Scholarly article on topic 'In-situ  tracking the structural and chemical evolution of nanostructured CuCr alloys'

In-situ tracking the structural and chemical evolution of nanostructured CuCr alloys Academic research paper on "Materials engineering"

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Acta Materialia
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{"CuCr nanocrystallines" / " in-situ TEM" / Spectroscopy / "Thermal stability" / Defects}

Abstract of research paper on Materials engineering, author of scientific article — Zaoli Zhang, Jinming Guo, Gerhard Dehm, Reinhard Pippan

Abstract We report the thermal stability of supersaturated CuCr nanocrystallines alloys at the atomic resolution using modern spherical aberration-corrected transmission electron microscopy (TEM) via performing in-situ structural and spectroscopy experiments. It is found that CuCr nanocrystallines are not only subjected to a structural change but also undergo a chemical evolution upon annealing. Chemical destabilization of supersaturated CuCr nanocrystallines occurs at a quite low temperature. Heating triggers a rapid separation of Cu and Cr grains at the forced intermixing zone, accompanied by an obvious decrease of average interface width whereas the grain growth is not significant. Elemental profiles and images recorded in real time reveal that the local compositions vary with heating, which in turn permits to derive the concentration of excess vacancy generated by deformation and observe its evolution with temperature, further to analyze the dynamic behavior in nanocrystalline materials. Electronic structure changes at the interface forced intermixing zone are revealed by the fine structure analysis. The study uncovers the interplay between the thermal stability and chemical decomposition process of bulk nanostructured materials in real-time.

Academic research paper on topic "In-situ tracking the structural and chemical evolution of nanostructured CuCr alloys"

Accepted Manuscript

In-situ tracking the structural and chemical evolution of nanostructured CuCr alloys Zaoli Zhang, Jinming Guo, Gerhard Dehm, Reinhard Pippan

PII: S1359-6454(17)30601-8

DOI: 10.1016/j.actamat.2017.07.039

Reference: AM 13938

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Received Date: Revised Date: Accepted Date:

Acta Materialia

24 February 2017 19 May 2017 17 July 2017

Please cite this article as: Z. Zhang, J. Guo, G. Dehm, R. Pippan, In-situ tracking the structural and chemical evolution of nanostructured CuCr alloys, Acta Materialia (2017), doi: 10.1016/ j.actamat.2017.07.039.

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In-situ tracking the structural and chemical evolution of nanostructured CuCr alloys

Zaoli Zhang1 *, Jinming Guo1 , Gerhard Dehm2, Reinhard Pippan1

1 Erich Schmid Institute of Materials Science, Austrian Academy of Science, Leoben, 8700, Austria 2 Max-Planck-Institut für Eisenforschung, D-40237 Düsseldorf, Germany

*To whom correspondence should be addressed. E-mail: zaoli.zhang@oeaw.ac.at.

Abstract

We report the thermal stability of supersaturated CuCr nanocrystallines alloys at the atomic resolution using modern spherical aberration-corrected transmission electron microscopy (TEM) via performing in-situ structural and spectroscopy experiments. It is found that CuCr nanocrystallines are not only subjected to a structural change but also undergo a chemical evolution upon annealing. Chemical destabilization of supersaturated CuCr nanocrystallines occurs at a quite low temperature. Annealing triggers a rapid separation of Cu and Cr grains at the forced intermixing zone, accompanied by an obvious decrease of average interface width whereas the grain growth is not significant. Elemental profiles and images recorded in real time reveal that the local compositions vary with annealing, which in turn permits to derive the concentration of excess vacancy generated by deformation and observe its evolution with temperature, further to analyze the dynamic behaviors in nanocrystalline materials. Electronic structure changes at the interface forced intermixing zone are revealed by the fine structure analysis. The study uncovers the interplay between the thermal stability and chemical decomposition process of bulk nanostructured materials in real-time.

Nanocrystalline (NC) materials which are of interest for both fundamental studies and engineering applications are recognized as substantially out of equilibrium. As a consequence they often rapidly coarsen at moderate temperature[1,2]. This coarsening tendency impedes the use of these materials at ambient and especially elevated temperatures [3,4]. The microstructure instability of NC materials is an important concern since it influences the resulting mechanical properties significantly. Especially, the large energy associated with the high volume fractions of GBs in NC materials represents a driving force for grain growth. Therefore, nano-grains in NC materials tend to coarsen. To address this important issue in NC materials, experimentally and theoretically attempts to find effective solutions to prevent coarsening have been made. Solute additions to NC metals can effectively retard the coarsening process but can't eliminate the tendency to coarsen[5]5. Theoretically, thermodynamic-based approaches have been invoked to discuss the stability of NC materials, which to some extent are successful for some alloy systems[6-13]. More recently, by considering different parameters such as composition, temperature, interaction parameters, enthalpies of segregation and mixing, formation of amorphous layers and two-phase nanocomposites, Murdoch and Schuh[7] proposed a nanostructure stability map with regions of stability, meta-stability and instability in binary alloys. It seems very attractive to estimate the possibility of high temperature stability from thermodynamic parameters by using the regular solution approximation.

