable extent and their size, number, and volume fraction were quantitatively determined during the deformation process. From the results, it can be concluded that supersaturated solid solutions up to 26 at.% Co in a fcc Cu–26 at.% Co alloy are obtained during deformation. However, the distribution of Co was found to be inhomogeneous even at the highest degree of investigated strain.

1. Introduction Methods of severe plastic deformation, such as high-pressure torsion (HPT), are very powerful tools for introducing significant grain refinement in a wide range of initially coarse grained single phase materials and alloys. Deformation of multiphase materials by HPT can lead to nanocomposites with an even smaller grain size in the order of  10 nm. In contrast to single phase materials and alloys, two or three phase materials deformed by HPT, whose microstructural evolution might be considerably different and dependent on its individual constituents, have not been intensively investigated. Metastable solid solutions can form during deformations if the multiphase material consists of components exhibiting no equilibrium solid solubility.

During ball milling, the phenomena of metastable and stable phase formation have been intensively investigated and explained by an interdiffusion reaction of the components for systems with a negative heat of mixing. For alloys with positive heat of mixing, a diffusion reaction generally results in decomposition. Although often reported in recent years for different alloy systems with positive heat of mixing (Cu–Cr, Cu–W, Cu–Fe) deformed by HPT, the formation of metastable solid solutions is far from being well understood for ball milling and HPT. In a recent paper, we have reported that the formation of supersaturated solid solutions during HPT of a Cu–Cr composite material, which is a combination of a soft and ductile matrix containing a hard and brittle second phase, is mainly controlled by a mechanical mixing mechanism and the amount of mixing is controlled by the strain level. Mainly based on energy-dispersive synchrotron diffraction measurements, in which the domain size, the microstrain and the change of the lattice parameter of both phases as a function of the applied strain was evaluated, detailed investigation of the microstructural evolution leading to the strain induced mixing during the HPT deformation was lacking. In this study, a similar type of composite consisting of a soft, ductile Cu matrix containing brittle Co particles is processed by HPT to understand the process of microstructural evolution during HPT deformation in detail.

The binary Cu–Co system, similar to the Cu–Cr system, possesses a large positive heat of mixing, has two different crystal structures at room temperature, the fcc Cu and hcp Co structure, and the formation of non-equilibrium solid solutions have already been observed in this alloy system during ball milling. The deformation strain can be directly calculated and linked to the microstructural observations.

Microstructural characterization was conducted on all length scales. First of all, extensive scanning electron microscopy (SEM) investigations including image analysis to quantitatively determine the size, numbers, and volume fractions of the Co particles and the microstructural evolution of the Cu matrix at different strains have been performed. Furthermore, transmission electron microscopy (TEM) investigations have been conducted to characterize the as-deformed microstructures. Analytical TEM investigations included energy-filtered transmission electron microscopy (EFTEM) and electron energy-loss spectroscopy (EELS). With EFTEM imaging, Co and Cu compositional maps were obtained. Quantification of EELS spectrum images in both phases and across Co and Cu phase boundaries yields information on the chemistry and amount of intermixing between Cu and Co phases with high spatial resolution. Atom probe tomography (APT) investigations were used to precisely determine the amount of intermixing at different levels of strain.

APT is a suitable technique for our case here since it offers uniquely high spatial resolution in three-dimensions for identifying local composition distribution and interfacial mixing at the atomic scale. Additionally to the microstructural characterization described above, the evolution of the microstructure was examined by hardness measurements. By using the example of a simple Cu–Co model composite material, this work has two major aims: describing the microstructural evolution leading to grain refinement of the composite material and dissolution of the Co particles in detail and to confirm the formation of supersaturated solid solutions in this system during HPT. Initial material and HPT processing The investigated material is a two-phase Cu–Co material containing about 26 at.% Co fabricated by RHP-Technology (Seibersdorf, Austria).

Co powders (purity 99.8%) and Cu powders (purity 99.9%) were mixed and subsequently precompacted into samples with cylindrical shape. For comparison purposes, pure Cu powder samples (with the same purity) were also compacted. The hardness of the initial, compacted material was measured by Vickers microhardness using a load of 500 g (H V0.5). The individual hardness of both phases of the Cu–Co material were measured with a nanoindenter (ASMEC UNAT-SEM2) fitted with a Berkovich Indenter. Vickers hardness is calculated using the measured indentation hardness (20 indentations in each individual phase using a load of 2 mN). Conversion into conventional Vickers hardness is done using the InspectorX software package, whereby the difference between indentation and Vickers hardness is below 10%. Processing of disk shaped samples by HPT deformation was conducted at room temperature using an HPT facility of the type described in earlier work.

The HPT processing was conducted under a pressure of 5 GPa with a rotation speed of 0.2 rpm. The disk samples (diameter 8.0 mm, thickness  0.6 mm) were processed for five different numbers of revolutions (N = 1, 5, 7, 25 and 100 turns).

All strains quoted in this study are given as von Mises equivalent shear strain ε eq. It is calculated according to. Hardness measurements Vickers microhardness measurements (H V0.5) were performed on deformed samples (N = 1, 5, 25 and 100) as function of the equivalent strain ε eq. Indents were made across the radii of the disk with a distance of 0.25 mm between the measurement positions and mean values of three individual indents at each position were determined. At selected positions, nanoindentation experiments were conducted with a cube corner diamond tip on a Hysitron TriboIndenter® with the Performech™ controller in displacement controlled mode with constant loading, hold and un-loading time (maximum displacement set to 50 nm).

Several individual indents were made following a quadratic square (10 × 10 or 12 × 12 indents) with 0.7 μm distance spacing between each separate indent. All of the measured nanoindentation values were then used to construct color-coded contour maps at the selected positions over a total area of 49.0 μm 2 or 70.6 μm 2, respectively. The principal goal of the nanoindentation measurements was to provide a qualitative and visual presentation of the hardness distributions at different deformation strains at a smaller scale and not to compare the measured indentation hardness to Vickers microhardness.

Microstructural characterization Microstructural examination was carried out by SEM using a Zeiss SIGMA™-VP field emission scanning electron microscope device equipped with an energy dispersive spectroscopy (EDS) detector. The microanalysis data was evaluated using the AZtec software (Oxford Instruments). All SEM micrographs were taken in tangential direction at selected positions (i.e. Different equivalent strain ε eq) across the radii of the HPT deformed samples.

Grain sizes were determined from back-scattered electron SEM micrographs with a high magnification at selected positions with the linear intercept method using the software Lince 2.4.2e. To study the process of fragmentation and microstructural evolution of the Co phase as a function of the equivalent strain ε eq, the software program ImageJ was employed. At each investigated position (i.e. Equivalent strain ε eq), five separate back-scattered electron SEM micrographs are recorded and investigated (see for the exact positions). The software automatically determines different phase parameters (such as a size, aspect ratio and area fraction) from the SEM images.

The total number of Co particles per area (μm 2) evaluated at each investigated position ranged from over 0.17 particles per μm 2 at low strains to 0.008 particles per μm 2 at a ε eq of 446. Overview of the total number of Co particles, the number of Co particles with a particle length  1 μm and the number of Co particles with a particle length of.

Hardness and microstructural evolution as function of the applied strain In (a) and (b), the initial microstructure of the Co and Cu phases in the Cu–Co material is shown. The microhardness of the undeformed, as-fabricated Cu–Co sample is 137 ± 39 HV. The grain size in the Cu phase is 486 ± 110 nm, and the grain size of the Co phase is significantly smaller (about 100 nm) with only single grains having a grain size above the nanometer range. XRD investigation of the initial, un-deformed material confirms its two-phase nature, consisting of two distinct hcp Co and fcc Cu phases ((a)). In the SEM image shown in (c), the phase distribution of hcp Co and fcc Cu in the initial condition is illustrated with a full-color overlay showing the variations in the EDS spectrum. Regions appearing green in the micrograph represent the Cu phase, regions appearing blue consist of the Co phase. It is apparent that the Co phase is inhomogeneously distributed and clustered.

The individual hardness of the Cu phase is 103 ± 12 HV in the initial condition. Due to its nanocrystalline grain size, the hardness of the Co phase is very high (394 ± 21 HV). Hence, a very hard second phase (Co) is embedded in a relatively soft matrix phase (Cu), similar as in a classical composite material. In the following, the Cu–Co material is treated as if it is such a composite material consisting of a Cu matrix containing hard Co particles. (a) XRD patterns recorded in the initial condition and after HPT deformation for 100 rotations. (b) Evolution of microhardness plotted as a function of the equivalent strain ε eq for Cu–Co samples deformed for N = 1, 5.