Substantial experimental studies of the thermal stability NC materials at high temperature have been carried out on metals and nanocomposites using a wide range of techniques, i.e. transmission electron microscopy (TEM), Electron-Energy-Loss Spectroscopy (EELS), atom probe tomography (APT), XRD and micro-hardness[14,15]. However, the experimental results seem divergent. The stability of NC materials differs from case to case when processed by different approaches. The thermal stability of NC materials related to different type of severe plastic deformation (SPD) is strongly influenced by the defects and stored energy generated from extreme deformation[16,17]. Among these, the studies by in-situ TEM are very promising to reveal the structural changes of NC materials in real-time. Grain growth and phase transformations can be directly visualized[18-20]. The recent experimental observation of grain growth in some NC materials during deformation[21,22] also attracted new and increasing attention. The stability aspects of bulk nanostructured materials were reviewed [23,24], more recently by Andrievski [25], who summarized the advances on this topic.

Although numerous achievements have been accomplished, the stability of NC materials, particularly in severely plastically deformed (SPD) materials, is still in its infancy and requires further high -spatial resolution/atom-scale insights. To date, all the studies on the stability of NC materials hardly dealt with a high-spatial resolution or atomic-scale in-situ observation, and by tracking the chemistry and structural evolution simultaneously. Such investigations are definitely essential to demonstrate how the bulk NC materials evolve with temperature chemically and structurally, and to understand the mechanism of thermal stability at a high spatial resolution.

Unfortunately, it is experimentally quite challenging to study the chemical and structural variations of NC materials simultaneously because the grain size in NC materials is usually quite small. New opportunities enabled by modern spherical aberration (CS)-corrected TEM with a sufficiently small probe and atomic-resolution to address this issue have appeared. The possibilities offered by modern instruments brought new inputs and fundamental understanding on deformation induced phenomena[26-30]. It is anticipated that new insights into the stability of NC materials will be gained with the advent of spherical aberration (CS)-corrected TEM.

Here, CuCr nanostructured alloy was utilized for in-situ monitoring the structural evolution and chemical composition changes simultaneously. The motivation is based on two considerations. From practical application point of view, CuCr alloys are used in numerous aspects, i.e. railway contact wires30 and electrodes for spot welding [31], and exhibit many extraordinary properties [32]. From scientific point of view, CuCr is a model system for studying the microstructure-properties relationship and thermal stability of nanostructured materials [33-38]. Especially, CuCr exhibits a positive Gibbs free energy change AGc for forming a solid solution alloy in the equilibrium state and is therefore nearly immiscible in the solid state, but can form supersaturated solid solution after severely deformation[33-38]. Experimentally, CuCr is well suited for in-situ tracking chemical composition (distinct Cr-L23 and Cu-L23 absorption edges, facilitate quantifying the composition with high spatial resolution by EELS at elevated temperature) and structural variation (exhibiting a distinct intensity in Z (atomic number) contrast scanning TEM (STEM) imaging, the intensity is ~ Z 2 ).

Experimental section Materials

A coarse-grained Cu-Cr composite material (43wt % Cr, 57wt % Cu, and corresponding to 48 at % Cr and 52 at% Cu, respectively) was deformed by high pressure torsion (HPT) at room temperature. The initial microstructure of the composite consists of a Cu matrix with Cr particles (volume fraction of about 50%, mean diameter of about 50 ^m) produced by PLANSEE (Reutte, Austria), similar to those in Ref [38]. Disks with a diameter of ~8 mm and a thickness (t) of ~0.8 mm were HPT deformed for different numbers of turns n under a constant pressure of 6.25 GPa with a rotation speed of 0.2 rotations/minute. And compressed air cooling was applied during the whole HPT process. HPT disks which were deformed for 25 turns (equal strains = 400) were used for in-situ annealing in TEM at 212°C and 414°C after preparing TEM specimens from the edge of the disk, where the highest strain has occurred.

Methods

For comparison, ex-situ annealing experiments were performed at the same temperature as in in-situ experiments. The corresponding TEM samples were prepared for the imaging and spectroscopy analysis.

TEM specimens were prepared by wedge Tripod polishing, and followed by final ion milling with 5 to 10 minutes. A TEM/STEM JEOL2100F operated at 200 kV and equipped with an image-side CS-corrector and an image filter (Tridiem) was used. The alignment of the CS-corrector was performed using the CEOS software based on the aberration measurements deduced from Zemlin tableaus. Eventually, the aberrations are sufficiently small. All HRTEM images shown here were recorded on a 2k x 4k pixel CCD camera at a magnification of 1.5 M* using an acquisition time of 1.0 second and a negative CS, under which all atoms are imaged as bright dots with different bright contrasts depending on their atomic numbers.

STEM images shown in this paper were recorded using an annular STEM detector, with the detector inner angle /outer angle were set to around 54 mrad /144 mrad. Under these conditions, the STEM high-angle annular dark field (HAADF) image is nearly Z-contrast image.