Processing the Cu–Co sample by HPT leads to strong grain refinement in the Cu phase and an increase in the overall hardness of the deformed material. The microstructure of the sample is homogeneously refined in the steady state region. This is confirmed by SEM investigations recorded at various positions in this region, which correspond to a saturation of the microstructural refinement and in which application of further HPT deformation will not induce further grain refinement. As an example, the microstructure at equivalent strains ε eq of 446 are shown in (d). Compared to the initial state, an ultrafine grained structure with a mean grain size of 101 ± 20 nm is visible. The microstructure in the steady state is additionally investigated by TEM. (a) shows a bright-field TEM micrograph, which was recorded at an equivalent strain of ε eq = 446.

It illustrates an ultrafine grained microstructure with well-defined grains, the majority having a size around 100 nm. Most of the grains are elongated in the shearing direction, but some smaller, truly nanocrystalline, equiaxed grains are also visible. Contrast variations inside the grains indicate a huge amount of defects in the as-deformed structure. In the SAD pattern shown in (b), only fcc Cu Debye–Scherrer rings of ( 111), ( 200), ( 220), ( 311) and ( 222) are observed.

The SAD pattern is consistent with observations from XRD measurements on the as-deformed Cu–Co material. Contrary to the initial state, only fcc Cu peaks, which are slightly broadened and shifted to higher angles, are detected in the XRD pattern (a).

TEM bright field image showing the microstructure of the Cu–Co material at an equivalent strain of 446 (a) and corresponding SAD pattern (b). Only Deybe–Scherrer rings of fcc Cu are indicated in the SAD pattern. A plot of the microhardness as a function of the applied strain illustrates the hardness evolution with increasing deformation strain in detail ((b)).

The hardness increases almost instantly from about 140 HV to about 175 HV in the beginning of the HPT deformation. A further slight hardness increase to  210 HV is visible up to an equivalent strain of 10. Between strains of 10–90, no significant change in the hardness is visible. Thereafter, a strong hardness increase to 270 HV, in a strain range from 90 to 150, is measured.

At higher strains, the microhardness remains nearly constant and a “second” steady state of the microhardness is reached. In, two-dimensional color-coded contour maps, illustrating the indentation hardness distributions for strains of (a) ε eq  0, (b) ε eq = 28, (c) ε eq = 56, (d) ε eq = 98, (e) ε eq = 153 and (f) ε eq = 446, are shown. The indentation hardness values (GPa) are presented by unique colors given at the right in the figure.

It is apparent that high indentation hardness differences are recorded at a strain of  0, which illustrates the distribution of the Co particles in the Cu matrix ((a)). The hardness ratio between the highest and the lowest measured indentation hardness is 2.9 at a strain of  0. At low deformation strain, regions with lower indentation hardness decrease in size. At a strain of 56, only very small regions exhibit an indentation hardness lower than 2 GPa ((c)). Additional deformation to a strain of 98 reduces the indentation hardness differences even more ((d)). For strains  150, a rather homogeneous high indentation hardness throughout the total investigated areas is visible ((e) and (f)). Two-dimensional color-coded contour maps of the hardness distributions at different positions (i.e.

Equivalent strain ε eq) of the HPT disk: (a) ε eq  0, (b) ε eq = 28, (c) ε eq = 56. In, SEM micrographs with a lower magnification are shown to illustrate the microstructural evolution of the Co particles in the Cu matrix with increasing deformation.

The shearing direction lies along the horizontal direction of all micrographs, which is indicated in (a). Co and Cu phases can be easily differentiated by their brightness levels — Cu regions are brighter, Co regions appear darker — in the back-scatter electron micrographs. At strains of 0 to 153 ((a)–(e)), a rather inhomogeneous microstructure is observed. Some large, but also very small Co particles are visible in the Cu matrix. The particles, which seem to exhibit a slight elongation in the shearing direction, are randomly distributed in the Cu matrix. With increasing strain level, their overall size is refined and their density continuously decreases. In the micrograph recorded at an equivalent strain of ε eq = 446, the microstructure appears to be free of micro-particles ((f)).

SEM analysis with higher magnification carried out at this position reveals that even at this amount of applied strain, single Co particles can be still detected in the micrographs. The process of fragmentation and microstructural evolution of the Co particles is quantitatively investigated in detail as a function of the equivalent strain ε eq with the image analysis software. In (a), the area fraction of the Co phase as a function of the equivalent strain ε eq is plotted.