The local composition variations with temperature were carefully measured by EELS using an A-size electron probe. STEM-EELS spectrum-images were acquired using a dispersion of 0.2 eV/ channel, a collection semi-angle of 10 mrad, and a convergence semi-angle of 7.5 mrad. The probe size under optimum conditions can reach 0.2 nm (or even smaller). Numerous EELS spectrum images were acquired during in-situ annealing. Drift compensation was applied when the spectrum images were acquired. The acquisition time for each spectrum is 0.5 s, and the total acquisition time is 15 s for one line-scan with a step interval of ~ 0.5 nm. For EELS spectrum images, the following process was performed. Firstly, multivariate statistical analysis (MSA) was applied to the raw data of all spectrumimages. Secondly, the background for each spectrum was removed using a power law function, and Hartree-Slater models were used to fit the L2 and L3 edges. The cross-sections were then subtracted within the signal window of 569.4 - 577.4eV for Cr-L3 and of 578 - 586 eV Cr-L2, respectively. To quantify the atomic ratio of Cr/Cu, the windows for background subtraction were set to 80 eV and 130 eV, and signal windows were set to 120 eV and 130eV for Cr and Cu respectively. All the spectra were analyzed under the same conditions, i.e. the same signal and background subtraction windows, in order to compare the change of Cr/Cu atomic ratio.

The initial microstructure was analyzed by TEM 30 days after HPT processing. In-situ heating in a TEM was carried out at three different temperatures, room temperature (25°C), 212°C for ~ 60 minutes (holding time), and 414°C for ~ 60 minutes (holding time) using a Gatan heating holder. In order to track the chemical composition evolution in real time, more than 30 interfaces and GBs were picked for extensive EELS spectrum image analysis. Some interfaces/GBs (about 20 interfaces) were then repeatedly measured at different annealing temperatures under an identical acquisition condition

so that the compositional changes together with the structural evolution of the nanostructured materials were tracked with annealing temperature and time.

For determining the grain sizes, about 30 Cu particles with a distinguishable boundary were statistically measured. As the grain shapes are rather random and irregular, Cu grain sizes were then determined based on the STEM images using Image J software by measuring grain area changes. The grain areas were then converted into the average grain size, D, (the diameter of a spherical grain, see the details in supplementary materials, Fig.2S). As the grain sizes are quite different over a large area, the error bar is a bit large, i.e. about 12 nm. The average grain size reflects the grain changes with the annealing. However, it should be noted that the Cu average grain sizes determined using HAADF image can be slightly overestimated because the diffraction contrast in HAADF image is not visible, which may lead to an overestimation of the actual grain size (a large Cu particle could comprise two or more grains since the grain boundaries are not clearly resolved in Z-contrast HAADF STEM image).

Apart from the structural evolution observations, in-situ chemical composition tracking was also performed on several individual GBs at different temperatures. Sufficient thin areas of the TEM specimens were selected for analytical purposes. These measurements were performed to bring new insights into the thermal stability of NC materials from a chemical viewpoint.

Results and discussions SEM and XRD measurements

Figure 1a shows a backscattered electron image of raw bulk material of 57 wt.%Cu - 43 wt.% Cr with Cu (bright) and Cr (dark) particles prior to deformation. As can be seen, Cr and Cu grain size are quite large, possessing an average size of about 50 ^m distributed homogeneously. Figure 1b shows XRD measurements on the raw-material (prior to deformation) and as-deformed CuCr nanocrystalline material (25 rotations, prior to annealing). The comparison clearly displays a difference, i.e. the peaks become broadening and reduced relative intensity ratios, i.e. (111)Cu and (110)Cr peaks, revealing the atom dissolution after severely deformation, i.e. forced chemical intermixing occurs. Such intermixing and dissolution amount are relevant to the numbers of rotation and initial composition ratios, and a strong relationship can thus be established.

HRTEM characterization

Figure 2a shows one high resolution TEM (HRTEM) image taken from the as-deformed sample (prior to annealing), resolving the atomic structure between Cr and Cu grains as well as the interface. A low magnification TEM image is inserted, in which numerous nano-grains are clearly visible. One Cr grain is surrounded by several Cu grains (one Cu grain with a twin), where the part of the interface

(labeled) with a somewhat distorted region (1~2 nm in width) is clearly visible, separating the two grains with an orientation relationship: Cr [111] //Cu [110] (Note that in the upper part there is another Cu grain along high-index zone). Except one small part of the coherent interface, however, the interface between Cr and Cu appears to be slightly blurred due to the presence of a chemically and/or structurally distorted region. One possible interpretation might be that the distorted zone is relevant to Cu-Cr intermixing generated by severe deformation. An overlapping of two grains is unlikely in this case as a Moiré pattern should become visible. The forced chemical intermixing zone ascribed to the diffusion of Cr (Cu) atoms into the Cu (Cr) lattice may emerge as a 'distorted region' between two grains. However, here, it is still hard to directly correlate such region with interface intermixing based merely on the phase contrast image. Combined with quantitative image analysis and theoretical simulations, the interpretation could be substantiated. Here, quantitative analysis by means of geometrical phase analysis was also carried out (Fig.1S(a)-(b),supplementary materials)) to gain extra information about the interface structure and reveal the strain distribution (strain tensor components exx) and rotation angle (rigid-body rotation axy). The strain states are quite different in the Cr and Cu.

To further characterize the atomic structure of CuCr nanocomposite, another representative HRTEM acquired from a different region of the sample shows in Figure 2b, where one Cr grain along [111] is fully surrounded by several Cu grains. The interfaces between Cu and Cr are either with Moiré patterns (as indicated by double arrows) or with 'coherent' regions (indicated by single arrows), where atomic planes running through the two neighboring grains. Again, different from Cr grain, Cu grains are with twins and obvious distorted areas. In sum, the interface atomic scale structures are of variety.