At low strains (0 1 μm, designated as “large” particles in the following, and particles with a length of their major axis. (a) Area fraction of the Co phase as a function of the equivalent strain ε eq. (b) Particle size distribution at different strains (ε eq = 0, 44, 88, 102, 167, 209 and 446). (c) Length of the major axis of an ellipse fitted. To illustrate the microstructural evolution of the Cu matrix in detail, SEM micrographs showing predominantly the Cu phase were recorded at the same positions as in.

In, the microstructure of the Cu matrix deformed to different strains of (a)  0, (b) 28, (c) 56, (d) 98, (e) 153 and (f) 446 is shown. Even at a strain of ε eq  0, an ultrafine grained structure in Cu has already developed.

Compared to the initial microstructure of the Cu phase (shown in (a)), the microstructure is refined even at this small amount of applied strain. The Cu grains are furthermore slightly elongated in the shearing direction, which is along the horizontal direction of all micrographs in. Due to a very fine microstructure, which is visible between the large grains, a quantitative determination of the grain size from back-scattered electron micrographs has not been performed at medium strains. Qualitatively, no significant change in the microstructure is visible to a strain of 56 ((c)). The size of the large grains is decreasing to a minor extent, the area of the fine microstructure is simultaneously increasing. At strains higher than 98, significant further grain refinement occurs ((d)).

At an equivalent strain of ε eq = 153 and ε eq = 446, no difference in the observed microstructures is visible ((e)–(f)). A “second” steady state in the microstructural refinement with a significant smaller grain size compared to the first steady state is reached. The microstructural change is consistent with the hardness change. The qualitative impression of a reduced grain size is further confirmed by grain size measurements from back-scattered electron micrographs at these positions. At an equivalent strain of ε eq = 153 and ε eq = 446, a mean grain size of 102 ± 11 nm and 101 ± 20 nm is measured, respectively. Back-scattered electron micrographs taken at different positions (i.e. Equivalent strain ε eq) illustrating the structural evolution of the Cu matrix: (a) ε eq  0, (b) ε eq = 28, (c) ε eq = 56.

The microstructure at a strain of 98 is additionally investigated by TEM. (a) shows a bright-field TEM micrograph, which was recorded in a homogeneous Cu region in between the Co particles. It illustrates an ultrafine grained microstructure with well-defined grains, most of the grains are elongated in the shearing direction as well. The majority of the grains exhibit a size above 100 nm. Compared to the microstructure in the “steady state” at a strain of 446, the microstructure is slightly coarser. In the corresponding SAD pattern, only fcc Cu Debye–Scherrer rings of ( 111), ( 200), ( 220), ( 311) and ( 222) are observed. (b) shows a bright-field TEM micrograph recorded in a two-phase region.

The transition between the Co particle and the Cu matrix is schematically indicated by the dashed white line. In the Co phase, a nanocrystalline structure, with huge contrast variations inside the grains due to a huge amount of defects, is visible. The corresponding SAD pattern illustrates Debye–Scherrer rings of the fcc Cu phase as well as the hcp Co phase. Analytical transmission electron microscopy In (a) and (b), zero loss filtered images and corresponding core-loss energy filtered Co (c) and Cu mappings (d) are shown, which qualitatively illustrate the distribution of Cu and Co phases after HPT processing to a ε eq of 102. To avoid grain overlap, thin edge regions of the samples were imaged and analyzed.

In both images, Co particles embedded in a Cu matrix are apparent. Inside both particles, Cu-rich regions are visible. Furthermore, smaller Co fragments and small Co particles with size below 100 nm are apparent in the Cu matrix. In (a), a zero loss filtered image at the same deformation strain is shown, in which a “large” Co particle on the edge of the sample is imaged. In (b), an energy filtered Cu map with a higher magnification of the position indicated in (a) illustrates Cu and Co-rich regions in this area.

The ADF STEM image, which is illustrated in (c), is positioned in such a way that a Co grain next to two Cu grains with a higher magnification is visible. EELS analysis was carried out to determine the local chemical distribution of Cu and Co at this position, which is shown in the concentration profiles obtained along line 1 and line 2 across the interphase boundaries marked in the STEM image ((d) and (e)). For the concentration profile along lines 1 and 2, 95 and 100 measurements with 0.5 nm spacing were conducted, respectively. In both composition profiles, Cu and Co phases are visible and nearly no Cu can be detected in the Co phase on the left side of the line scan in both cases. Along line 1, a Cu phase containing 12–28 at.% Co is subsequently measured. The concentration gradient along the line is quite small.

Along line 2, a Cu phase in which 8–24 at.% Co is dissolved can be seen next to the Co phase. Although the scan is conducted at the edge region of the sample, the concentration gradient at the Co/Cu interface along line 2 seems to be affected by the overlapping of the Cu and Co grains. However, the data strongly suggests, that a supersaturated solid solution of Co in fcc Cu is formed during HPT processing, whereas nearly no Cu is dissolved in the Co phase. Atom probe tomography Based on TEM investigations, the formation of supersaturated solid solutions with a single phase fcc Cu structure is supposed. In the TEM-based analysis, the EELS signal is averaged over the thickness of the sample which is always a potential source of error.

Hence, further correlative investigations using APT have been conducted at this strain to confirm and to measure the chemical distribution at a finer scale for Co in the Cu matrix. For evaluating the amount of mechanical intermixing at a higher deformation degree, the microstructural evolution at a strain of 446 (ε eq446) was additionally investigated in detail by APT. Site specific APT specimens at a strain of 98 were prepared from an homogeneous Cu single phase region (ε eq98-I), from a Cu-rich and Co-rich region (ε eq98-II) and directly from a Co particle (ε eq98-III) (see ). Show 3D reconstructed sub-volumes from the homogeneous regions of the Cu–Co alloy samples deformed to both strains (ε eq98-I ((a)), ε eq98-II ((a)) and ε eq446 ((b)).

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In all 3D reconstructed sub-volumes, Cu atoms are displayed in green and Co atoms are illustrated in blue. In (b), a 2D composition map illustrating the distribution of Co in a larger reconstructed cross-section of ε eq98-II is shown. The Co concentration is given in at.% indicated according to color (0 = blue; 90 = red). The selected sub-volume displayed in (a) is taken from the Cu-rich region (small boxes inserted in (b)).

Considering the analyzed volumes illustrated in, the distribution of Co atoms in the Cu matrix seems to be fairly homogeneous and no spatial correlations in the distribution of the Co atoms are visually apparent at both positions and strains. The local compositions, determined in the displayed subvolumes in, is 23.9 ± 0.1 at.% Co and 76.0 ± 0.1 at.% Cu for (a) and 25.6 ± 0.1 at.% Co and 74.2 ± 0.1 at.% Cu for (b). The local composition of the subvolume displayed in (a) is 21.7 ± 0.1 at.% Co and 78.1 ± 0.1 at.% Cu. Hence, the agreement with the nominal composition of the alloy (74 at.% Cu, 26 at.% Co) is quite good. (a) 3D reconstructed sub-volume and 2D composition map of a Cu–Co alloy sample deformed to a strain of 98 (ε eq98-II) and (b) through-thickness representation (2D composition map) displaying the position of the sub-volume shown in (a). Additionally, the frequency distributions of Co concentration are presented in comparison to the theoretical binominal distribution for random solid solution to investigate the degree of homogenization of Co in the Cu matrix for each sub-volume.

Each dataset has the same size (40 × 10 × 50 nm 3) and is divided into equal blocks of 200 atoms. In (a) and (a), the experimental distribution is widened to both sides and does not fit the random distribution, indicating that the Cu–Co supersaturated solid solution is not homogeneous in both cases. The experimental distribution shown in (b) deviates less from random compared to those at a strain of 98, which reveals less Co clustering effects in the analyzed dataset. For a more quantitative interpretation of the frequency distributions, the Pearson coefficient μ is determined for each dataset, which is 0.88 for ε eq98-I ((a)), 0.87 for ε eq98-2 ((a)) and 0.67 for ε eq446 ((b)).

The calculated μ reflects the difference between the datasets obtained at a strain of 98 and 446. A Pearson coefficient close to 1 in samples ε eq98-I and ε eq98-II confirms the visible non-random Co distribution at that strain. In contrast, the much lower μ for sample ε eq446 indicates its more random Co distribution. The 2D composition maps, which are additionally displayed in (a) show the distribution of the Co concentration in the same region of interest illustrated in the 3D reconstructed sub-volumes.