STEM image analysis

The Z-contrast STEM images recorded at 25°C (room temperature, prior to heating), 212°C and 414°C are shown in Figure 3a-3c, which expose the dynamic evolution process of grain morphology. The grains possess irregular shapes, and the grain size ranges from a few nanometers to tens of nanometers. Close examination of several grains (labelled by arrows) clearly reveals that the grain shape and contrast at different temperatures have experienced change. The main changes can be summarized as follows: i) Grain shape (as indicated by arrows) becomes less sharp or round as the Cu retracts due to the diffusion. It is very remarkable from 212°C to 414°C as compared to heating from 25°C to 212°C; ii) the contrast difference between Cr and Cu is enhanced; iii) Surprisingly, some grains have undergone remarkable changes appearing a 'pre-melting' phenomenon, and completely disappear upon 414°C (as labeled by arrows). It happened in a few seconds for certain grains at the thin region. Only some Cu grains were subjected to such 'pre-melting\ independent of their sizes. At the thick region, however, the contrast becomes more complicated owing to grain

overlapping. Actually, the intensity change in the Z contrast image between two neighboring Cr and Cu grains is related to the inter-diffusion, which causes the vanishing of the supersaturation phenomenon and a reduction of forced intermixing zones.

By roughly comparing the STEM images, no remarkable grain growth was observed. To evaluate the subtle change of grains with irregular shapes (please note that the morphology seems different from low-magnification image in Figure1a, due to the different sample position and imaging formation mechanism), the statistical measurements on numerous grains were performed. The average grain size, D, was used to describe the grain size change at different temperature. Figure 4 demonstrates the change of the average grain size (D) with the temperature. It clearly displays that D slightly increases from 25°C to 414 °C, grain growth process exists, but less significant. It should be mentioned that not all the grains grow, and some grains actually decrease in size (Fig.2S, supplementary materials).

Chemical analysis

Figure 5 shows the chemical composition variations from one designated grain, two composition profiles crossing one Cu grain were acquired at 25°C and 212°C (corresponding images inserted). In-situ spectroscopy tracking at the same interfaces reveals that: i) the examined Cu grain slightly shrinks; ii) the Cu and Cr concentration profiles across the interface are broader prior to annealing, with an interface width of about 3~4 nm, suggesting the presence of a strong disorder at the interface after severe plastic deformation; iii) most significantly, the concentration profiles across the interfaces become more abrupt at 212°C with a reduced interface width. The in-situ chemical profiles reflect the concentration variation with temperature, i.e. Cu in Cr lattice (or Cr in Cu lattice). Rising temperature causes a reduction of concentration, consequently, the interfacial intermixing zone decreases. And it is also an indication that rapid diffusion of Cu and Cr atoms across the interface took place. By comparing the two profiles, one may conclude that the super-saturation especially of Cu in Cr partially vanishes, the concentration (supersaturation concentration) Cu in Cr (or Cr in Cu) has obviously changed, and an obvious decrease of Cu concentration in Cr grain has occurred. However, both the STEM image contrast and grain size show no significant change when the temperature approaches 212°C. Chemically speaking, the stability of NC materials is already lost at 212°C. It means that the destabilization process starts at a lower temperature than usually expected, while grain sizes remain almost unaltered.

Figure 6 presents average chemical elemental profiles at 212°C and 414°C, respectively. It shows how the chemical composition distribution crossing the interface develops with temperature. Apparently, the composition profiles become even steeper at 414°C than at 212°C. This reflects that the interface becomes more chemically abrupt and a sharp interface forms (as illustrated by the atomic model, inserts), interface sharpening phenomenon is observed in such immiscible system. Correspondingly,

the contrast difference in STEM images (Fig.3a and Fig.3c) increases at the thin region. These features suggest that the decomposition process at the intermixing zone is nearly or fully completed.

The structural and chemical evolution from one single interface at three different temperatures in real-time was also successfully tracked (Fig.3S). It is found that in this case the structural and chemical fluctuations can occur both at the interface and in the grain interior. It signifies that at high temperature (414°C) the dynamic process becomes very pronounced, consequently, both the composition and the image contrast undergo significant changes. Taken together, the measured composition profiles across the interface (Fig.5 and Fig.6) at 25°C, 212°C and 414°C demonstrate clear differences. The profiles remarkably change upon annealing, developing a chemical abrupt Cu-Cr interface, associated with the diffusion. As the diffusion requires vacancies, and severe plastic deformation (SPD) is known to enhance the concentration of vacancies [27], it can be thus speculated that the numerous vacancies contribute to the interfacial de-mixing process.

Furthermore, to rule out the possibility of artifact introduced by in-situ annealing, ex-situ annealing experiments were also carefully performed at 414°C prior to fabrication of electron transparent TEM samples. The results are shown in Fig.4S (supplementary materials), in which both the concentration profile and corresponding image are displayed. To a large extent, the profile shows a similarity to Fig. 5 and Fig. 6. From the ex-situ concentration profile at 414°C, diffusion coefficients obtained via fitting are quite close to that from in-situ heating (Table1) at same temperature. The comparison verifies in-situ experimental observations, signifying that the artifact, i.e. surface diffusion, is not significant under current experimental conditions. It seems somewhat surprising because it is usually assumed that surface diffusions may occur and play an important role during in-situ experiment.

Dynamic process analysis

With the chemistry profiles, diffusion coefficients can be worked out by diffusion simulation analysis using a non-linear curve fit approach. Using a modified function (Equation 1), the measured composition profiles in Figure 5 and Figure 6 were carefully fitted. Diffusion coefficient (D), at different temperatures can be thus determined.

£ _ CCr+CCU I CCr~CCU crfX~X0 ________________________________(1)

2 2 2VÔt ................................^ '

Where the CCr and CCu are the initial concentrations in the respective Cr and Cu grain. By fitting, (Dt) and D is solved (D is instantaneous diffusion coefficient), as listed in Table 1 (referring to notes in supplementary materials). For a comparison, the results from ex-situ heating experiments performed at 25°C and 414 °C are incorporated.

Table 1 Diffusion coefficients from composition profiling experiments. The data in last two columns are from the ex-situ heating experiments.

25°C 212°C 414°C 25°C (ex-situ) 414°C (ex-situ)

VDt, (nm) 0.616 ± 0.095 0.647 ± 0.053 0.317 ± 0.023 0.473 ± 0.180 0.514 ± 0.030

Dt, (cm2) 0.38 x 10 "14 0.42x10 -14 0.10 x 10 -14 0.22 x 10 -14 0.26 x 10 -14

D, (cm2/s) (1.95 ± 0.3) x10 -21 (1.75 ± 0.3) x 10 -18 (0.84 ± 0.3 ) x 10 -18 (1.16 ± 0.13) x 10 -21 (1.50 ± 0.01) x 10 -18

D#1, (cm2/s) 7.8x10 -38 1.35x10 -23 6.7x10 -17

D#2, (cm2/s) 2.60 x10-24

D#\ (cm2/s) 4.0 x10 -44

D means instantaneous diffusion coefficients from this study. D#1 denotes Cu self-diffusion coefficients in polycrystalline[39]. Data for ex-situ heating experiments are referred to supplementary materials. D#2, D#3—Cu self-diffusion coefficients in polycrystalline and single crystals at 20°C [40] . Here, simplification is made (referring to notes in the Supplementary materials).

Diffusion coefficients interestingly shows : i) a slightly high diffusion coefficient at 212°C compared to 414°C; ii) diffusion coefficient is significantly higher in nanostructured materials than in bulk Cu coarsened polycrystalline [39] and nanostructured polycrystalline [40], i.e. at room temperature D = 2.6*10"24 cm2/s. However, it should be noted that diffusion coefficients via the composition profile represent solute diffusion in the bulk lattice across the specific interfaces, different from diffusions along the grain boundary, dislocations etc... Since the chemistry changes across the interface were locally considered in current situations, and grain growth is not significant, in this case, only bulk diffusion (or volume diffusion) by means of defects (mainly vacancies) is involved [41]. Grain boundary (GB) diffusion can be very pronounced in nanocrystalline materials [42], however, it was reported that it decrease several orders of magnitude in HPT due to the sample subjected to the high pressure [43]. Moreover, HPT produces a high dislocation density, mostly located at GBs, which are also acted as sinks for vacancy annihilation. This is different from the diffusion addressed here. It is clearly shown by Wang et al that interfacial diffusions in nanostructured materials prepared by SPD can be significant in nanostructured materials due to the excess free energy, i.e. GB diffusivity being ~ 2 orders of magnitude higher than twin boundary diffusivity[44,45], and also slightly higher than that along conventional GB at temperature <373K. This highlights that 'non-equilibrium' GBs (with higher excess energies and diffusivities [46]) can be ultra-fast paths. However, as compared to the line profile measurement crossing the Cu-Cr interface by a small electron probe, it obviously differs.

In contrast to the chemical intermixing during the deformation, the diffusion process upon annealing is a kind of de-mixing process (or decomposition), in particular at the initial stage of heating. The de-mixing process is an uphill diffusion, and Cr and Cu atoms diffuse back the respective lattice. The driving force originates primarily from the positive AG caused by the supersaturated solid solution.

On the other hand, the excess vacancies produced by severe plastic deformation[27], which greatly assists for the kinetic process, play an essential role in the diffusion process. With increasing temperature, the excess vacancy reduces, and its influence diminishes and even completely vanishes. Such dynamic process of NC materials with annealing can be schematically illustrated by atomic models (insets in Fig.4, only two stages displayed). It simplified the dynamic process occurred at the forced intermixing interface.

A plot of D versus T is shown in Figure 7a, where Cu equilibrium coefficient (Deq, represents Cu self-diffusion coefficient) is inserted. As seen, experimental measured D and Deq crossed at a certain temperature, i.e. at ~620K. The cross-point regime indicates a transition from the excess-vacancy-dependent to the equilibrium-vacancy-dependent. It clearly demonstrates vacancy effects.