The concentration is given in at.%, which is indicated according to color (15 = blue; 35 = red). As already expected from the experimental distributions, the composition map at a strain of 98 ((a)) show strong phase separation into Co-enriched areas. In between, the Co concentration is fairly high, but areas with very low Co concentration (. 4. Discussion An ultrafine grain structure in the Cu matrix quickly develops similar as in the pure metals in the very early stages of HPT deformation. A saturation of the hardness and refinement in the Cu matrix is reached at a strain of about 10. The grain size in this strain regime is significantly smaller compared to pure, bulk Cu deformed by HPT, but comparable to the steady state microstructure of pure HPT deformed powder Cu (microhardness of 207 ± 3.5 HV 0.5, grain size 168 ± 10 nm). The Cu–Co composite material consists of a softer, plastically deformable Cu phase with hard Co particles as second phase.

The grain size of the Co phase is on the average below 100 nm in the initial state. From the view of the deformation process, Cu and Co exhibit also different crystal structures (Cu with a fcc structure and Co with a hcp structure).

Hence, the imposed plastic strain is not distributed equally between the phases due to its different flow stresses and work hardening behavior resulting in complicated flow patterns and local stress concentrations. The Cu matrix accommodates the major part of the applied strain, which implies that the Cu matrix has to undergo additional deformation to accommodate the harder Co phase. Load transfer to the Co phase can be either accommodated by plastic flow or fracture. “Large” Co particles (length  1 μm) are slightly deformed in the beginning of the deformation indicated by their length increase ((c)). However, the length of the large Co particles continuously decreases at larger strains and the total number of particles per area exhibiting a size larger than 1 μm is continuously decreasing as well. The microstructural refinement progresses slowly since the Co particles are mainly fragmented and fractured. An example illustrating the fragmentation and deformation behavior of the Co particles in the Cu matrix at medium deformation strain is shown in (a).

The Co particles are steadily flattened and slightly elongated (like the large particle in the micrograph), fractured (marked in the micrograph) and rebound during continuing deformation similar as in ball milling or surface severe plastic deformation. The same observations of the structural evolution of the Co particles have been made in the EFTEM investigations. (2) the fractal dimensions D of N fragments with a size larger than r can be calculated. In Eq., C is a constant. At low and medium strains, the fragmentation process of large Co particles in the Cu matrix can be described by the fractal distribution as well.

Calculating the fractal dimensions from the plot shown in (b) results in similar fractal dimensions of D = 2.37, D = 2.51, D = 2.78, D = 3.28, D = 2.75 and D = 2.18 at strains of ε eq  0, ε eq = 44, ε eq = 88, ε eq = 102, ε eq = 167 and ε eq = 209, respectively. For a higher strain (ε eq = 446), the particle size distribution deviates from the fractal distribution model due to the low amount of detected particles and very small particle sizes. Particles with a very small size (.

5. Conclusion SEM and hardness measurements combined with state-of-the-art nanoscale characterization methods were conducted to clarify structural evolution and phase formations in immiscible composites deformed by HPT on all length scales using the example of a simple Cu–Co model material. The findings can be summarized as follows:. (i) Supersaturated solid solutions up to 26 at.% Co in a fcc Cu–26 at.% Co alloy were obtained by HPT, contrary to the Co phase, which is not affected by Cu. (ii) Although Co has an extended solubility in Cu, the formed supersaturated solid solutions are not homogeneous.

Small Co particles remain in the Cu matrix and Co is still enriched in nanometer-sized clusters even at the highest amount of investigated strain. (iii) The fragmentation process of large Co particles follows a fractal distribution at low and medium strains. (iv) Dissolution of small Co particles starts at low strains. Further grain size refinement due to alloying induces a higher hardness of the Cu matrix, which accelerates the fragmentation and fracturing of large Co particles and promotes the dissolution of small Co particles at medium deformation strains. (v) Shear induced mixing is suppressed by localized deformation of the ductile-brittle components and inhibition of dislocation transfer across fcc– hcp interfaces in this model system.

The authors gratefully acknowledge the financial support by the Austrian Science Fund (FWF): J3468-N20. The atom probe instrument was financed by the DFG (INST 256/298-1 FUGG) and the Federal State Government of Saarland. We thank the Erich Schmid Institute of Materials Science in Austria, especially R. Pippan and P. Kutlesa, for providing all HPT deformed samples and fruitful discussions, G. Felber for TEM sample preparation, J. Schmauch for experimental assistance at the TEM and C.

Soldera for specimen preparation by FIB.