Moreover, with the help of above diffusion coefficient, deformation-induced vacancy concentration at different temperature can be worked out. A cumulative diffusion coefficient is used to model the impact of non-equilibrium dynamic. Under the presence of non-equilibrium interstitials and vacancies, the diffusion coefficient (instantaneous diffusion coefficient, D ) for a certain species can be characterized by: [47]

D = De«* ---------(2)

Where Cv is instantaneous vacancy concentration; C$q (CjQ) is equilibrium vacancy (interstitial) concentration; C$x is excess vacancy concentration. Deq denotes the diffusion coefficient under equilibrium condition (here, Cu self-diffusion coefficients used). [39] If f = 1, it means the diffusion works via interstitials while f = 0, it implies that the diffusion works via vacancies. Evidently, it is difficult to tell the exact amount of excessive vacancies and interstitials generated during HPT. However, many previous studies have shown that significantly additional vacancies are produced during the deformation[41,43,48-51], and these excessive vacancies enhance diffusions [41]. Furthermore, the study on an immiscible CuFe nanocomposite obtained by HPT also reveals that a high level vacancy concentrations advance the diffusion coefficients of Cu and Fe [41]. On the other hand, substitutional Cu or Cr atoms shift to interstices in cubic crystal during HPT is less likely as compared to interstitial atoms. In consideration of these, one may reasonably speculate that diffusions in the deformed CuCr nanocrystals are via the vacancies.

The above equation is then simplified as follows:

~ex -V

D CV? + CVX

Deq cevq

(i.i9+0.03)ey

Equilibrium vacancy concentration for Cu is expressed by [52]: = ( 27.1~14.8) x e 2t . According to the equation (2) and (3), C$X values are derived as shown in Table 2. The excess

vacancies in the specimen primarily originate from the deformation process [49,53]. The table reveals that excess vacancy concentration generated by HPT deformation in Cu is about 5.0* 10"3, and decreases with increasing the temperature. Up to 414°C, the concentration of excess vacancy is quite low, much less than that of equilibrium vacancy. The magnitude of measured excess vacancy concentration after HPT deformation at room temperature, i.e. 5.0* 10-3 for as-deformed, basically in the range of 10"3~10"5 (referring to the review paper by Lotkov et al, [51]), is quite close to those previously measured using macroscopic techniques, such as differential scanning calorimetry [48], positron annihilation spectroscopy [54]. This, from another point of view, provides a clue that the assumption made on the diffusion mechanism (mainly via vacancy diffusion) in severely deformed substitutional Cu-Cr alloy is valid.

A relationship of excess vacancy with annealing temperature is plotted in Figure 7b. The relation of Cu equilibrium vacancy concentration with temperature is inserted for a comparison. The diagram shows that: 1) non-equilibrium excess vacancy concentration dramatically decreases with increasing the temperature while the equilibrium vacancy concentration increases; 2) at room temperature, the excess vacancy concentration in as-deformed materials is quite high, significantly affecting diffusion of Cu (Cr) atoms at initial stage. The plot also illustrates that at ~ 680 K both excess and equilibrium vacancies approach the same concentration if assuming the excess vacancy (Ln C$x) linearly changes temperature (T). It means that the effect of the excess vacancies diminishes at this temperature. In addition, decay function for the excess vacancy changes: C$x = 1.54*

10-3* eO-nev/kT^ can be obtained.

The estimated vacancy migration energy (in Cu) is about 0.51eV, close to theoretical prediction [55]. This is actually the migration energy of vacancy in non-equilibrium state, which is clearly lower than the migration energy of vacancies in the absence of deformation at the equilibrium state, i.e. 0.98 eV[56] .

Table 2 Calculated equilibrium and derived excess vacancy concentration at different temperatures.

25°C 212°C 414°C Remarks

req rV 2.0x10 -19 1.17x10 -11 5.08 x10 -8 [52]

rex rV 5.0x10 -3 1.50x10 -6 < 5.0x10 -8 Eq. 1 and Eq. 2

Note that at 414°C the excess vacancy concentration is already lower than equilibrium concentration. Therefore, no value is input. Here, only non-equilibrium excess vacancy concentration in Cu is evaluated. Cr frequently acts as a diffusion barrier for Cu [57], and Cu diffusion coefficient in Cr matrix is not available in literature.

Similar to the study of Cu-Fe composites by X-ray absorption near edge structure (XANES) spectroscopy [58], it may be expected that the electronic structures at the strongly intermixed zones is different by closely viewing the energy loss near-edge structure (ELNES) of Cr-L23 and Cu-L23

absorption edges. However, the experimental measurements using the existing spectrometer with a limited energy resolution of ~0.8 eV hardly reveals a significant change of the fine structure as shown in Figure 8. The ELNES spectra of Cr-L23 and Cu-L23 were acquired from one same Cr/Cu interface at an step of 0.7 nm under different annealing temperatures. Here, a comparison between Cr-L23 edges along the same interface at room temperature (Figure 8a) and 212°C (Figure 8b) were made. Although the fine structures during the annealing not significantly change, however, the subtle changes are still distinguishable in fine structures between those prior to annealing and annealed at 212°C. As shown, 1) when crossing the interface from Cr to Cu, no clear chemical shift is visible, and the only remarkable change seems to occur at the Cr-L3 peak, which becomes quickly reduced at the interface region as compared to Cr bulk; 2) ELNES at 212°C shows more abruptly changes at the interface as compared to room temperature (25°C), which corresponds to the sharp contrasts in STEM images; (3) fine structure changes from 25°C to 212°C also reveal the interface region (width) is reduced, as labeled in Cr-L23 spectra, i.e. nearly changes from 2.8 nm (4 steps) to 2.1 nm (3 steps). Cu-L23 edge features at 25°C and 212°C show quite similar. However, at 25°C concurrent of Cu-L2,3 and Cr-L2,3 edge at the interface spans about 2.1nm (three steps) while at 212°C it is about 0.7 nm (one step). All signify the evolution of the interface forced intermixing zone with in-situ heating.

The measured L2/L3 ratio in the Cr-L23 edge also presents a gradient-reduced from Cr to Cu grain at the interface at 25°C, L2/L3 s from 0.25 to 0. It could be an indication that the electronic structure at the interfacial intermixing zone is varied. But the ratio changes slightly at 25°C and 212°C (not shown here). A theoretical calculation may further verify the observations, and a close study on the fine structures of Cr and Cu at the intermixing zone enabled using a high-energy resolution spectrometer might bring more details in the absorption edge of Cr-L2,3.

This experiment reveals how deformed CuCr nanocrystalline, in particular deformation-generated excess vacancies, evolve with in-situ heating. It provides a direct and approximate atomic-level measure on the stability of nanostructured materials. However, it should be noted that deformation can also induce some other phenomena, such as dislocations, microstructure irregularities etc... All these are relevant to the recovery behavior of plastically deformed materials. Here, we only show how to apply in-situ TEM technique to understand the stability of nanostructured materials via a combination of morphological observations and composition tracking, and further analyze the evolution of dynamic process.

Conclusions

To summarize, the structural and chemical composition evolution of a Cu-Cr NC material is revealed using advanced TEM via real-time. The experiments show that the simultaneous changes in chemical composition and local structure lead to the supersaturation concentration variation in the bcc and fcc lattices and eventually triggers the loss of thermal stability. Through the analysis of compositional

profiles recorded under different annealing conditions, and the dynamic process in NC materials was visualized in detail. The main conclusions are summarized as follows: (1) Chemical composition tracking of the dynamic process reveals that Cu-Cr NC materials destabilization started at a relative low temperature (~212°C), although the structural morphology hardly changed and grain size remained almost unaltered, the chemistry at the interface started fluctuation. (2)The chemistry profiles also shows that annealing leads to an obvious interface sharpening, and the width of intermixing zone in width become less. (3) Atomic mobility is dramatically enhanced by HPT deformation, and measured diffusion coefficient is tens of orders of magnitude larger than Cu self-diffusion coefficient (room temperature). (4) Vacancy concentration during HPT process is estimated. Moreover, a direct relation of deformation induced excess vacancy and temperature is derived. The effect of deformation induced excess vacancy would completely vanish at ~ 400°C. (4) The electronic structures at the interface are slightly different for 25°C and 212°C, being ascribed to the interface intermixing zone change. The study is of significance for understanding the thermal stability of nanocrystalline materials at an approximately atomic-level.

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Acknowledgements: The authors thank Dr. Andrea Bachmaier and Oliver Renk for helpful

discussion. Financial support by the FWF (Austrian Science Fund) within project number P27034-

N20 is gratefully acknowledged.

Supplementary materials

Supporting Information are available on detailed analysis of HRTEM image, grain sizes calculations

and ex-situ experiments.

Figure captions:

Figure 1 (a) A backscattered electron image of raw bulk material with Cu (bright) and Cr (dark) particles prior to deformation. (b) XRD measurements on the raw-material (prior to deformation) and after deformation (as-deformed CuCr nanocrystalline, 25 rotations, prior to annealing).

Figure 2 (a) Typical HRTEM images of nanocrystalline grain recorded along Cu [110] (inside the red curve) and Cr [111] directions, the orientation relationship: Cu [110]// Cr [111], two fast Fourier transform patterns are attached. A disordered (labeled) interface between Cr and Cu grains is clearly seen. One low-magnification TEM image is inserted on the top-right corner. Note that in the upper part there is another Cu grain which is along high-index zone. (b) HRTEM image of single Cr grains along [111] direction surrounded by several Cu grains.

Figure 3 HAADF STEM image demonstrating the nanocrystalline structural evolution process, (a) 25°C, (b) 212°C, (c) 414°C, note that the contrast changes with temperature. Particular locations are labelled by arrows. Some pores are formed immediately as annealing temperature upon 414°C.

Figure 4 Evolution of the average Cu grain size (D) with the annealing temperature; the D value at each temperature was obtained by measuring about 30 grains with irregular shapes and different dimensions. The error bars is about ±12 nm due to a large scatter in the grain size. The microhardness was measured from a HPT sample in the as-deformed state (400% deformation strain) and subsequently ex-situ annealed at 200°C, 300°C, 400°C, and 550°C for 30 min. The microhardness in the as-deformed state is surprisingly smaller compared to annealed conditions. Please note that, here, Cu grain growth was particularly measured due to the bright contrast.

Figure 5 In-situ chemical composition profile changes recorded at the same positions (identical interface) at 25°C (a), and 212°C (b). The curves are non-linear fits of the data. The corresponding images recorded at two temperatures are on the left-hand side. The arrows in the images indicate the position of the profile measured (along the arrow direction). The sample thickness is around 30 nm.

Figure 6 Chemical composition profiles recorded in-situ at 212°C and 414°C across the different Cr-Cu interface at a nearly identical thickness. The curves are non-linear fit of the two profiles. The profiles were taken at a specimen thickness of about 30 nm. Error function curves are fitted to the two profiles. The open and solid symbols represent experimental measurements at 414°C and 212°C, respectively. Inserted atomic models schematically show the dynamic process with annealing: 212°C with a relative large intermixing zone while it becomes narrower at 414°C. The variation of

intermixing zone, vacancy concentration and element diffusion with temperature are clearly displayed. Note that green and blue circles represent Cr and Cu while Blue Square represents vacancy.

Figure 7 (a) Diffusion coefficients derived by non-linear function fit are plotted as a function of temperature. Red curve represents measured diffusion coefficient while the black curve represents equilibrium diffusion coefficient. (b) The changes of derived excess vacancy concentration and equilibrium vacancy concentrations with temperature. Note that at temperatures exceeding ~ 320°C the excess vacancy is lower than the equilibrium vacancy concentration.

Figure 8 ELNES of Cr-L2, 3 and Cu-L2,3 absorption edges from Cr to Cu grains at a step of ~ 0.7 nm. (a) at room temperature, 25°C; (b) at 212°C. The numbers on the spectra correspond to the measured positions in Figure 5a and Figure 5b as labeled as starting and ending points.

Figures and Figure captions:

Figure 1 (a) A backscattered electron image of raw bulk material with Cu (bright) and Cr (dark) particles prior to deformation. (b) XRD measurements on the raw-material (prior to deformation) and after deformation (as-deformed CuCr nanocrystalline, 25 rotations, prior to annealing).

red curve) and Cr [111] directions, the orientation relationship: Cu [110]// Cr [111], two fast Fourier transform patterns are attached. A disordered (labeled) interface between Cr and Cu grains is clearly seen. One low-magnification TEM image is inserted on the top-right corner. Note that in the upper part there is another Cu grain which is along high-index zone. (b) HRTEM image of single Cr grains along [111] direction surrounded by several Cu grains.

Figure 3 HAADF STEM image demonstrating the nanocrystalline structural evolution process, (a) 25°C, (b) 212°C, (c) 414°C, note that the contrast changes with temperature. Particular locations are labelled by arrows. Some pores are formed immediately as annealing temperature upon 414°C.

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| Microhardness ] Averaged grain size

.y 30 Ico

300 400 500 600 700

Temperature (K)

(f) y) cu c

05 sz O

Figure 4 Evolution of the average Cu grain size (D) with the annealing temperature; the D value at each temperature was obtained by measuring about 30 grains with irregular shapes and different dimensions. The error bars is about ±12 nm due to a large scatter in the grain size. The microhardness was measured from a HPT sample in the as-deformed state (400% deformation strain) and subsequently ex-situ annealed at 200°C, 300°C, 400°C, and 550°C for 30 min. The microhardness in the as-deformed state is surprisingly smaller compared to annealed conditions. Please note that, here, Cu grain growth was particularly measured due to the bright contrast.

r>nrai üaiükixnsi blv^i >j i üi ii

Starting point

_l_i_1_

0 5 10 15

Distance (nm)

5 10 15

Distance (nm)

Figure 5 In-situ chemical composition profile changes recorded at the same positions (identical interface) at 25°C (a), and 212°C (b). The curves are non-linear fits of the data. The corresponding images recorded at two temperatures are on the left-hand side. The arrows in the images indicate the position of the profile measured (along the arrow direction). The sample thickness is around 30 nm.

Cr-Cu interface distance (nm)

Figure 6 Chemical composition profiles recorded in-situ at 212°C and 414°C across the different Cr-Cu interface at a nearly identical thickness. The curves are non-linear fit of the two profiles. The profiles were taken at a specimen thickness of about 30 nm. Error function curves are fitted to the two profiles. The open and solid symbols represent experimental measurements at 414°C and 212°C, respectively. Inserted atomic models schematically show the dynamic process with annealing: 212°C with a relative large intermixing zone while it becomes narrower at 414°C. The variation of intermixing zone, vacancy concentration and element diffusion with temperature are clearly displayed. Note that green and blue circles represent Cr and Cu while Blue Square represents vacancy.

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3.0x10

2.0x101

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Q) O it

g 1.0x10 O

CO it b

Temperature /K

800700 600 500

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-5 -10 -15 -20

(J -25 c:

-Ij -30 -35 -40 -45

induced excess vacancy

Equilibiurm vacancy concentration

103/T(K"1)

Figure 7 (a) Diffusion coefficients derived by non-linear function fit are plotted as a function of temperature. Red curve represents measured diffusion coefficient while the black curve represents equilibrium diffusion coefficient. (b) The changes of derived excess vacancy concentration and equilibrium vacancy concentrations with temperature. Note that at temperatures exceeding ~ 320°C the excess vacancy is lower than the equilibrium vacancy concentration.

IWai at iJi il at alif M Ai itfai >1I>]|J

-t—> w c

a> -*->

Cr Cu-L

560 570 580 590 600 610 Energy loss (eV)

750 9001050 Energy loss (eV)

-1—« if) c <D

560 570 580 590 600 Energy loss (eV)

610 750 9001050

tergy loss (eV)

Figure 8 ELNES of Cr-L2, 3 and CU-L2 3 absorption edges from Cr to Cu grains at a step of ~ 0.7 nm. (a) at room temperature, 25°C; (b) at 212°C. The numbers on the spectra correspond to the measured positions in Figure 5a and Figure 5b as labeled as starting and ending points.